Classifications

• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
  • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
  • C22F1/04—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of aluminium or alloys based thereon
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C21/00—Alloys based on aluminium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C45/00—Amorphous alloys
  • C22C45/08—Amorphous alloys with aluminium as the major constituent

Definitions

• 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.
• 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 Ll 2 phase precipitates.
• 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, Nolume 2, 10 th Edition, ASM International].
• 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).
• 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.
• intermetallic precipitation dispersions 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 Ll 2 phase (fee ordered) is one of the best-known precipitates to achieve a good combination of high strength and high toughness.
• RSP crystalline aluminum alloys not designed to form glass with RSP, utilize a metastable and coherent Ll 2 dispersion (about 1%) which converts into other phase structures at high temperature.
• Ll 2 dispersion about 17% which converts into other phase structures at high temperature.
• 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 Ll 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 Ll 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.
• 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
• U.S. Pat. No. 6,248,453B1 describes the use of an Ll 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 amo ⁇ hous state as a means to refine the
• the present invention is directed to a new class of aluminum alloys characterized by formation from an intermediate amo ⁇ hous state to a final face centered cubic (fee) 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 fee matrix strengthened by a precipitate phase with high cubic symmetry, preferably the Ll phase.
• An appropriate melt of aluminum with crystalline formers is first processed to achieve an intermediate amo ⁇ hous state to preserve the crystalline formers.
• the preferred method to achieve a primarily (above 70% in volume) amo ⁇ hous 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 amo ⁇ hous microstructure through a solid-state process, such as mechanical milling, may also be used.
• the intermediate amo ⁇ hous alloys are then devitrified to a final primarily (above 70% in volume) fcc/Ll crystalline microstructure with at least about 30% fee phase in volume and at least about 10% Ll 2 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.
• 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 amo ⁇ hous or glassy state and then subsequently to a crystalline state by a devitrification process.
• a melt which is converted to an amo ⁇ hous or glassy state and then subsequently to a crystalline state by a devitrification process.
• 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.
• 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 amo ⁇ hous or glassy material from the melt.
• Other materials included, optionally, in the melt comprise additional transition metals, lithium (Li) and magnesium (Mg).
• Si scandium
• Y yttrium
• Ti titanium
• Zr zirconium
• V vanadium
• Cr chromium
• Mn manganese
• Li lithium
• magnesium (Mg), zinc (Zn), silver (Ag) and niobium (Nb) may be inco ⁇ orated into the melt or alloy to accommodate fee misfit in the crystalline state.
• the selection of alloying components is based on the function or criteria of (1) good
• Ll 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/Ll 2 coherency with low misfit which lead to a fine microstructure and good coarsening resistance at high temperature.
• Lanthanide rare earth Ll 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 Ll 2 formation during devitrification.
• Scandium (Sc) is a strong Ll 2 former which can be used in conjunction with rare earth components to promote Ll 2 formation and reduce the misfit between fee and Ll 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 fee aluminum to further reduce the fcc/Ll misfit.
• 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.
• a further object of the invention is to provide aluminum alloys with usable strength at or above 300°C by selecting (1) slow diffusing Ll 2 formers to prevent fast coarsening kinetics at high temperature, (2) appropriate matrix fee and Ll 2 components to reduce the lattice misfit between them to further slow down the coarsening kinetics and to promote the initial finer Ll 2 mo ⁇ hology.
• Another object of the invention is to utilize the cubic crystal structure of Ll 2 phase to promote the interfacial coherency with an aluminum fee matrix and the inherent ductile behavior of the Ll 2 crystalline phase itself, due to high number of slip systems in cubic Ll 2 structure, to maintain reasonable overall alloy ductility.
• a further object of the invention is to combine alloying components so that the Ll 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.
• 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 amo ⁇ hous aluminum alloy structure, preserving the Ll 2 formers in the solution before a subsequent devitrification process.
• a further object of the invention is to utilize the intermediate, highly supersaturated amo ⁇ hous state, slow Ll 2 coarsening kinetics, and low Ll 2 /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 amo ⁇ hous 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 Ll 2 dispersion strengthened aluminum alloy fee crystalline structure after devitrification.
• 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 AlNi 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 1A as melt-spun with positions of fee pure aluminum reflections indicating a fully amo ⁇ hous state;
• Figure 6 is an isochronic DSC thermogram of the alloy of Example 1A indicating devitrification steps upon heating;
• Figure 7 is an X-ray diffractogram of the alloy of Example 1 A after devitrification at 425°C for 22 hours, with positions of reflections of pure fee Al, Al 3 Yb, and Al 3 Er phases, indicating the desired phases: FCC+L1 2 ;
• Figure 8 is a SIMS elemental map of devitrified alloy of Example 1A indicating Er and Yb partition together to Ll 2 phase, and NI and Co which do not partition to Ll 2 ;
• Figure 9 is a scanning electron microscope (SEM) secondary electron image of devitrified alloy of Example 1A indicating submicron mo ⁇ hology;
• Figure 10 is a 3DAP microscopy reconstruction of Al 3 (Yb,Er)Ll 2 phase of the alloy of Example 1 A after devitrification at 424°C for 19 hours;
• Figure 11 is an X-ray diffractogram of the alloy of Example 2 A as melt-spun, with positions of fee pure Al and Ll 2 reflections, indicating fully amo ⁇ hous status;
• Figure 12 is an X-ray diffractogram of the alloy of Example 2A after devitrification at
• Figure 13 is an X-ray diffractogram of the alloy of Example 2B as melt-spun, with positions of fee pure Al and Ll 2 reflections, indicating fully amo ⁇ hous status;
• Figure 14 is an X-ray diffractogram of the alloy of Example 2B after devitrification at
• Figure 15 is an X-ray diffractogram of the alloy of Example 2C as melt-spun, with positions of fee pure Al and Ll 2 reflections, indicating almost fully amo ⁇ hous status;
• Figure 16 is an X-ray diffractogram of the alloy of Example 2C after devitrification at
• 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 fee Al, Al Yb, and Al Er phases, indicating the desired phases: FCC+L1 2 and wherein the lattice parameter of fee is expanded by Mg and that of Ll 2 is reduced by Sc;
• Figure 19 is an X-ray diffractogram of the alloy of Example IC as melt-spun, with positions of fee pure Al and Ll 2 reflections, indicating non-amo ⁇ hous status due to the absence of transition metal elements;
• Figure 20 is an X-ray diffractogram of the alloy of Example IC after devitrification at
• Figure 21 is a scanning electron microscope (SEM) secondary electron image of a devitrified alloy of Example IC 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.
• SEM scanning electron microscope
• 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 amo ⁇ hous or glassy state followed by devitrification of the material from the amo ⁇ hous state to a mixed crystalline state comprised of face centered cubic (fee) and Ll 2 . (an ordered fee) state wherein the ratios of the crystalline states is within certain preferred ranges.
• the resultant alloy has at least about 30% by volume fee phase and at least about 10% by volume Ll 2 phase material with limited residual amo ⁇ hous or other phase material such as quasi-crystalline or ordered amo ⁇ hous material.
• the choice of starting materials may be varied as may the compounding processes, the glass phase formation processes and the devitrification processes.
• the amo ⁇ hous state there may be some crystalline material contained therein, but preferably no more than about 30% by volume in the fee phase.
• 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).
• 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.
• amo ⁇ hous state processes for mixing or forming the starting materials in the amo ⁇ hous 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 amo ⁇ hous state is preferably at least about 70% by volume and preferably greater.
• the amo ⁇ hous 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 fee 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.
• 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.
• processing techniques achieve a fully amo ⁇ hous 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 fee 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 Ll 2 allows the precipitate to be coherent or semi-coherent with the fee 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.
• transition metals such as Fe, Co, Ni, Cu
• TM transition metals
• RE Rare-earth elements
• Ce, Gd, Yb, Er Rare-earth elements with large atomic size exhibit low diffusivity in Al and thus retard crystal nucleation.
• 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):
• the variables x, and G m represent the mole fraction and molar Gibbs energy of the liquid, respectively.
• the viscosity of pure components, ⁇ ° is available from literature, see Table 1.
• T m melting temperature
• V m molar volume at T m
• T m viscosity at T ⁇
• 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.
• 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.
• 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 Figure 2 for comparison. The calculation shows a significant increase in the viscosity due to the formation of liquid AlNi 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.
• T g glass transition temperature
• the desired performance for applications 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 Ll 2 phase dispersion in the Al matrix after devitrification of the glassy Al, as long as a sufficient volume fraction of Ll 2 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) Ll 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 Ll 2 decomposition or congruent melting temperature.
• 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.
• Alloys of the invention utilize Yb and Er as preferred Ll 2 formers rather than Tm and Lu.
• Sc is the only transition metal (TM) element that can form a stable Ll 2 with Al. In comparison with RE Ll 2 formers, Sc can form Ll 2 with a smaller lattice parameter of 4.101 A, thus leading to smaller misfit between LI 2 and the Al matrix.
• TM transition metal
• 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. Alo. 5 Sco..25Tio.i2 5 .
• Ti has the lowest diffusion coefficient in solid Al among the transition metals.
• the diffusivity in Al at 300°C is 9.0x10 "20 m 2 /s for Sc, 6.3x10 "24 m 2 /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 Ll 2 precipitates.
• 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 Ll 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 Ll 2 compounds show significant smaller lattice parameters than those stable Ll 2 phases.
• alloys of the invention inco ⁇ orate Yb, Er and Sc as base Ll 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.
• the compound Al 3 Ti is metastable as the Ll 2 phase but stable as D0 22 . It can be transformed to the Ll 2 structure if Al is substituted partially by Cr, Mn, Fe, Co or Ni.
• the lattice parameter a can be calculated based on a linear model similar to Vegard's law.
• Ll 2 has a linear thermal expansion
• ⁇ (Ll 2 ) represents the thermal expansion estimated according to the following reciprocal relation: a(M) _ T m (Al 3 M)
• Thermodynamic modeling and assessments have been performed for Al-TM-RE systems with relevant elements inco ⁇ orated in alloys of the invention.
• the assessed Al-Yb binary phase diagram is shown in Figure 3.
• the Ll 2 Al 3 Yb is the strengthening phase that is used in alloys of the invention.
• 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 amo ⁇ hous state due to the larger driving force. In addition, the precipitation process can be controlled through devitrification. Quantum mechanical calculations have been performed in Al 3 Er-Al 3 Co and Al 3 Yb-
• the present invention alloys, through computational design of multi-component Al- TM-RE systems inco ⁇ orate, desired processing properties - glass forming ability and the desired microstructure - a fine dispersion of Ll 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 fee and the LI 2 phases and the misfit between Ll 2 and Al matrix.
• alloys with TM elements are fully glass or almost fully glass in the as melt-spun conditions. These alloys include alloy 1A, IB, 2 A, 2B, 2C and 2D. Alloys without Ni and Co primarily show FCC Al and Ll 2 crystalline phases although they are partially amo ⁇ hous in the as melt-spun conditions. These alloys include alloy IC, ID, IE, IF, IG, IH and II. 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 Ll 2 in the Al matrix.
• TM elements e.g. Ni and Co
• 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-K ⁇ 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.
• XRD X-ray diffraction
• DSC differential scanning calorimetry
• 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.
• Alloy 1A mainly consists of FCC Al and Ll 2 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 AI+LI 2 microstructure of Alloy 1A was investigated by the secondary ion mass spectrometry (SIMS) technology. Figure 8 indicates that Yb and Er partition together into LI 2 precipitates, while Ni and Co have limited partitions.
• 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, "Amo ⁇ hous, nanoquasicrystalline and nanocrystalline alloys in Al-based systems", Progress in Materials Science 43 (1998) 365-520). This result indicates high efficient strengthening of Ll 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-K ⁇ radiation ) of as-spun Alloy 2 A 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 2 A after devitrification at 425°C for 19 hours. At this point the ribbon is fully devitrified in the sense that no significant amount of amo ⁇ hous 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 LI 2 structure.
• Alloy 2A mainly consists of FCC Al and Ll 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.
• 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-K ⁇ 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 amo ⁇ hous material can be detected by X-ray diffraction (XRD).
• XRD X-ray diffraction
• 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.
• 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-K ⁇ 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
• 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-K ⁇ 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
• 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 Ll 2 .
• Alloy 2D achieved the lowest misfit so far:2.42% without losing the glass forming ability.
• Alloy 2D mainly consists of FCC Al and Ll 2 phases as devitrified microstructure.
• 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 IE was characterized by X-ray diffraction. The diffractometer trace (Cu-K ⁇ radiation ) of as-spun Alloy IE in Figure 19 shows crystalline reflection peaks, indicating non-amo ⁇ hous 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 IE after devitrification at 425°C for 19 hours. This diffractogram shows that Alloy IE only has two phases: fee A1+L1 2 . It also demonstrates that Yb, Er and Sc can be completely intersoluble to form Ll 2 , and the lattice parameter has been significantly reduced by adding Sc. The misfit between FCC Al and Ll 2 is 3.03%.

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