Patent Publication Number: US-2004055671-A1

Title: Nanophase precipitation strengthened Al alloys processed through the amorphous state

Description:
[0001] 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  
       [0002] In a principal aspect, the present invention relates to aluminum-based alloys processed through an amorphous state or stage, preferably by means of a rapid solidification process from molten alloy, and then devitrified to a primarily crystalline structure by thermal or thermomechanical processing. In order to achieve a desired microstructure during processing, the aluminum alloy comprises a combination of rare earth and transition metal components with the aluminum. The final crystalline microstructure has stable high temperature strength (at or above about 300° C.) characterized by nanoscale intermetallic precipitates, preferably L1 2  phase precipitates.  
       [0003] Currently available commercial aluminum 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 the fan components in turbine engines where short-term tensile strengths of the order of 500 MPa (40-60 ksi) are desired. Precipitation hardening introduced by aging is a known method to strengthen aluminum alloys. Thus, conventional high strength aluminum alloys in commercial applications utilize GP (Guinier-Preston) zones and subsequent precipitation when seeking to achieve high material strength. With this type of strengthening, precipitates are usually formed during aging at or below 250° C. in order to produce appropriate strengthening precipitate dispersions.  
       [0004] Examples of aluminum alloys processed with relatively high aging temperatures in commercial practice include 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 improved microstructure stability relative to other commercial aluminum alloys. These aluminum alloys, when age-hardened, usually possess a room temperature yield strength of about 600 Mpa (85 ksi). However, at temperatures approaching 300° C., the precipitation strengthening 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.  
       [0005] In order to achieve a combination of high strength and usable high temperature properties in aluminum alloys, researchers have investigated a variety of intermetallic precipitation dispersions. Since selected intermetallic precipitates may determine the strength and ductility of the final aluminum alloy, it is desirable to select cubic precipitate phases with high crystallographic symmetry to promote ductility. The aluminum-based L1 2  phase (fcc ordered) is one of the best-known precipitates to achieve a good combination of high strength and high toughness.  
       [0006] However, since crystalline aluminum has very limited solubility for most intermetallics, including L1 2  forming components, it is difficult to produce fine intermetallic dispersions through crystalline solid-state heat treatments that utilize a step to dissolve the strengthening phase particles prior to reprecipitation during aging. However, rapid solidification processing (RSP) from the liquid state has been attempted or utilized. With an RSP process, it is possible to either (1) directly produce fine crystalline structure, or (2) produce partially or completely amorphous aluminum alloys. Nonetheless, crystalline RSP aluminum alloys have not been able to meet the high temperature strength requirements due to the difficulty of producing small, stable particle sizes at adequate volume fraction. Alternatively, the focus on amorphous aluminum alloys has been primarily on the glass formability. These glassy alloys in general have low ductility, and, when devitrified, they typically form crystalline microstructures that are far from optimal for fracture toughness and other requirements.  
       [0007] For example, RSP crystalline aluminum alloys, not designed to form glass with RSP, utilize a metastable and coherent L1 2  dispersion (about 1%) which converts into other phase structures at high temperature. [Parameswaran, Weertman and Fine, Scripta Metall., vol.23, pp.147, 1989] U.S. Pat. Nos. 4,950,452, 5,074,935, 5,334,266, 5,458,700, 5,431,751, 5,053,085 and its derived divisional patents (U.S. Pat. Nos. 5,240,517, 5,320,688, and 5,368,658) utilize RSP to achieve either amorphous, a mixture of amorphous and crystalline, or crystalline microstructures, without the use of an L1 2  strengthening dispersion. The aluminum alloys of U.S. Pat. No. 4,889,582 achieve the high temperature strength with RSP process and age hardening to produce Al 3 Fe (which has non-cubic DO 3 , not L1 2  structure) strengthening precipitates, using different alloying components. U.S. Pat. No. 4,464,199 describes high strength aluminum alloys with 4 to 15% iron processed with RSP to achieve high temperature stability. Use of iron promotes Al 3 Fe precipitation.  
       [0008] The aluminum alloys of U.S. Pat. No. 4,787,943 are composed of transition metals and rare earth metal components. However, with titanium and iron as the preferred transition metals, and gadolinium and cerium as rare earth elements, the alloy does not comprise a stable L1 2  strengthening dispersion.  
       [0009] U.S. Pat. No. 6,248,453B1 describes the use of an L1 2  strengthening dispersion for high strength aluminum alloys utilizing certain alloying components, but does not utilize a transition metal such as nicker and cobalt and an intermediate amorphous state as a means to refine the L1 2  microstructure.  
       [0010] Finally, the high strength and high temperature aluminum alloys referenced have not found widespread application due the lack of robust processing to produce those alloys and the deficiency of other alloy properties, such as toughness and ductility.  
       SUMMARY OF THE INVENTION  
       [0011] The present invention is directed to a new class of aluminum alloys characterized by formation from an intermediate amorphous state to a final face centered cubic (fcc) aluminum matrix strengthened by a precipitate dispersion in order to achieve high temperature strength with usable ductility. The disclosed aluminum alloys thus are comprised of a fine crystalline structure of primarily aluminum fcc matrix strengthened by a precipitate phase with high cubic symmetry, preferably the L1 2  phase. An appropriate melt of aluminum with crystalline formers is first processed to achieve an intermediate amorphous state to preserve the crystalline formers. The preferred method to achieve a primarily (above 70% in volume) amorphous state is a rapid solidification process (RSP) from the molten alloy by process techniques such as powder atomization, melt spinning, and spray casting. The RSP process preferably 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 devitrified to a final primarily (above 70% in volume) fcc/L1 2  crystalline microstructure with at least about 30% fcc phase in volume and at least about 10% L1 2  phase in volume. Thermal treatment, thermomechanical treatment, or a combination of both, are the preferred methods to produce the final crystalline microstructure.  
       [0012] The alloys of the invention demonstrate high strength capability after both short and long exposure to a high temperature at or above 300° C. by maintaining a fine and stable microstructure. For example, a short-term strength requirement of 40-60 ksi yield strength is expected. For the long-term strength, as characterized by creep strength, it has been demonstrated that optimal creep resistance at 300° C. for aluminum alloys is obtained from a dispersion of 25 nm diameter L1 2  particles [E. A. Marquis,  Microstructural Evolution and Strengthening Mechanisms in Al—Sc and Al—Mg—Sc Alloys , Ph.D. thesis, Northwestern University, 2002.]. Also, it has been demonstrated [Ohtera, Inoue and Masumoto, in  First International Conference on Processing Materials for Properties , Ed. Henein and Oki, The Minerals, Metals &amp; Materials Society, p.713, 1993] that aluminum-based alloys devitrified by extrusion from amorphous atomized powder can achieve aluminum fcc grain size of 80 to 170 nm, with intermetallic phase dispersions as fine as 10 to 80 nm in diameter. Such grain or particle sizes may therefore encompass the desired optimal particle diameter for creep strengthening by an L1 2  phase intermetallic dispersion. Accordingly, a process and an aluminum alloy product combining a sufficient amount of stable L1 2  refined or converted by devitrification from a glass phase material by appropriate thermal or thermomechanical processing has been discovered. The alloy meets or exceeds high temperature strength requirements with optimized creep resistance as a result of intermetallic, small particle phase dispersions.  
       [0013] The aluminum alloys of this invention are produced by combining aluminum, selected metals, particularly transition metals, rare earth elements, and other optional elements into a melt which is converted to an amorphous or glassy state and then subsequently to a crystalline state by a devitrification process. In the broadest sense about 2 to 15 atomic percent in the aggregate of at least one rare earth selected from the group consisting of erbium (Er), thulium (Tm), ytterbium (Yb) and lutetium (Lu) is provided for the melt. Further about 2 to 7 atomic percent in the aggregate of at least one transition metal selected from the group consisting of nickel (Ni), cobalt (Co), iron (Fe) and copper (Cu) is provided to enhance or facilitate the process step of forming an amorphous or glassy material from the melt. Other materials included, optionally, in the melt comprise additional transition metals, lithium (Li) and magnesium (Mg). For example, of less than about 5% in the aggregate by atomic fraction scandium (Sc), yttrium (Y), titanium (Ti), zirconium (Zr), vanadium (V), chromium (Cr), manganese (Mn) and lithium (Li) incorporated into the melt or alloy as a modifier. Less than about 7% in the aggregate by atomic fraction magnesium (Mg), zinc (Zn), silver (Ag) and niobium (Nb) may be incorporated into the melt or alloy to accommodate fcc misfit in the crystalline state.  
       [0014] The selection of alloying components is based on the function or criteria of (1) good L1 2  formers, (2) adequate glass formability with RSP—this leads to the selection of components with strong short range ordering effects and slow long-range diffusing kinetics in molten aluminum, and (3) good fcc/L1 2  coherency with low misfit which lead to a fine microstructure and good coarsening resistance at high temperature.  
       [0015] Lanthanide rare earth L1 2  formers such as erbium (Er), thulium (Tm), ytterbium (Yb), and lutetium (Lu) are slow diffusers and hence enhance glass formability while at the same time allowing L1 2  formation during devitrification. Scandium (Sc) is a strong L1 2  former which can be used in conjunction with rare earth components to promote L1 2  formation and reduce the misfit between fcc and L1 2  phases. Transition metals such as nickel (Ni), cobalt (Co), iron (Fe), and copper (Cu) can enhance short range ordering in the molten aluminum, leading to the increase of glass transition temperature and promoting glass formability. Magnesium (Mg), as a component with relatively high solubility in crystalline aluminum, can be used in fcc aluminum to further reduce the fcc/L1 2  misfit.  
       [0016] With efficient precipitation strengthening by the L1 2  dispersion, the alloys of the subject invention achieve a hardness of 490 VHN at significantly higher aging temperature compared to conventional aluminum alloys indicating a usable strength at or beyond 300° C.  
       [0017] Thus, it is an object of the invention to provide a new class of high-strength L1 2  phase strengthened aluminum alloys processed into the amorphous state, preferably with a rapid solidification process, and then subsequently devitrified into a crystalline structure with a thermomechanical or thermal processes.  
       [0018] A further object of the invention is to provide aluminum alloys with usable strength at or above 300° C. by selecting (1) slow diffusing L1 2  formers to prevent fast coarsening kinetics at high temperature, (2) appropriate matrix fcc and L1 2  components to reduce the lattice misfit between them to further slow down the coarsening kinetics and to promote the initial finer L1 2  morphology.  
       [0019] Another object of the invention is to utilize the cubic crystal structure of L1 2  phase to promote the interfacial coherency with an aluminum fcc matrix and the inherent ductile behavior of the L1 2  crystalline phase itself, due to high number of slip systems in cubic L1 2  structure, to maintain reasonable overall alloy ductility.  
       [0020] A further object of the invention is to combine alloying components so that the L1 2  formation is enhanced while the likelihood of forming other undesirable intermetallics is reduced, in order to further promote alloy ductility of an aluminum alloy.  
       [0021] Another object of the invention is to provide adequate glass formability during the rapid solidification process such as powder atomization or melt spinning to form an amorphous aluminum alloy structure, preserving the L1 2  formers in the solution before a subsequent devitrification process.  
       [0022] A further object of the invention is to utilize the intermediate, highly supersaturated amorphous state, slow L1 2  coarsening kinetics, and low L1 2 /fcc misfit to produce a fine microstructure in a strong aluminum alloy.  
       [0023] A further object of the invention is to provide aluminum alloys having high strength at high temperatures with appropriate toughness when processed with powder metallurgy techniques through an amorphous state.  
       [0024] Another object of the invention is to combine transition metals and rare earth components to provide sufficient glass formability during a rapid solidification process and a nanoscale L1 2  dispersion strengthened aluminum alloy fcc crystalline structure after devitrification.  
       [0025] These and other objects, advantages and features will be set forth in the detailed description which follows.  
     
    
    
     BRIEF DESCRIPTION OF THE DRAWING  
     [0026] In the detailed description which follows, reference will be made to the drawing comprised of the following figures:  
     [0027]FIG. 1 is a chart setting forth criteria for the design of aluminum alloys of the invention;  
     [0028]FIG. 2 is a diagram illustrating the viscosity of liquid aluminum nickel alloys calculated at 1700° C. in comparison with experimental data wherein the comparison is based upon calculations made using different thermodynamic descriptions of the aluminum nickel liquid (substitutional solution model without associate or associate model with AlNi associate);  
     [0029]FIG. 3 is an assessed Al—Yb phase diagram in comparison with experimental data;  
     [0030]FIG. 4 a  is a graph of the enthalpy of formation of cobalt and a rare earth, i.e. Er and Yb predicted by quantum mechanical calculation;  
     [0031]FIG. 4 b  is a graph of the mixing enthalpy of the materials set forth in FIG. 4 a;    
     [0032]FIG. 4 c  is a graph of the molar volume of the materials set forth in FIG. 4 a;    
     [0033]FIG. 5 is an X-ray diffractogram of the alloy of Example 1A as melt-spun with positions of fcc pure aluminum reflections indicating a fully amorphous state;  
     [0034]FIG. 6 is an isochronic DSC thermogram of the alloy of Example 1A indicating devitrification steps upon heating;  
     [0035]FIG. 7 is an X-ray diffractogram of the alloy of Example 1A after devitrification at 425° C. for 22 hours, with positions of reflections of pure fcc Al, Al 3 Yb, and Al 3 Er phases, indicating the desired phases: FCC+L1 2 ;  
     [0036]FIG. 8 is a SIMS elemental map of devitrified alloy of Example 1A indicating Er and Yb partition together to L1 2  phase, and NI and Co which do not partition to L1 2 ;  
     [0037]FIG. 9 is a scanning electron microscope (SEM) secondary electron image of devitrified alloy of Example 1A indicating submicron morphology;  
     [0038]FIG. 10 is a 3DAP microscopy reconstruction of Al 3 (Yb,Er)L1 2  phase of the alloy of Example 1A after devitrification at 424° C. for 19 hours;  
     [0039]FIG. 11 is an X-ray diffractogram of the alloy of Example 2A as melt-spun, with positions of fcc pure Al and L1 2  reflections, indicating fully amorphous status;  
     [0040]FIG. 12 is an X-ray diffractogram of the alloy of Example 2A after devitrification at 425° C. for 19 hours, with positions of reflections of pure fcc Al, Al 3 Yb, and Al 3 Er phases, indicating the desired phases: FCC+L1 2 ;  
     [0041]FIG. 13 is an X-ray diffractogram of the alloy of Example 2B as melt-spun, with positions of fcc pure Al and L1 2  reflections, indicating fully amorphous status;  
     [0042]FIG. 14 is an X-ray diffractogram of the alloy of Example 2B after devitrification at 425° C. for 19 hours, with positions of reflections of pure fcc Al, Al 3 Yb, and Al 3 Er phases, indicating the desired phases: FCC+L1 2  wherein the lattice parameter of L1 2  is reduced by Sc;  
     [0043]FIG. 15 is an X-ray diffractogram of the alloy of Example 2C as melt-spun, with positions of fcc pure Al and L1 2  reflections, indicating almost fully amorphous status;  
     [0044]FIG. 16 is an X-ray diffractogram of the alloy of Example 2C after devitrification at 425° C. for 19 hours, with positions of reflections of pure fcc Al, Al 3 Yb, and Al 3 Er phases, indicating the desired phases: FCC+L1 2  wherein the lattice parameter of L1 2  is expanded by Mg;  
     [0045]FIG. 17 is an X-ray diffractogram of the alloy of Example 2D as melt-spun, with positions of fcc pure Al and L1 2  reflections, indicating fully amorphous status;  
     [0046]FIG. 18 is an X-ray diffractogram of the alloy of Example 2D after devitrification at 425° C. for 19 hours, with positions of reflections of pure fcc Al, Al 3 Yb, and Al 3 Er phases, indicating the desired phases: FCC+L1 2  and wherein the lattice parameter of fcc is expanded by Mg and that of L1 2  is reduced by Sc;  
     [0047]FIG. 19 is an X-ray diffractogram of the alloy of Example 1C as melt-spun, with positions of fcc pure Al and L1 2  reflections, indicating non-amorphous status due to the absence of transition metal elements;  
     [0048]FIG. 20 is an X-ray diffractogram of the alloy of Example 1C after devitrification at 425° C. for 19 hours, with positions of pure fcc Al, Al 3 Yb, and Al 3 Er phases, indicating elements ER, Yb, and Sc can be completely intersoluble to form L1 2  and wherein the lattice parameter of L1 2  is reduced by Sc; and  
     [0049]FIG. 21 is a scanning electron microscope (SEM) secondary electron image of a devitrified alloy of Example 1C showing coarse microstructure in comparison with the alloy of Example 1A (see FIG. 9) and indicating that passing through the glassy state can reduce the particle size of precipitation.  
    
    
     DESCRIPTION OF THE PREFERRED EMBODIMENTS  
     [0050] General Summary  
     [0051] In general, the subject matter of the invention comprises a product (alloy) as well as a product by process and a process all directed to an aluminum alloy in crystalline form having higher or greater strength particularly at elevated temperatures, i.e. greater than at least 250° C. and preferably 300° C. The alloy of aluminum is made by compounding a mixture of aluminum with selected rare earths (RE) and transition metals (TM) in an amorphous or glassy state followed by devitrification of the material from the amorphous state to a mixed crystalline state comprised of face centered cubic (fcc) and L1 2 . (an ordered fcc) state wherein the ratios of the crystalline states is within certain preferred ranges. Preferably the resultant alloy has at least about 30% by volume fcc phase and at least about 10% by volume L1 2  phase material with limited residual amorphous or other phase material such as quasi-crystalline or ordered amorphous material.  
     [0052] The choice of starting materials may be varied as may the compounding processes, the glass phase formation processes and the devitrification processes. Importantly, when in the amorphous state, there may be some crystalline material contained therein, but preferably no more than about 30% by volume in the fcc phase. Thus forming the mixture in the amorphous or glassy intermediate state and subsequent devitrification constitute very important aspects of the invention.  
     [0053] The alloy materials, in addition to aluminum (Al), include one or more lanthanide rare earths (RE) selected or taken from the group of erbium (Er), thulium (Tm), ytterbium (Yb), and lutetium (Lu), and one or more transition metals (TM) taken or selected from the group of copper (Cu), nickel (Ni), cobalt (Co), titanium (Ti), iron (Fe), yttrium (Y), and scandium (Sc). Further additives in minor amounts (less than about five percent (5%) atomic percent) may be included such as magnesium (Mg). The range of various constituents is also important to maintain processability and preserve transition characteristics. Thus RE materials are utilized in the range of about 2 to 15 atomic percent (%), and the TM and metals are utilized in the aggregate range of about 2 to 12 atomic percent (%). Transition metals (TM) in the range of about 2 to 7 atomic percent are utilized.  
     [0054] 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 based upon examples as reported hereinafter. The completeness of the amorphous state is preferably at least about 70% by volume and preferably greater. The amorphous state is a solid noncrystalline state or phase prior to devitrification.  
     [0055] Development Technique  
     [0056] Devitrification is accomplished via heat treating (thermal) processing, thermo-mechanical processing as well as other processing techniques.  
     [0057] Precipitation strengthened aluminum (Al) alloys are difficult to develop for high strength due to limited solubility of alloying elements in fcc Al. The low solubility of strengthening dispersion elements prevents the solid-state solution treatment of the alloy to bring strengthening components into solution. It is hence not necessarily possible to allow subsequent aging to produce a desirable strengthening microstructure. Al alloys with high fractions of precipitate that cannot be completely solution treated therefore have very coarse particle sizes that tend to limit strength, corrosion resistance and toughness. In contrast, the Al alloys exhibit composition and microstructure characteristics that enable the Al alloys of invention to exhibit high strength, good ductility, and high temperature stability at or above 300° C.  
     [0058] By selecting an appropriate composition, processing techniques achieve a fully amorphous state after rapid cooling of the Al alloy composition. Furthermore, with careful selection of composition this glass can then be thermal or thermo-mechanically processed such that the glass devitrifies into a fully crystalline fcc matrix with intermetallic precipitates having a very fine size scale. By passing through the glass state, the equilibrium solidification that would produce coarse precipitates is avoided. Furthermore, the selection of the precipitate phase as L1 2  allows the precipitate to be coherent or semi-coherent with the fcc matrix improving the ductility of the alloy and the resistance of the precipitate to coarsen during subsequent thermal treatment.  
     [0059]FIG. 1 is a system flow-block diagram that illustrates the processing/structure/properties relationships for alloys of the invention. The two-stage nature of the chart emphasizes the generally equal importance of processability and final performance in the design of alloys of the invention. The processability is governed by the glass forming ability; strength, ductility, and thermal stability are the required end properties.  
     [0060] Certain transition metals (TM) (such as Fe, Co, Ni, 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. Rare-earth (RE) elements (such as Ce, Gd, Yb, 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.  
     [0061] Viscosity of liquid metals also has a strong effect on the glass forming ability. Liquid phase metals having high viscosity usually becomes glass through fast cooling. A binary model to describe liquid&#39;s viscosity is set forth in Equation (1):  
             η   =       -         x   1          x   2         R                 T              (         x   1          η   1                o         +       x   2          η   2                o           )              ∂   2          G   m         ∂     x   2   2                   (   1   )                       
 
     [0062] The variables x i  and G m  represent the mole fraction and molar Gibbs energy of the liquid, respectively. The viscosity of pure components, η i   o , is available from literature, see Table 1. From the above equation it is seen that the composition dependence of the viscosity is fully determined by the thermodynamic properties of the liquid.  
               TABLE 1                          Viscosity and related data for pure liquid metals       [L. Battezzati and A. L. Greer, “The Viscosity of Liquid Metals and Alloys,” Acta       metall. 37, 1791 (1989).]                                     V m     η(T m )   C A                                                       T m     (10 −6     (10 −3     [10 −7  (J/K   E   B   η o                                               Element   (K)   m 3 /mol)   Pa s)   mol ⅓ ) ½ ]   (kJ/mol)   (E/RT m )   (10 −3  Pa s)                                                     Al   933   11.3   1.30   1.30   16.5   2.13   0.149       Co   1765   7.59   4.18   1.58   44.4   3.03   0.255       Cr   2133   8.28   5.7   ˜2.2   ˜185   ˜10.43   1.7 × 10 −4         Cu   1356   7.94   4.0   1.72   30.5   2.71   0.301       Fe   1809   7.96   5.5   2.18   41.4   2.75   0.370       Mn   1517   9.58   5   2.5   20-46.5   1.6-3.7   0.12-1.02       Ni   1728   7.43   4.90   1.85   50.2   3.49   0.166       V   2175   8.9   2.4   0.98   ˜73   ˜4.04   ˜0.042       Ti   1958   11.6   2.2   1.18   ˜68   ˜4.18   ˜0.034       Zr   2125   ˜15.7   3.5   1.58   ˜88   ˜4.98   ˜0.024                                                                          
 
     [0063] Equation (1) has been validated for several binary systems where both experimental viscosity data and thermodynamic descriptions are available. An example is the Al—Ni system where an intermediate compound AlNi melts congruently at about 1650° C. FIG. 2 shows the calculated viscosity in Al—Ni liquid at 1700° C. as a function of Ni content, in comparison with experimental data. It is seen that the calculated viscosity shows a similar trend as the experimental data, but the latter indicates a stronger increase of the viscosity around 50 at. % Ni. The strong increase in viscosity is believed to be caused by molecular associates of Al—Ni in the liquid. In order to illustrate the effect of the AlNi associates on the viscosity, the Al—Ni system was thermodynamically reassessed with an associate model for the liquid free energy. The viscosity of Al—Ni liquid calculated using the AlNi associate model is also plotted in FIG. 2 for comparison. The calculation shows a significant increase in the viscosity due to the formation of liquid AlNi associates.  
     [0064] This viscosity model has further been extended to multi-component systems, which provides guidance of glass forming ability for the design of alloys of the invention.  
     [0065] A model to predict glass transition temperature, T g , has been developed for the Al—TM—RE system. The model is based on the parabolic relationship between T g  and room temperature glass shear modulus, G, and it assumes that G is linear in atomic fraction, X, of each element according to equation (2):  
                     T   g     =       (     14100        ∑       K   i          X   i           )     0.49                   i   =   Al     ,   Mg   ,   Ni   ,   Co   ,   Er   ,   Yb   ,   Sc   ,   Y   ,     etc   .                   (   2   )                       
 
     [0066] The coefficient for each element, K i , has been calibrated to experimental DSC data of melt-spun amorphous ribbons. These coefficients quantitatively represent the relative effectiveness of the aforementioned molecular associates in promoting glass forming ability by varying T g . This model also provides guidance of glass forming ability for the design of alloys of the invention. The prediction error of equation (2) compared to available experimental data is about 9 K with a standard deviation of 10 K, which is within the experimental precision.  
     [0067] The desired performance for applications (e.g. aircraft structural components, high temperature turbine components, etc.) determines a set of alloy properties required, including strength, ductility and thermal stability resistance up to 300° C. The desired properties can be achieved through a fine, stable, multicomponent L1 2  phase dispersion in the Al matrix after devitrification of the glassy Al, as long as a sufficient volume fraction of L1 2  and low misfit with the Al matrix and thus low coarsening rate can be obtained.  
     [0068] The elements Er, Lu, Tm and Yb are known as the only stable rare earth (RE) L1 2  formers. The melting point and cost of these rare earth elements are presented in Table 2. Some properties of corresponding trialuminides are also given in Table 2, including lattice parameter, density, and L1 2  decomposition or congruent melting temperature. Among the four RE elements, according to Table 2, Yb has the smallest lattice parameter and relatively lower cost. Er has the lowest cost. Both Yb and Er have relatively lower melting point. Usually, the lower the melting temperature, the higher the glass forming ability. As a consequence, Alloys of the invention utilize Yb and Er as preferred L1 2  formers rather than Tm and Lu.  
               TABLE 2                          Stable L1 2  formers and their properties.                                         Melting       Lattice       L1 2  Max.                                         Element   Point   Cost a     L1 2     Parameter   Density   Temp.               Er   1529° C.   $725/kg   Al 3 Er   4.226 Å   5.507 g/cm 3     1070° C.       Lu   1663° C.   $7500/kg   Al 3 Lu   4.186 Å   5.783 g/cm 3     1200° C.       Tm   1545° C.   $6500/kg   Al 3 Tm   4.203 Å   5.591 g/cm 3         Yb    819° C.   $1600/kg   Al 3 Yb   4.202 Å   5.682 g/cm 3      980° C.       Sc   1541° C.   $18,000/kg   Al 3 Sc   4.101 Å   3.03 g/cm 3     1320° C.                          
 
     [0069] Sc is the only transition metal (TM) element that can form a stable L1 2  with Al. In comparison with RE L1 2  formers, Sc can form L1 2  with a smaller lattice parameter of 4.101 Å, thus leading to smaller misfit between L1 2  and the Al matrix. However, Sc is by far the most expensive of the elements considered, as seen in Table 2. The alloy of the invention therefore seeks to limit the use of Sc as much as possible. Efforts have been made to search for other transition metals to substitute for Sc. A preliminary requirement for such substitution is solubility. Experiments have shown that Ti has a substantial solubility in Al 3 Sc, up to 12.5 at. %, i.e. Al 0.75 Sc 0.125 Ti 0.125 . In addition, Ti has the lowest diffusion coefficient in solid Al among the transition metals. For instance, the diffusivity in Al at 300° C. is 9.0×10 −20  m 2 /s for Sc, 6.3×10 −24  m2/S for Zr, and 2.7×10 −25  m 2 /S for Ti. Adding Ti to Al 3 Sc will thus reduce the coarsening rate of L1 2  precipitates. Moreover, addition of Ti decreases the lattice parameter of Al 3 (Sc,Ti) and hence minimizes the lattice misfit with Al. The lattice parameters for seven hypothetical L1 2  compounds with transition metals are calculated at zero degree Kelvin (0 K) using the first-principle quantum mechanical approach. Results are shown in Table 3. All seven hypothetic L1 2  compounds show significant smaller lattice parameters than those stable L1 2  phases. Thus, alloys of the invention incorporate Yb, Er and Sc as base L1 2  formers but are not limited to these elements. TMs such as Ti, V, Zr, etc, which will result in low misfit and thus lower interfacial energy to yield a low driving force for particle coarsening are considered useful.  
               TABLE 3                          Lattice Parameters of hypothetical L1 2  TM Compounds through First       Principle Calculations at 0 K.                             TMs Hypothetic L1 2     Lattice Parameter, Å                                         Al 3 Ti   3.967           Al 3 Zr   4.085           Al 3 V   4.045           Al 3 Fe   3.8005           Al 3 Ni   3.8469           Al 3 Co   3.7946           Al 3 Cu   3.9363                      
 
     [0070] It should be pointed out that the compound Al 3 Ti is metastable as the L1 2  phase but stable as D0 22 . It can be transformed to the L1 2  structure if Al is substituted partially by Cr, Mn, Fe, Co or Ni.  
     [0071] For a multicomponent L1 2  phase, the lattice parameter a can be calculated based on a linear model similar to Vegard&#39;s law. As an example, we have the following equation for L1 2  with the formula Al 3 (Er x Yb y Sc z ) is utilized: 
       a [Al 3 (Er x Yb y Sc z   ]=x*a (Al 3 Er)+ y*a (Al 3 Yb)+ z*a (Al 3 Sc)  (3) 
     [0072] Assuming L1 2  has a linear thermal expansion, its lattice parameter at elevated temperature can be calculated as follows (T in ° C.): 
       a ( L 1 2   ,T )= a ( L 1 2   , RT )×[1+( T −25)×α( L 1 2 )]  (4) 
     [0073] where α(L1 2 ) represents the thermal expansion estimated according to the following reciprocal relation:  
                 α        (   M   )         α        (       Al   3        M     )         =         T   m          (     A                   l   3        M     )           T   m          (   M   )                 (   5   )                       
 
     [0074] where M represents the corresponding pure metal in Al 3 M and T m  is the melting point (° K.). Table 4 summarizes melting point, thermal expansion of relevant elements in alloys of the invention and the estimated thermal expansion as well as lattice parameter at 25° C. and 300° C. for stable L1 2  or meta-stable L1 2 .  
               TABLE 4                          Estimation of Thermal Expansion Coefficients and Lattice Parameters.                                 Thermal   Thermal                                                         Melting   Expansion,   Expansion,       Melting   Thermal   Lattice   Lattice           Point,   1/° K.   1/° K.   L1 2     Point,   Expansion,   Parameter   Parameter,                                                 Element   ° K.   20˜100° C.   20˜300° C.   Alloys   ° K.   20˜300° C.   Å, 25° C.   Å, 300° C.                                                         Ni   1728   13.3 × 10 −6     14.4 × 10 −6     Al 3 Ni   913   27.254 × 10 −6     3.8639   3.8929       Co   1768   12.5 × 10 −6     13.6 × 10 −6     Al 3 Co   1408   17.077 × 10 −6     3.8114   3.8293       Yb   1092    25 × 10 −6      25 × 10 −6     Al 3 Yb   1253   21.788 × 10 −6     4.202   4.2272       Er   1802    9.2 × 10 −6      9.2 × 10 −6     Al 3 Er   1343   12.344 × 10 −6     4.226   4.2403       Sc   1814    12 × 10 −6      12 × 10 −6     Al 3 Sc   1593   13.665   4.101   4.1164       Ti   1943    8.9 × 10 −6      9.2 × 10 −6     Al 3 Ti   1623   11.014   3.967   3.979       Y   1795   10.8 × 10 −6     10.8 × 10 −6     Al 3 Y   1253   15.472   4.299   4.3173       Al           25.3 × 10 −6                 4.0497   4.0779                  
 
     [0075] In the practice of the invention, decrease of the misfit by expanding the lattice parameter of the Al matrix is a useful technique. There are few elements that have significant solubility in Al. Mg is the only one that can substantially increase the lattice spacing of Al. According to literature information, in dilute AlMg alloys, every one atom percent of Mg will expand the Al lattice by 0.0045 Å. Further, every one atom percent of Mg will increase the linear coefficient of thermal expansion of Al by 1.179×10 −7 /° C.  
     [0076] Thermodynamic modeling and assessments have been performed for Al—TM—RE systems with relevant elements incorporated in alloys of the invention. The assessed Al—Yb binary phase diagram is shown in FIG. 3. The L1 2  Al 3 Yb is the strengthening phase that is used in alloys of the invention. According to FIG. 3, Al 3 Yb is not solution treatable. The only way to avoid coarse precipitation from Al liquid is using fast cooling. It is possible to get fine crystalline precipitates directly by rapid solidification processing (RSP) from the melt. However, much finer precipitates can be generated through devitrification of amorphous state due to the larger driving force. In addition, the precipitation process can be controlled through devitrification.  
     [0077] Quantum mechanical calculations have been performed in Al 3 Er—Al 3 Co and Al 3 Yb—Al 3 Co systems. The predicted enthalpy of formation, mixing enthalpy and molar volume in the corresponding systems are plotted in FIG. 4. According to FIG. 4, there is a strong repulsive Co—RE interaction indicating negligible Co solubility in L1 2  Al 3 RE. This sets a challenge in producing good glass forming alloys through conventional technology in the same composition intended to precipitate a significant amount of L1 2  precipitation through devitrification.  
     [0078] Experimental Results  
     [0079] 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 L1 2  after devitrification in the Al matrix. Alloys of the invention are summarized in Table 5 for their composition, status in melt-spun and devitrification conditions, as well as the lattice parameter of the Al fcc and the L1 2  phases and the misfit between L1 2  and Al matrix.  
               TABLE 5                          Summary of Composition and Microstructure of Al Alloys                                         Alloy Composition       Main phases   Lattice Parameter   Misfit                                         No.   at. %   As-spun status   after devitrfication   L1 2 , a(Å)   fcc, a(Å)   %                                                 1A   Al-2.5Ni-2.5Co-   fully glass   fcc, L1 2  +   4.215   4.063   3.735           3.5Yb-3.5Er       minor x phase       1B   Al-2.75Ni-2.75Co-2.17Yb-   fully glass   fcc, L1 2  + Al 3 Y +   4.217   4.050   4.123           2.17Er-2.17Y       minor x phase       1C   Al-6Yb-6Er   fcc, L1 2  +   fcc, L1 2     4.208   4.057   3.726               some glass       1D   Al-4Yb-4Er-4Y   fcc, L1 2  +   fcc, L1 2  + Al 3 Y   4.229   4.059   4.192               some glass       1E   Al-4.5Yb-4.5Er-3Sc   fcc, L1 2  +   fcc, L1 2     4.183   4.060   3.032               some glass       1F   Al-12Er-3Mg   fcc, L1 2  +   fcc, L1 2     4.216   4.059   3.865               some glass       1G   Al-12Er-6Mg   fcc, L1 2  +   fcc, L1 2     4.217   4.065   3.739               some glass       1H   Al-12Yb-3Mg   fcc +   fcc, L1 2     4.216   4.076   3.435               some glass       1I   Al-12Yb-6Mg   fcc, L1 2  +   fcc, L1 2     4.191   4.061   3.201               some glass       2A   Al-2Ni-2Co-4Yb-4Er   almost fully   fcc, L1 2  +   4.201   4.049   3.754               glass + minor fcc   minor x phase       2B   Al-2.5Ni-2.5Co-2Yb-   almost fully   fcc, L1 2  +   4.150   4.048   2.520           2Er-3Sc   glass + minor fcc   minor x phase       2C   Al-3Mg-2.5Ni-2.5Co-   mostly glass +   fcc, L1 2  +   4.197   4.060   3.374           4.5Yb-2.5Er   minor fcc   minor x phase       2D   Al-3Mg-2.5Ni-2.5Co-   fully glass   fcc, L1 2  +   4.155   4.057   2.416           2.5Yb-1.5Er-3Sc       minor x phase                  
 
     [0080] According to Table 5, alloys with TM elements (e.g. Ni and Co) are fully glass or almost fully glass in the as melt-spun conditions. These alloys include alloy 1A, 1B, 2A, 2B, 2C and 2D. Alloys without Ni and Co primarily show FCC Al and L1 2  crystalline phases although they are partially amorphous in the as melt-spun conditions. These alloys include alloy 1C, 1D, 1E, 1F, 1G, 1H and 1I. The results confirm the prediction that TM elements such as Ni and Co are critical elements to promote the glass forming ability in the alloys of the invention. Through devitrification, all alloys precipitate mainly L1 2  in the Al matrix. It should be emphasized that 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 L1 2  already present in the as-spun condition.  
     [0081] Following is a detailed demonstration of characteristics of alloys in Table 5.  
     EXAMPLE 1  
     [0082] Alloy 1A in Table 5 was fabricated into both powders using gas atomization and ribbons using melt spinning technology (wheel speed 55 m/s). In the as melt-spun condition, Alloy 1A was initially characterized by X-ray diffraction (XRD) and differential scanning calorimetry (DSC). The diffractometer trace (Cu—Kα radiation ) of as-spun Alloy 1A in FIG. 5 shows a broad peak characteristic of glass, and the ribbons show very good ductility, also characteristic of the glassy state. A DSC trace run at 40° C./min. shown in FIG. 6 indicates three stages of glass devitrification at temperatures near 200° C., 315° C. and 370° C.  
     [0083] DSC thermograms indicated that for heating at a constant rate of 10-40° C./min, the dominant crystallization reactions occur below about 400° C. for all alloys under investigation. Based on this information, a temperature of 425° C. was selected for the devitrification anneal. Specifically, alloy samples were encapsulated in quartz tubes under a protective atmosphere of approximately 500 mbar Argon with 99.999% purity. The encapsulated specimens were kept at 425° C. for periods of 19 or 22 hrs, respectively, to ensure completion of the crystallization reaction and grain growth to a size that facilitates experimental characterization.  
     [0084]FIG. 7 shows an X-ray diffractogram of Alloy 1A after devitrification at 425° C. for 22 hours. At this point transformations I-III are fully completed and the ribbons are fully devitrified in the sense that no significant amount of amorphous material can be detected by XRD and TEM. The peak positions of (111), (200), and (220) reflections calculated for FCC pure Al (lattice parameter 4.0496 Å) are indicated in red. Within the accuracy of this experiment (that is 0.01° at 2θ=38.505° (Al (111)), the measured peak positions are identical to pure Al, indicating the equilibrium solubility of Ni, Co, Yb, and Er in Al is very small. Also marked in red are the positions of the ‘forbidden’ reflections (110), (210), (211), and (221) for FCC Al, which are superlattice reflections of the L1 2  structure. The peak positions for L1 2  Al 3 Yb (lattice parameter 4.202 Å marked in green, and that for L1 2  Al 3 Er (lattice parameter 4.215 Å) are marked in brown. It can be seen that the measured peak positions are perfectly matched by L1 2  structure reflections. Therefore, Alloy 1A mainly consists of FCC Al and L1 2  phases as devitrified microstructure. It is also noticed that there is a minor third phase x, which crystalline structure is to be identified.  
     [0085] The FCC Al+L1 2  microstructure of Alloy 1A was investigated by the secondary ion mass spectrometry (SIMS) technology. FIG. 8 indicates that Yb and Er partition together into L1 2  precipitates, while Ni and Co have limited partitions.  
     [0086] The morphology of FCC+L1 2  microstructure of Alloy 1A is shown in FIG. 9. This SEM secondary electron image indicates that the majority of L1 2  precipitates are less than 1 micron after devitrification at 425° C., 22 hours. Based upon previously discussed experience in devitrified Al alloys, this indicated the potential to obtain nanoscale L1 2  precipitates with thermomechanical processing.  
     [0087] The three dimension atom probe microscopy (3DAPM) analysis confirms that L1 2  exists in Alloy 1A with formula Al 0.72 (Yb,Er) 0.28 . Yb and Er have about the same site fraction. Ni and Co are negligible in L1 2  phase. See FIG. 10.  
     [0088] Alloy 1A was also studied through microhardness testing in ribbon conditions after thermal cycles required for powder consolidation and extrusion. Results are presented in Table 6. Through devitrification at a temperature near 300° C., Alloy IA has demonstrated a hardness of 490 VHN, which is equivalent to the room temperature tensile strength exceeding 1000 MPa according to Inoue&#39;s work in similar alloy systems. (Akihisa Inoue, “Amorphous, nanoquasicrystalline and nanocrystalline alloys in Al-based systems”, Progress in Materials Science 43 (1998) 365-520). This result indicates high efficient strengthening of L1 2 .  
               TABLE 6                          Microhardness testing of Alloy 1A                             Condition   Hardness (VHN)                       As-spun   369           177° C. 1 hr   276           177° C. 10 hrs   310           177° C. 20 hrs   335           177° C. 20 hrs then 288° C. 2 hrs.   490           177° C. 20 hrs then 316° C. 2 hrs.   340           177° C. 20 hrs then 400° C. 2 hrs.   275                      
 
     EXAMPLE 2  
     [0089] Alloy 2A in Table 5 was fabricated into ribbons using melt spinning technology (wheel speed 55 m/s). In the as melt-spun condition, Alloy 2A was characterized by X-ray diffraction. The diffractometer trace (Cu—Kα radiation ) of as-spun Alloy 2A in FIG. 11 shows a broad peak characteristic of glass, and the ribbons show reasonable good ductility, also characteristic of the glassy state.  
     [0090]FIG. 12 shows an X-ray diffractogram of Alloy 2A after devitrification at 425° C. for 19 hours. At this point the ribbon is fully devitrified in the sense that no significant amount of amorphous material can be detected by XRD. The peak positions of (111), (200), and (220) reflections calculated for FCC pure Al (lattice parameter 0.40496 Å) are indicated in red. The measured peak positions are identical to pure Al, indicating the equilibrium solubility of Ni, Co, Yb, and Er in Al is very small. Also marked in red are the positions of the ‘forbidden’ reflections (110), (210), (211), and (221) for FCC Al, which are superlattice reflections of the L1 2  structure. The peak positions for L1 2  Al 3 Yb (lattice parameter 4.202 Å) are marked in green, and that for L1 2  Al 3 Er (lattice parameter 4.215 Å) are marked in brown. It can be seen that the measured peak positions are perfectly matched by L1 2  structure reflections. Therefore, Alloy 2A mainly consists of FCC Al and L1 2  phases as devitrified microstructure. It is noticed that the third phase x, in Example 2A, is less than in Example 1A, reflecting the desire to reduce the amount of x phase by reducing Ni and Co in Alloy 2A while retaining reasonable glass forming ability.  
     EXAMPLE 3  
     [0091] Alloy 2B in Table 5 was fabricated into ribbons using melt spinning technology (wheel speed 55 m/s). In the as melt-spun condition, Alloy 2B was characterized by X-ray diffraction. The diffractometer trace (Cu—Kαradiation) of as-spun Alloy 2B in FIG. 13 shows a broad peak characteristic of glass, and the ribbons show reasonable good ductility, also characteristic of the glassy state.  
     [0092]FIG. 14 shows an X-ray diffractogram of Alloy 2B after devitrification at 425° C. for 19 hours. At this point the ribbon is fully devitrified in the sense that no significant amount of amorphous material can be detected by X-ray diffraction (XRD). The peak positions of (111), (200), and (220) reflections calculated for FCC pure Al (lattice parameter 4.0496 Å) are indicated in red. The measured peak positions are identical to pure Al, indicating the equilibrium solubility of Ni, Co, Yb, Er and Sc in Al is very small. Also marked in red are the positions of the ‘forbidden’ reflections (110), (210), (211), and (221) for FCC Al, which are superlattice reflections of the L1 2  structure. It can be seen that the measured peak positions are matched by L1 2  structure reflections with lattice parameter 4.150 Å. It demonstrates that Yb, Er and Sc can be completely intersoluble to form L1 2 , and the lattice parameter has been significantly reduced by adding Sc, which is another means of reducing misfit to keep the L1 2  precipitates coarsening resistant. The misfit between FCC Al and L1 2  is 2.52% for Alloy 2B.  
     EXAMPLE 4  
     [0093] Alloy 2C in Table 5 was fabricated into ribbons using melt spinning technology (wheel speed 55 m/s). In the as melt-spun condition, Alloy 2C was characterized by X-ray diffraction. The diffractometer trace (Cu—Kα radiation ) of as-spun Alloy 2C in FIG. 15 shows a broad peak characteristic of glass, and the ribbons show reasonable good ductility, also characteristic of the glassy state.  
     [0094]FIG. 16 shows an X-ray diffractogram of Alloy 2C after devitrification at 425° C. for 19 hours. At this point the ribbon is fully devitrified in the sense that no significant amount of amorphous material can be detected by XRD. The peak positions of (111), (200), and (220) reflections calculated for FCC pure Al (lattice parameter 4.0496 Å) are indicated in red. The measured peak positions indicate that the lattice space of Al has been dilated to 4.060 Å, reflecting another feature of the invention: adding Mg to expand the lattice space of Al and thus decrease the misfit between Al and L1 2 . Also marked in red are the positions of the ‘forbidden’ reflections (110), (210), (211), and (221) for FCC Al, which are superlattice reflections of the L1 2  structure. It can be seen that the measured peak positions are perfectly matched by L1 2  structure reflections. Therefore, Alloy 2C mainly consists of FCC Al and L1 2  phases as devitrified microstructure.  
     EXAMPLE 5  
     [0095] Alloy 2D in Table 5 was fabricated into ribbons using melt spinning technology (wheel speed 55 m/s). In the as melt-spun condition, Alloy 2D was characterized by X-ray diffraction. The diffractometer trace (Cu—Kα radiation ) of as-spun Alloy 2D in FIG. 17 shows a broad peak characteristic of glass, and the ribbons show very good ductility, also characteristic of the glassy state.  
     [0096]FIG. 18 shows an X-ray diffractogram of Alloy 2D after devitrification at 425° C. for 19 hours. At this point the ribbon is fully devitrified in the sense that no significant amount of amorphous material can be detected by XRD. The peak positions of (111), (200), and (220) reflections calculated for FCC pure Al (lattice parameter 4.0496 Å) are indicated in red. The measured peak positions indicate that the lattice space of Al has been dilated to 4.057 Å. Also marked in red are the positions of the ‘forbidden’ reflections (110), (210), (211), and (221) for FCC Al, which are superlattice reflections of the L1 2  structure. It can be seen that the measured peak positions are matched by L1 2  structure reflections with lattice parameter 4.155 Å. It demonstrates that Yb, Er and Sc can be completely intersoluble to form L1 2 , and the lattice parameter has been significantly reduced by adding Sc. Alloy 2D reflects the means of reducing misfit by adding Mg to dilate the lattice space of Al and adding Sc to decrease the lattice parameter of L1 2 . Thus Alloy 2D achieved the lowest misfit so far:2.42% without losing the glass forming ability. Alloy 2D mainly consists of FCC Al and L1 2  phases as devitrified microstructure.  
     EXAMPLE6  
     [0097] Alloy 1E in Table 5 was fabricated into ribbons using melt spinning technology (wheel speed 55 m/s). In the as melt-spun condition, Alloy 1E was characterized by X-ray diffraction. The diffractometer trace (Cu—Kα radiation ) of as-spun Alloy 1E in FIG. 19 shows crystalline reflection peaks, indicating non-amorphous or partially non-amorphous state. It demonstrates that TMs are necessary elements to promote glass in Al—TM—RE system. In addition, the ribbons are very brittle, also characteristic of the non-glassy state.  
     [0098]FIG. 20 shows an X-ray diffractogram of Alloy 1E after devitrification at 425° C. for 19 hours. This diffractogram shows that Alloy 1E only has two phases: fcc Al+L1 2 . It also demonstrates that Yb, Er and Sc can be completely intersoluble to form L1 2 , and the lattice parameter has been significantly reduced by adding Sc. The misfit between FCC Al and L1 2  is 3.03%.  
     [0099] The morphology of fcc+L1 2  microstructure of Alloy 1E is shown in FIG. 21. This SEM secondary electron image indicates that the particle size of L1 2  precipitates is about 3 times larger than that of Alloy 1A, see FIG. 9, although the misfit of Alloy 1A is a little bit larger:3.735%. It indicates that devitrification of glassy state can significantly reduce the particle size of precipitation in comparison with that by crystallization directly from fast cooling at the same rate.  
     [0100] In review, the range of composition processing parameters, microstructure characterization and physical properties is summarized in the following Table 7:  
                           TABLE 7                                      Composition                                 Transition Metals                                     and Mg, Li       Processing Parameters                                         Late       FCC       thermo-                                                 Rare   TM for       elem.       thermal   mechanical           Earth   glass   L12   For       aging temp.   hot                                                     Elem.   forming   modifier   misfit       glass   for thermal   extrusion           (1)   (2)   (3)   (4)   at. %   RSP   devitrifying   parameters                                                             Maximum   Er, Lu,   Ni, Co,   Sc, Y   Mg, Zn,   2%&lt;sum(group   cooling   &gt;300° C.           Ranges   Yb, Tm   Fe, Cu   Ti, Zr,   Ag   (1))&lt;15%   rate           U       V, Cr,       2%&lt;sum(group   &gt;10 4 ° C./                   Mn, Li,       (2))&lt;7%   sec                   Nb       sum(group                           (3))&lt;5%       Preferred   Er, Lu,   Ni, Co,   Sc   Mg   5%&lt;sum(group   cooling   &gt;325° C.   &gt;350° C.       Ranges   Yb, Tm               (1))&lt;12%   rate   &lt;475° C.   &lt;450° C.                           2%&lt;sum(group   &gt;10 4 ° C./       5, 9, 11:1                           (2))&lt;7%   sec       ratio                           Sc&lt;4%                           Mg&lt;6%       Optimum   Er, Yb   Ni, Co,   Sc       7%&lt;sum(Er, Yb,   &gt;10 5 ° C./   &gt;370° C.   400° C.       Ranges                   Sc)&lt;8%   sec,   &lt;450° C.   11:1 ratio                           3%&lt;sum(Ni, Co)   atomized                           &lt;6%   powder                               with                               −325 mesh       Examples   Er, Yb   Ni, Co   Sc, Y   Mg   7%&lt;sum(Er, Yb)   about   425° C.   400° C.                           &lt;8% sum(Ni, Co)   10 6 ° C./       11:1 ratio                           &lt;5.5% sum(Sc, Y)   sec                           &lt;5%, Mg&lt;3%                         Preferred Methods:           Atomization and consolidation           with extrusion                                     Properties                         high           temperature                                 microstructure characterization (vol %)   room   UTS of                                     amorphous   devitrified state   temperature   275-410                                                         fcc   glass   others   fcc   L12   glass   others   UTS   Mpa                                                                         Maximum   &lt;15%   bal.   &lt;10%   bal.   &gt;10%   &lt;15%   &lt;20%   &gt;500Mpa   &gt;250° C.           Ranges           Preferred   &lt;8%   bal.   &lt;3%   bal.   &gt;20%-   &lt;8%   &lt;15%   &gt;800Mpa   &gt;275° C.           Ranges                   25%           Optimum   &lt;5%   bal.   &lt;1%   bal.   &gt;30%   &lt;3%   &lt;5%   &gt;1000Mpa   &gt;300° C.           Ranges           Examples       100% or       bal.   &gt;28%   &lt;3%   &lt;15-   490VHN                   &lt;20%           &lt;32%       20%   (1A)                                           which is                                           &gt;1000MPa                      
 
     [0101] Variations of the described aluminum alloy as well as the process for manufacture thereof and the product created by the process are 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.