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 L12 phase precipitates.

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 500MPa (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.

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'h 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 600Mpa (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.

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 L12 phase (fcc ordered) is one of the best-known precipitates to achieve a good combination of high strength and high toughness.

However, since crystalline aluminum has very limited solubility for most intermetallics, including L12 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.

For example, RSP crystalline aluminum alloys, not designed to form glass with RSP, utilize a metastable and coherent L12 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 L12 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 Al3Fe (which has non-cubic D03, not L12 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 Al3Fe precipitation.

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 L12 strengthening dispersion.

U. S. Pat. No. 6,248, 453B1 describes the use of an L12 strengthening dispersion for high strength aluminum alloys utilizing certain alloying components, but does not utilize a transition metal such as nickel and cobalt and an intermediate amorphous state as a means to refine the L12 microstructure.

Finally, the high strength and high temperature aluminum alloys referenced have not found widespread application due to 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 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 L12 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 103°C/sec, preferably at least 104°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/L12 crystalline microstructure with at least about 30% fcc phase in volume and at least about 10% L12 phase in volume. Thermal treatment, thermomechanical

treatment, or a combination of both, are the preferred methods to produce the final crystalline microstructure.

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 L12 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 & 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 L12 phase intermetallic dispersion. Accordingly, a process and an aluminum alloy product combining a sufficient amount of stable L12 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.

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.

The selection of alloying components is based on the function or criteria of (1) good L12 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/Ll2 coherency with low misfit which lead to a fine microstructure and good coarsening resistance at high temperature.

Lanthanide rare earth L12 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 L12 formation during devitrification. Scandium (Sc) is a strong L12 former which can be used in conjunction with rare earth components to promote L12 formation and reduce the misfit between fcc and L12 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/Ll2 misfit.

With efficient precipitation strengthening by the L12 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.

Thus, it is an object of the invention to provide a new class of high-strength L12 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.

A further object of the invention is to provide aluminum alloys with usable strength at or above 300°C by selecting (1) slow diffusing L12 formers to prevent fast coarsening kinetics at high temperature, (2) appropriate matrix fcc and L12 components to reduce the lattice misfit between them to further slow down the coarsening kinetics and to promote the initial finer L12 morphology.

Another object of the invention is to utilize the cubic crystal structure of L12 phase to promote the interfacial coherency with an aluminum fcc matrix and the inherent ductile

behavior of the L12 crystalline phase itself, due to high number of slip systems in cubic L12 structure, to maintain reasonable overall alloy ductility.

A further object of the invention is to combine alloying components so that the L12 formation is enhanced while the likelihood of forming other undesirable intermetallics is reduced, in order to further promote alloy ductility of an aluminum alloy.

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 L12 formers in the solution before a subsequent devitrification process.

A further object of the invention is to utilize the intermediate, highly supersaturated amorphous state, slow L12 coarsening kinetics, and low L12/fcc misfit to produce a fine microstructure in a strong aluminum alloy.

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.

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 L12 dispersion strengthened aluminum alloy fcc crystalline structure after devitrification.

These and other objects, advantages and features will be set forth in the detailed description which follows.

BRIEF DESCRIPTION OF THE DRAWING In the detailed description which follows, reference will be made to the drawing comprised of the following figures: Figure 1 is a chart setting forth criteria for the design of aluminum alloys of the invention ; Figure 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 AINi associate); Figure 3 is an assessed Al-Yb phase diagram in comparison with experimental data;

Figure 4a is a graph of the enthalpy of formation of cobalt and a rare earth, i. e. Er and Yb predicted by quantum mechanical calculation; Figure 4b is a graph of the mixing enthalpy of the materials set forth in Figure 4a; Figure 4c is a graph of the molar volume of the materials set forth in Figure 4a; Figure 5 is an X-ray diffractogram of the alloy of Example lA as melt-spun with positions of fcc pure aluminum reflections indicating a fully amorphous state; Figure 6 is an isochronic DSC thermogram of the alloy of Example lA indicating devitrification steps upon heating; Figure 7 is an X-ray diffractogram of the alloy of Example lA after devitrification at 425°C for 22 hours, with positions of reflections of pure fcc Al, Al3Yb, and Al3Er phases, indicating the desired phases: FCC+L12 ; Figure 8 is a SIMS elemental map of devitrified alloy of Example 1A indicating Er and Yb partition together to L12 phase, and NI and Co which do not partition to Llz ; Figure 9 is a scanning electron microscope (SEM) secondary electron image of devitrified alloy of Example 1A indicating submicron morphology; Figure 10 is a 3DAP microscopy reconstruction of A13 (Yb, Er) L12 phase of the alloy of Example 1A after devitrification at 424°C for 19 hours; Figure 11 is an X-ray diffractogram of the alloy of Example 2A as melt-spun, with positions of fcc pure Al and L12 reflections, indicating fully amorphous status; Figure 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, Al3Yb, and Al3Er phases, indicating the desired phases: FCC+L12 ; Figure 13 is an X-ray diffractogram of the alloy of Example 2B as melt-spun, with positions of fcc pure Al and L12 reflections, indicating fully amorphous status; Figure 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, Al3Yb, and Al3Er phases, indicating the desired phases: FCC+L12 wherein the lattice parameter of L1z is reduced by Sc; Figure 15 is an X-ray diffractogram of the alloy of Example 2C as melt-spun, with positions of fcc pure Al and L12 reflections, indicating almost fully amorphous status; Figure 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, Al3Yb, and Al3Er phases, indicating the desired phases: FCC+L12 wherein the lattice parameter of L12 is expanded by Mg;

Figure 17 is an X-ray diffractogram of the alloy of Example 2D as melt-spun, with positions of fcc pure Al and L12 reflections, indicating fully amorphous status; Figure 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 AI, Al3Yb, and Al3Er phases, indicating the desired phases: FCC+L12 and wherein the lattice parameter of fcc is expanded by Mg and that of L12 is reduced by Sc; Figure 19 is an X-ray diffractogram of the alloy of Example 1C as melt-spun, with positions of fcc pure Al and L12 reflections, indicating non-amorphous status due to the absence of transition metal elements; Figure 20 is an X-ray diffractogram of the alloy of Example 1C after devitrification at 425°C for 19 hours, with positions of reflections of pure fcc Al, Al3Yb, and Al3Er phases, indicating elements ER, Yb, and Sc can be completely intersoluble to form L12 and wherein the lattice parameter of L12 is reduced by Sc; and Figure 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 Figure 9) and indicating that passing through the glassy state can reduce the particle size of precipitation.

DESCRIPTION OF THE PREFERRED EMBODIMENTS General Summary 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 L12. (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 L12 phase material with limited residual amorphous or other phase material such as quasi-crystalline or ordered amorphous material.

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.

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.

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.

Development Technique Devitrification is accomplished via heat treating (thermal) processing, thermo- mechanical processing as well as other processing techniques.

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.

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 L12 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.

Figure 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.

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.

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's viscosity is set forth in Equation (1) : <BR> <BR> <BR> <BR> z G<BR> <BR> <BR> <BR> # = ####(x1#1o+x2#2o)#### (1)<BR> <BR> <BR> <BR> ')-2 The variables xi and Gm represent the mole fraction and molar Gibbs energy of the liquid, respectively. The viscosity of pure components, °, 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). ] Element Tm Vm #(Tm) CA E B #o (K) (10-6 (10-3 [10-7 (kJ/mol) (E/RTm) (10-3 Pa s) m3/mol) Pa s) (J/K mol1/3)1/2] 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. 7x10 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 Tm : melting temperature Vm molar volume at Tm TI (Tm) : viscosity at Tm CA : constant E: activation energy for viscous flow B : B=E/RTm 71,, : pre-exponential viscosity 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 AINi melts congruently at about 1650°C. Figure 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 AINi 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 AINi associate model is also plotted in Figure 2 for comparison. The calculation shows a significant increase in the viscosity due to the formation of liquid AINi associates.

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.

A model to predict glass transition temperature, Tg, has been developed for the Al-TM- RE system. The model is based on the parabolic relationship between Tg 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):

i = Al, Mg, Ni, Co, Er, Yb, Sc, Y, etc.

The coefficient for each element, K ;, 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 Tg. 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.

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 L12 phase dispersion in the AI matrix after devitrification of the glassy AI, as long as a sufficient volume fraction of L12 and low misfit with the Al matrix and thus low coarsening rate can be obtained.

The elements Er, Lu, Tm and Yb are known as the only stable rare earth (RE) L12 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 L12 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 L12 formers rather than Tm and Lu.

Table 2. Stable L12 formers and their properties. Element Melting Costa L12 Lattice Density L12 Max. Point Parameter Temp. Er 1529°C $725/kg Al3Er 4. 226 A 5.507 g/cm3 1070°C Lu 1663°C $7500/kg A13Lu 4. 186 A 5.783 g/cm3 1200°C Tm 1545°C $6500/kg Al3Tm 4.203 Å 5.591 g/cm3 Yb 819°C $1600/kg Al3Yb 4.202 Å 5.682 g/cm3 980°C Sc 1541°C $18, 000/kg Al3Sc 4. 101 Å 3.03 g/cm3 1320°C a For 1-5 kg cast metal ingots from 99.9%-grade oxides Sc is the only transition metal (TM) element that can form a stable L12 with Al. In comparison with RE L12 formers, Sc can form L12 with a smaller lattice parameter of 4.101 A,

thus leading to smaller misfit between L12 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 Al3Sc, up to 12.5 at. %, i. e. Al0. 75Sc0. 25Tio.) 25. 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 m2/s for Sc, 6. 3#10-24 m2/s for Zr, and 2. 7x 10-25 m2/s for Ti. Adding Ti to AIsSc will thus reduce the coarsening rate of L12 precipitates. Moreover, addition of Ti decreases the lattice parameter of Al3 (Sc, Ti) and hence minimizes the lattice misfit with Al. The lattice parameters for seven hypothetical L12 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 L12 compounds show significant smaller lattice parameters than those stable L12 phases. Thus, alloys of the invention incorporate Yb, Er and Sc as base L12 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 L12 TM Compounds through First Principle Calculations at 0 K. TMs Hypothetic L12 Lattice Parameter, A Al3Ti 3. 967 Al3Zr 4. 085 A13V 4. 045 Al3Fe 3. 8005 Al3Ni 3. 8469 Al3Co 3. 7946 Al3Cu 3. 9363 It should be pointed out that the compound Al3Ti is metastable as the L12 phase but stable as D022. It can be transformed to the L12 structure if Al is substituted partially by Cr, Mn, Fe, Co or Ni.

For a multicomponent L12 phase, the lattice parameter a can be calculated based on a linear model similar to Vegard's law. As an example, we have the following equation for L12 with the formula Al3 (ErXYbyScz) is utilized: a [Al3 (ErxYbyScz] = x * a (Al3Er) + y * a (Al3Yb) + z * a (Al3Sc) (3)

Assuming L1 has a linear thermal expansion, its lattice parameter at elevated temperature can be calculated as follows (T in °C) : a (Llz, T) = a (Ll 2, RT) X [1 + (T-25) x « (L12)] (4) where a (Ll2) represents the thermal expansion estimated according to the following reciprocal relation: cr (M) Tu (5) a (Al3M) T », (M) where M represents the corresponding pure metal in A13M and Tm 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 LI 2 or meta-stable L12.

Table 4. Estimation of Thermal Expansion Coefficients and Lattice Parameters. Thermal Thermal Expansion, Expansion, Melting Thermal Lattice Lattice Melting 1/°K, 1/°K, LI 2 Point, Expansion, Parameter, Parameter, Element Point, °K 20-100°C 20-300°C Alloys °K 20-300°C A, 25°C A, 300°C Ni 1728 13. 3#10-6 14. 4#10-6 Al3Ni 913 27. 254x 10-6 3.8639 3.8929 Co 1768 12. 5#10-6 13. 6#10-6 Al3Co 1408 17. 077x10-6 3.8114 3.8293 Yb 1092 25x10-6 25x10-6 Al3Yb 1253 21. 788x 10-6 4.202 4.2272 Er 1802 9. 2x 10-6 9. 2x10-6 Al3Er 1343 12. 344x 10-6 4.226 4.2403 Sc 1814 12x 10-6 12x 10-6 Al3Sc 1593 13.665 4.101 4.1164 Ti 1943 8. 9x10-6 9. 2#10-6 Al3Ti 1623 11.014 3.967 3.979 Y 1795 10. 8x 10-6 10. 8x 10-6 A13Y 1253 15.472 4.299 4.3173 A1 25.3#10-6 4. 0497 4. 0779 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 A. Further, every one atom percent of Mg will increase the linear coefficient of thermal expansion of Al by 1. 179X 10-7/°C.

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 Figure 3. The L12 Al3Yb is the strengthening phase that is used in alloys of the invention. According to Figure 3, Al3Yb 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.

Quantum mechanical calculations have been performed in Al3Er-Al3Co and Al3Yb- Also systems. The predicted enthalpy of formation, mixing enthalpy and molar volume in the corresponding systems are plotted in Figure 4. According to Figure 4, there is a strong repulsive Co-RE interaction indicating negligible Co solubility in L12 A13RE. This sets a challenge in producing good glass forming alloys through conventional technology in the same composition intended to precipitate a significant amount of L12 precipitation through devitrification.

Experimental Results 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 L12 after devitrification in the AI 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 L12 phases and the misfit between L12 and Al matrix.

Table 5 Summary of Composition and Microstructure of A1 Alloys No. Alloy Composition As-spun status Main phases Lattice Parameter Misfit at. % after devitrfication L12, a (A) fcc, a (A) % 1A Al-2. 5Ni-2.5Co fully glass fcc, Ll2+ 4.215 4.063 3.735 -3. 5Yb-3. 5Er minor x phase 1 B Al-2. 75Ni-2.75Co-2. 17Yb#fully glass fcc, L12+Al3Y+ 4.217 4.050 4.123 -2. 17Er-2. 17Y minor x phase 1C Al-6Yb-6Er fcc, L12+ fcc, L12 4.208 4.057 3.726 some glass 1D Al-4Yb-4Er-4Y fcc, L12 + fcc, L12 + Al3Y 4.229 4.059 4.192 some glass 1E Al-4, 5Yb-4.5Er-3Sc fcc, L12 + fcc, L12 4.183 4.060 3.032 some glass 1F Al-12Er-3Mg fcc, L12 + fcc, L12 4.216 4.059 3.865 some glass 1G Al-12Er-6Mg fcc, L12 + fcc, L12 4.217 4.065 3.739 some glass 1H Al-12Yb-3Mg fcc + fcc, L12 4.216 4.076 3.435 some glass 1I Al-12Yb-6Mg fcc, L12 + fcc, L12 4.191 4.061 3.201 some glass 2A Al-2Ni-2Co-4Yb-4Er almost fully fcc, L12 + 4.201 4.049 3.754 glass+minor fcc minor x phase 2B Al-2. 5Ni-2. 5Co-2Yb- almost fully fcc, L12 + 4.150 4.048 2.520 2Er-3Sc glass+minor fcc minor x phase 2C Al-3Mg-2. 5Ni-2. 5Co- mostly glass + fcc, L12 + 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, L12 + 4.155 4.057 2.416 2.5Yb-1. 5Er-3Sc minor x phase

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 L12 crystalline phases although they are partially amorphous in the as melt-spun conditions. These alloys include alloy 1C, 1D, 1E, 1F, 1G, 1H and 11. 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 L12 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 L12 already present in the as-spun condition.

Following is a detailed demonstration of characteristics of alloys in Table 5.

Example 1 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-Ka radiation) of as-spun Alloy 1A in Figure 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 Figure 6 indicates three stages of glass devitrification at temperatures near 200°C, 315°C and 370°C.

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.

Figure 7 shows an X-ray diffractogram of Alloy 1A after devitrification at 425°C for 22 hours. At this point transformations 1-111 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 A) are indicated in red. Within the accuracy of this experiment (that is 0. 01° at 20 = 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 L12 structure. The peak positions for L12 Al3Yb (lattice parameter 4.202 A marked in green, and that for L12 Al3Er (lattice parameter 4. 215 A) are marked in brown. It can be seen that the measured peak positions are perfectly matched by L12 structure reflections. Therefore, Alloy 1A mainly consists of FCC AI and L12 phases as devitrified microstructure. It is also noticed that there is a minor third phase x, which crystalline structure is to be identified.

The FCC Al+L 12 microstructure of Alloy 1A was investigated by the secondary ion mass spectrometry (SIMS) technology. Figure 8 indicates that Yb and Er partition together into L12 precipitates, while Ni and Co have limited partitions.

The morphology of FCC+L12 microstructure of Alloy 1A is shown in Figure 9.

This SEM secondary electron image indicates that the majority of L12 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 L12 precipitates with thermomechanical processing.

The three dimension atom probe microscopy (3DAPM) analysis confirms that L12 exists in Alloy 1A with formula Alo. 72 (Yb, Er) 0. 28. Yb and Er have about the same site fraction. Ni and Co are negligible in L12 phase. See Figure 10.

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 1A has demonstrated a hardness of 490 VHN, which is equivalent to the room temperature tensile strength exceeding 1000 MPa according to Inoue'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 L12 Table 6 Microhardness testing of Alloy 1 A Condition Hardness (VHN) As-spun 369 177°C 1hr 276 177°C 10hrs 310 177°C 20 hrs 335 177°C 20hrs then 288°C 2hrs. 490 177°C 20hrs then 316°C 2hrs. 340 177°C 20hrs then 400°C 2hrs. 275 Example 2 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-Ka radiation) of as-spun Alloy 2A in Figure 11 shows a broad

peak characteristic of glass, and the ribbons show reasonable good ductility, also characteristic of the glassy state.

Figure 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 A) 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 L12 structure. The peak positions for L12 Al3Yb (lattice parameter 4.202 A) are marked in green, and that for L12 Al3Er (lattice parameter 4.215 A) are marked in brown. It can be seen that the measured peak positions are perfectly matched by L12 structure reflections. Therefore, Alloy 2A mainly consists of FCC Al and L12 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 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-Ka radiation) of as-spun Alloy 2B in Figure 13 shows a broad peak characteristic of glass, and the ribbons show reasonable good ductility, also characteristic of the glassy state.

Figure 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 A) 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 L12 structure. It can be seen that the measured peak positions are matched by L12 structure reflections with lattice parameter 4.150 A. It demonstrates that Yb, Er and Sc can be completely intersoluble to form L12, and the lattice parameter has been

significantly reduced by adding Sc, which is another means of reducing misfit to keep the L12 precipitates coarsening resistant. The misfit between FCC Al and L12 is 2.52% for Alloy 2B.

Example 4 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-Ka radiation) of as-spun Alloy 2C in Figure 15 shows a broad peak characteristic of glass, and the ribbons show reasonable good ductility, also characteristic of the glassy state.

Figure 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 A) are indicated in red. The measured peak positions indicate that the lattice space of Al has been dilated to 4.060 A, reflecting another feature of the invention: adding Mg to expand the lattice space of Al and thus decrease the misfit between AI and L12. 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 L12 structure. It can be seen that the measured peak positions are perfectly matched by L12 structure reflections. Therefore, Alloy 2C mainly consists of FCC Al and L12 phases as devitrified microstructure.

Example 5 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-Ka radiation) of as-spun Alloy 2D in Figure 17 shows a broad peak characteristic of glass, and the ribbons show very good ductility, also characteristic of the glassy state.

Figure 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 A) are indicated in red. The measured peak positions indicate that the lattice space of Al has been dilated to 4.057 A. 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 L12 structure. It can be seen that the measured peak positions are matched by L12 structure reflections with lattice parameter 4.155 A. It demonstrates that Yb, Er and Sc can be completely intersoluble to form L12, 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 L12. 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 L12 phases as devitrified microstructure.

Example 6 Alloy IE 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-Ka radiation) of as-spun Alloy 1E in Figure 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.

Figure 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+Ll2. It also demonstrates that Yb, Er and Sc can be completely intersoluble to form L12, and the lattice parameter has been significantly reduced by adding Sc. The misfit between FCC Al and L12 is 3.03%.

The morphology of fcc+L12 microstructure of Alloy 1E is shown in Figure 21. This SEM secondary electron image indicates that the particle size of L12 precipitates is about 3 times larger than that of Alloy 1A, see Figure 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.

In review, the range of composition processing parameters, microstructure characterization and physical properties is summarized in the following Table 7: Table 7 Composition Processing Parameters microstructure characterizatlon vol %) Properties Rare thermo- Earth Trane (tioh MetaIS mechan- Elem. and M, LI lass thermal ical amorphous devitrified state Late FCC high TM for elem. aging tempera- glass L12 For temp. hot room tun3 UTS fonming modi-mistit RSP forthermal extruslon tempera-of275- (1) (2) fier (3) (4) at. % devitrifying parameter fcc glass others fcc L12 glass ottlers ture UTS 410 Mpa 2% <sum (group Suc, Y (1)) <10% Tl, Zr, 2% esum (group cooling Er, Lu, TI, Zr, Maxtmum NI, Co, Mg, Zn, (2)) <7°/ rate >3, C <15° bal. <10°/bal. >109ro <16°Ao <20% >500MPa >250°C Rangos U : Tm Fe Gu Mn,Li,' sum (group >104°C Nb (3)) <5% sec sum (group (4)) <7% 5% <sum (group (1)) <12% cooling >350°C Preferred Er, Lu, 2% <sum (group rate >325°C <450°C ° o >20-° ° > Ranges b ; Tm N, Co Sc. M9 (2)) <7% >10°°C <475"C 5. 8, 11 : 1 8/° bal. <3/a bl. 26% <8/0 <15/>800MPa 215°C St<4%/ec ratio 8050c 105C sec, 7% <sum (Er. Yb, S atomize Optlmum c8% d >370°C i00°C ° o'> < ° <5'/0 >100. OMP >300°C Ranges 3% <sum (Ni, Co) < powder <450°C 11 : 1 ratlo /° bai. <1/° bal. 3096 3/0 6% with 325 mesh 7% <sum (EO, Yb) < 9oVHN ° 100 1 Exam <sum (Er. Yb) <.-,,, 480VHN about % /0 11 : 1 ratio <20 <32/0 20/> sum (Sc, Y) <5%, sec % 1000 Mg<3% Preferred Methods : tomizatlon and consolidation" ith extrusidn

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.