Grain-boundary strengthened cast aluminum alloys for enhanced creep resistance

Abstract

An alloy composition includes: 4.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1 -0.5 wt. % Mn; 0.1 -0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; balance Al, wherein the alloy with the average grain diameters ranging between 10 and 1000 μm produces minimum creep rates from 10-10 to 3x10-9 s-1 at stresses up to 30 MPa, and from 10-10 s-1 to 2x10-8 s-1 at stresses up to 70 MPa at 300 °C. A method of making an alloy is also disclosed.

Description

GRAIN-BOUNDARY STRENGTHENED CAST ALUMINUM ALLOYS FOR ENHANCED CREEP RESISTANCE STATEMENT REGARDING FEDERALLY SPONSORED RESEARCH AND DEVELOPMENT

[0001] The technologies described herein were developed with governmentsupport under Contract No. DE-AC05-00OR22725 awarded by the U.S. Department of Energy. The government has certain rights in the described technologies. CROSS-REFERENCE TO RELATED APPLICATIONS.

[0002] This application claims priority to 63 / 559,270, filed on February 29,2024, entitled “GRAIN-BOUNDARY STRENGTHENED CAST ALUMINUM ALLOYS FOR ENHANCED CREEP RESISTANCE”, the entire disclosure of which incorporated herein by reference. FIELD OF THE INVENTION

[0003] The present invention relates generally to aluminum alloys, and moreparticularly to aluminum alloys having enhanced creep resistance. BACKGROUND OF THE INVENTION

[0004] Lightweight, high-strength aluminum alloys are widely used in theautomotive and aerospace industries. It is widely recognized that most aluminum alloys exhibit poor mechanical properties at temperatures above ~200 °C, which hinders their use for load-bearing high-temperature applications. A new alloy design strategy has recently emerged, according to which the metastable primary -Al2Cu) in cast Al-Cu-based alloys can be retained up to350 °C. It has been demonstrated that slow-diffusing transition metals, such as Zr-Al2Cu up to 350 °C by their thermally activated segregation to -Al2 -Al interfaces in Al-Cu alloys.

[0005] Extensive efforts have been made to characterize the creep behaviorof cast Al-Cu-Mn-Zr alloys and identify new microstructural stabilization strategiesthat can lead to further improved alloy creep performance. The -Al2Cu precipitates,owing to their platelet morphology and high aspect ratio (~ 300-400 nm in length and~ 5-30 nm in thickness), can effectively suppress thermally activated dislocation1 168773228.1climb. As a result, dislocation motion within the grain cores is restricted at stressesup to the Orowan stress since these precipitates are also resistant to shearing. Varying grain size affects the creep strain rate in diffusional and dislocation creep regimes, with finer grains producing higher creep strain rates. Higher creep strain rates for fine-grained microstructures in the diffusional creep regime were attributed to the accelerated vacancy diffusion along grain boundaries, that is, enhanced Coble creep. Enhanced dislocation creep for fine-grained microstructures was associatedwith the increased vol -PFZ where plastic strain localizes, accelerating theoverall creep strain rate.

[0006] Al-Cu-Mn-Zr alloys containing Ni and Co, commercially known asRR350, have shown higher creep resistance as compared to Al-Cu-Mn-Zr alloysbecause Ni / Co-rich intermetallic precipitates that form during its solidification ( )effectively retard grain-boundary sliding needed to accommodate both diffusionaland PFZ-dislocation creep and ( -Al, producingcomposite strengthening effects. Unlike Ni / Co- -Al2Cuphase, obtained at high volume fractions (~ 6%) by increasing Cu content of Al-Cu-Mn- -Al, has shown no beneficial effect onalloy creep performance in both diffusional and dislocation creep regimes. This is -Al2Cu precipitates have low coarsening resistance, unlike Ni / Co-rich -PFZ readily expands -Al2Cu but not around Ni / Co-rich precipitates. -PFZ becomes a weak region in the microstructure, where localized plastic strain occurs at -Al2Cu ineffective in both retarding grain-boundary sliding and load- -Al. -PFZ-Al2Cu precipitates at grain boundaries, making the alloy more prone to cavitation in tension, further deteriorating the alloy creep performance.

[0007] Grain-boundary microstructure remains an essential consideration foralloy creep performance, given the difficulty of producing extremely coarse-grained aluminum alloys by conventional manufacturing processes. Also, Ni and Co, which together form coarsening-resistant intermetallic phases in RR350 alloy and significantly improve its creep resistance, add to the cost of the alloy. Further 2 168773228.1improvements are, therefore, required to make grain-boundary intermetallic phases effective in improving alloy creep resistance. 3 168773228.1SUMMARY OF THE INVENTION

[0008] An alloy composition comprises:4.0-24.0 wt. % Cu 0.5-3.0 wt. % Fe 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti;balance Al, wherein the alloy has average grain diameters ranging between 10 and1000 μm and produces minimum creep rates from 10-10to 3x10-9s-1at stresses up to 30 MPa, and from 10-10s-1to 2x10-8s-1at stresses up to 70 MPa at 300oC.

[0009] The total amount of impurities can be present at less than 0.5 wt. %based on total weight of the alloy. The impurity can be Ag and can be present atless than 0.05 wt. %. Each impurity element other than Ag can be present at lessthan 0.1 wt. % based on the total weight of the alloy.

[0010] The alloy can comprise 0.05-0.15 wt.% Si based on the total weight ofthe composition. The alloy composition can have x= Fe + Ni + Co, and x = 0.5 to 3wt. % based on the total weight of the alloy. The ratio of Cu / x can be from 5:1 to 8:1in wt. % based on the total weight of the alloy. The concentration of intragranular Cu(Cu content in the grain interiors) can be from 2.5 wt. % to 6 wt. %, based on thetotal weight of the alloy.

[0011] -Al2Cu can be within 0.005-0.04(0.5-4%) based on the total volume of the alloy. The alloy can comprise-Al2Cu phase precipitates and intergranular Al7Cu2(Fe,Co,Ni),Al9(Fe,Co,Ni)2 and Al3(Fe,Co,Ni)2 precipitates. The alloy can comprise in weightpercent based on the total -Al2Cu phase than -Al2Cu phase. The linear fraction of grain boundary being covered by intergranularAl7Cu2(Fe,Co,Ni), Al9(Fe,Co,Ni)2 and Al3(Fe,Co,Ni)2 precipitates can be at least 20%.

[0012] An alloy composition consists essentially of:4 168773228.14.0-24.0 wt. % Cu 0.5-3.0 wt. % Fe 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; balance Al, wherein the alloy with the average grain diameters ranging between 10 and 1000 μm and produces minimum creep rates from 10-10to 3x10-9s-1at stresses up to 30 MPa, and from 10-10s-1to 2x10-8s-1at stresses up to 70 MPa at 300oC.

[0013] An alloy composition consists of:4.0-24.0 wt. % Cu 0.5-3.0 wt. % Fe 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti;balance Al, wherein the alloy has the average grain diameters ranging between 10and 1000 μm and produces minimum creep rates from 10-10to 3x10-9s-1at stresses up to 30 MPa, and from 10-10s-1to 2x10-8s-1at stresses up to 70 MPa at 300oC.

[0014] A method of making an alloy, comprises the steps of:combining 4.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1-0.5 wt. % Mn; 0.1- 0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; and balance Al; melting the composition; casting and solidifying the alloy to create a solidified alloy; and, 5 168773228.1heat-treating the solidified alloy, the heat treating comprising at least one selected from the group consisting of solutionizing heat treatment, quenching in hot water, artificial aging and preconditioning heat treatment.

[0015] The heat-treating can comprise a solutionizing heat treatment. Thesolutionizing heat treatment can comprise heating to 530-545oC for at least 4 hours. The heat-treating step can comprise quenching in hot water having a temperature of60-85 oC. The heat-treating step can comprise aging at 230-240 oC for at least 5hours. The heat-treating step can comprise preconditioning at 300 oC for at least100 hours. 6 168773228.1BRIEF DESCRIPTION OF THE DRAWINGS

[0016] There are shown in the drawings embodiments that are presentlypreferred it being understood that the invention is not limited to the arrangementsand instrumentalities shown, wherein:

[0017] Figure 1(A) is an electron backscatter diffraction (EBSD) all-Euler anglemap showing the grain structure of ACMZ and FIG.1(B) is a plot of average grain diameter measured for all alloys. EBSD scans were made on alloys after heat treatment.

[0018] Figure 2 shows the XRD spectra for the six alloys and the peaksassigned to the associated phases.

[0019] Figure 3(A) is a bright-field TEM micrograph, acquired in two-beam-Al2Cu precipitates at graininteriors of base ACMZ alloy. Figures 3(B,C,E-H) show backscattered electron SEMmicrographs showing grain-boundary microstructures of FIG. 3(B) ACMZ; FIG. 3(C)ACMZ-Co, FIG.3(E) ACMZ-Ni; FIG.3(F) ACMZ-Co / Ni; FIG.3(G) ACMZ-Fe, andFIG. 3(H) ACMZ-Cu / Fe. FIG. 3(D) is a phase-contrast EBSD map showing thedistribution of intermetallic phases indexed as Al7Cu2Fe and Al9Co2 in ACMZ-Co alloy.

[0020] Figure 4 shows double-logarithmic plots of steady-state strain rateversus compressive stress at 300 °C for ACMZ alloys (base and modified).

[0021] Figures 5(A-F) show SEM backscattered electron micrographs showingtypical post-creep microstructures of five ACMZ alloys (base and modified): FIG. 5(A) ACMZ; FIG.5(B) ACMZ; FIG.5(C) ACMZ-Fe, FIG.5(D) ACMZ-Cu / Fe; FIG. 5(E) ACMZ-Co and FIG.5(F) ACMZ-Co / Ni.

[0022] Figures 6(A-F) are EBSD maps acquired from base ACMZ alloy beforeand after creep deformation (with ~ 10% total strain accumulated within the applied stress range of 20-133 MPa during the stress-jump creep test), where FIG.6(A) is an inverse pole figure, before creep deformation; FIG.6(B) is a kernel average misorientation (KAM) map, before creep deformation; FIG.6(C) is a grain reference orientation deviation (GROD) map, before creep deformation; FIG.6(D) is an inverse pole figure, after creep deformation; FIG.6(E) is a kernel average misorientation 7 168773228.1(KAM) map, after creep deformation; and FIG.6(F) is a grain reference orientation deviation (GROD) map, after creep deformation.

[0023] Figure 7 is a double-logarithmic plot of steady-state strain rate versuscompressive stress at 300 °C for the three unmodified ACMZ alloys with various Cu concentrations, and therefore, with different fraction ƒ of grain-boundary -Al2Cu precipitates; also shown are modified ACMZ-Cu / Fe, RR350, and Al-1Mn-0.3Zr- 0.3Er-0.05Si (wt.%).

[0024] Figures 8(A,B) show SEM backscattered electron micrographs showingthe typical post-creep microstructures of base ACMZ, with FIG.8(B) showing a higher-magnification view of an area within the dotted rectangle in FIG.8(A). Figure 8(C) is a schematic diagram illustrating the main microstructural factors controlling the creep behavior of ACMZ alloy and the presence of increased GND density in the -PFZ.

[0025] Figure 9(A-B) are plots of equilibrium fraction of phases vs.temperature, as calculated using Thermo-Calc (TCal 8 database) for FIG.9(A) ACMZ-Co alloy and FIG.9(B) ACMZ-Cu / Fe compositions.

[0026] Figures 10(A-D) are schematic illustrations of primary microstructuralconstituents before and after the compressive creep in FIG.10(A) baseline ACMZbefore creep; FIG. 10(B) baseline ACMZ after creep; FIG. 10(C) Fe / Ni / Co-modifiedACMZ alloys before creep; and FIG.10(D) shows Fe / Ni / Co-modified ACMZ alloysafter creep.

[0027] Figure 11 is a plot of minimum strain rate versus applied stress forbaseline and modified ACMZ alloys.

[0028] Figures 12(A-F) are EBSD all-Euler angle maps showing the grainstructures of FIG.12(A) ACMZ; FIG.12(B) ACMZ-Fe; FIG.12(C) ACMZ-Cu / Fe; FIG. 12(D) ACMZ-Ni; FIG.12(E) ACMZ-Co; and FIG.12(F) ACMZ-Co / Ni.

[0029] Figure 13(A) is a phase-contrast map and FIG. 13(B) is an inverse polefigure EBSD map, acquired from ACMZ-Co alloy after preconditioning heat treatment. 8 168773228.1

[0030] Figures 14(A-D) shows a comparison of the typical high-magnificationmicrostructures of FIG.14(A) baseline ACMZ; FIG.14(B) ACMZ-Ni; FIG.14(C) ACMZ-Fe; and FIG.14(D) ACMZ-Cu / Fe.

[0031] Figure 15 is an SEM backscattered electron micrograph showing thetypical post-creep microstructure of ACMZ-Ni crept in compression at 300 °C for a total strain of ~10%.

[0032] Figures 16(A-B) provides a comparison of the typical microstructures ofFIG.16(A) ACMZ-Co / Ni and FIG.16(B) RR350. 9 168773228.1DETAILED DESCRIPTION OF THE INVENTION

[0033] An alloy composition according to the invention comprises:4.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti;balance Al, wherein the alloy has average grain diameters ranging between 10 and1000 μm and produces minimum creep rates from 10-10to 3x10-9s-1at stresses upto 30 MPa, and from 10-10 s-1 to 2x10-8 s-1 at stresses up to 70 MPa at 300 oC. Thealloy composition can consist essentially of these elements.

[0034] The alloy composition can comprise 4.0-24.0 wt. % Cu. The alloycomposition can comprise 4, 5, 6, 7, 8, 9, 10, 11, 12, 13, 14, 15, 16, 17, 18, 19, 20, 21, 22, 23, or 24 wt.% based on the total weight of the alloy. The wt. % Cu can be within a range of any high value and low value selected from these values.

[0035] The alloy composition can comprise 0.5-3.0 wt. % Fe. The alloycomposition can comprise 0.5, 0.6, 0.7, 0.8, 0.9, 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9 or 3.0 wt. % Fe based on the total weight of the alloy. The wt. % Fe can be within a range of any high value and low value selected from these values.

[0036] The alloy composition can comprise 0.1-0.5 wt. % Mn. The alloycomposition can comprise 0.1, 0.15, 0.2, 0.25, 0.3, 0.35, 0.4, 0.45, or 0.5 wt. % Mn based on the total weight of the alloy. The wt. % Mn can be within a range of any high value and low value selected from these values.

[0037] The alloy composition can comprise 0.1-0.3 wt. % Zr. The alloycomposition can comprise 0.1, 0.125, 0.15, 0.175, 0.2, 0.225, 0.25, 0.275, or 0.3 wt. % Zr based on the total weight of the alloy. The wt. % Zr can be within a range of any high value and low value selected from these values. 10 168773228.1

[0038] The alloy composition can comprise 0-0.15 wt.% Si. The alloycomposition can comprise 0, 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.11, 0.12, 0.13, 0.14, or 0.15 wt. % Si based on the total weight of the alloy. The wt. % Si can be within a range of any high value and low value selected from these values. The alloy composition can have a Si content in the range of 0.05-0.15 wt.% to produce a more -Al2Cu phase.

[0039] The alloy composition can comprise 0-0.3 wt. % Mg. The alloycomposition can comprise 0, 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.11, 0.12, 0.13, 0.14, 0.15, 0.16, 0.17, 0.18, 0.19, 0.2, 0.21, 0.22, 0.23, 0.24, 0.25,0.26, 0.27, 0.28, 0.29 or 0.3 wt. % Mg based on the total weight of the alloy. The wt.% Mg can be within a range of any high value and low value selected from these values.

[0040] The alloy composition can comprise 0-0.3 wt. % Ti. The alloycomposition can comprise 0, 0.01, 0.02, 0.03, 0.04, 0.05, 0.06, 0.07, 0.08, 0.09, 0.1, 0.11, 0.12, 0.13, 0.14, 0.15, 0.16, 0.17, 0.18, 0.19, 0.2, 0.21, 0.22, 0.23, 0.24, 0.25, 0.26, 0.27, 0.28, 0.29 or 0.3 wt. % Ti based on the total weight of the alloy. The wt. % Ti can be within a range of any high value and low value selected from these values.

[0041] The alloy composition can comprise 0 – 2.5 wt. % Ni based on the totalweight of the alloy. The alloy composition can comprise 0, 0.1, 0.2, 0.3, 0.4, 0.5, 0.6, 0.7, 0.8, 0.9, 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, or 2.5 wt. % Ni based on the total weight of the alloy. The wt. % Ni can be within a range of any high value and low value selected from these values.

[0042] The alloy composition can comprise 0 – 2.5 wt. % Co based on thetotal weight of the alloy. The alloy composition can comprise 0, 0.1, 0.2, 0.3, 0.4, 0.5, 0.6, 0.7, 0.8, 0.9, 1.0, 1.1, 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 2.0, 2.1, 2.2, 2.3, 2.4, or 2.5 wt. % Co based on the total weight of the alloy. The wt. % Co can be within a range of any high value and low value selected from these values.

[0043] The ratio between Cu and combined Fe, Ni, and Co can be controlledto obtain a Cu concentration of at least 2.5 wt. % and up to ~6 wt. % in the grain interiors based on the total weight of the alloy, a maximum achievable amount in Al crystal lattice of the grain interiors. The Cu concentration in the Al crystal lattice of 11 168773228.1the grain interiors can be 2.5, 2.6, 2.7, 2.8, 2.9, 3.0, 3.1, 3.2, 3.3, 3.4, 3.5, 3.6, 3.7, 3.8, 3.9, 4.0, 4.1, 4.2, 4.3, 4.4, 4.5, 4.6, 4.7, 4.8, 4.9, 5.0, 5.1, 5.2, 5.3, 5.4, 5.5, 5.6,5.7, 5.8, 5.9 or 6.0 wt. % based on the total weight of the alloy, or can be within arange of any high value and low value selected from these values.

[0044] The ratio between Cu and combined Fe, Ni, and Co can also becontrolled to form grain-boundary phases containing at least one of Fe, Ni, and Co at a minimum amount of 1 wt.% based on the total weight of the grain-boundary phase. The alloy provides low solid-solubility and slow diffusivity of Fe, Ni, and Co to ensure coarsening resistance of grain-boundary phases, whereas too high Cu content and low Fe, Ni, and Co presence causes the formation of fast-coarsening grain-boundary Al2Cu phase, with limited Fe, Ni, and Co, that forms weak links and deteriorates mechanical properties.

[0045] The alloy can comprise impurities which are elements other than thoselisted above. The total amount of these impurities that are present is less than 0.5wt. % based on total weight of the alloy. The impurity can be Ag, and Ag should bepresent at less than 0.05 wt. % based on the total weight of the alloy. The totalamount of Ag can be 0 or equal to or less than 0.005, 0.01, 0.015, 0.02, 0.025, 0.03,0.035, 0.04, 0.045, or 0.05 wt. % based on the total weight of the alloy, or can be within a range of any high value and low value selected from these values. All otherimpurities should each be present at less than 0.1 wt. % based on the total weight ofthe alloy. The alloy can comprise in weight percent based on the total weight of the-Al2 -Al2Cu phase. The control of Cu, Mg,-Al2Cu phase than -Al2Cu phase in the grain interiors.

[0046] The alloy composition can have x= Fe + Ni + Co, and x = 0.5 to 3 wt. %based on the total weight of the alloy. The ratio of Cu / x can be from 5:1 to 8:1 in wt.% based on the total weight of the alloy. The ratio of Cu / x can be 5:1, 5.25:1, 5.5:1,5.75:1, 6:1, 6.25:1, 6.5:1, 6.75:1, 7:1, 7.25:1, 7.5:1, 7.75:1, or 8:1, and can be within a range of any high value and low value selected from these values.

[0047] -Al2Cu can be within a range of0.005-0.04 (0.5-4%) range, based on the total volume of the alloy. The volume-Al2Cu can be 0.5, 0.6, 0.7, 0.8, 0.9, 1.0, 1.1, 1.2, 1.3, 1.4, 12 168773228.11.5, 1.6, 1.7, 1.8, 1.9, 2, 2.1, 2.2, 2.3, 2.4, 2.5, 2.6, 2.7, 2.8, 2.9, 3, 3.1, 3.2, 3.3, 3.4, 3.5, 3.6, 3.7, 3.8, 3.9, or 4 volume %, based on the total volume of the alloy, or can be within a range of any high value and low value selected from these values.

[0048] The alloy can include -Al2Cu phase precipitates andintergranular Al7Cu2(Fe,Co,Ni), Al9(Fe,Co,Ni)2 and Al3(Fe,Co,Ni)2 precipitates. Thealloy linear fraction of grain boundary being covered by intergranularAl7Cu2(Fe,Co,Ni), Al9(Fe,Co,Ni)2 and Al3(Fe,Co,Ni)2 precipitates is at least 20%. Thelinear fraction of grain boundary being covered by these precipitates refers to a 2Dslice through the alloy, where the total length of grain boundaries in such a 2D sliceis determined, and the percentage of this length that is covered by the precipitates is the linear fraction of grain boundary coverage. The alloy linear fraction of grain boundary coverage can be 20, 21, 22, 23, 24, 25, 25, 27, 28, 29, 30, 31, 32, 33, 34,35, 36, 37, 38, 39, 40, 41, 42, 43, 44, 45, 46, 47, 48, 49, 50, 51, 52, 53, 54, 55, 56,57, 58, 59, 60, 61, 62, 63, 64, 65, 66, 67, 68, 69, 70, 71, 72, 73, 74, 75, 76, 77, 78, 79, 80, 81, 82, 83, 84, 85, 86, 87, 88, 89, 90, 91, 92, 93, 94, 95, 96, 97, 98, 99 or 100 %, and can be within a range any high value and low value selected from these values.

[0049] A method of making an alloy according to the present invention caninclude the step of combining 4.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; and balance Al. This combination of elements is melted. The melted combination is cast and solidified. The solidified composition is then subjected to a heat treatment. The heat treatment can include a solutionizing heat treatment, quenching in hot water, artificial aging and a preconditioning heat treatment.

[0050] The step of solution heat treatment can be heating the solidifiedcomposition to within 530-545 °C for a duration of at least 4 hrs. The step ofquenching can be immersing the composition in hot water within a range of 20-85 °C, where 60-95 °C range is preferred to avoid quenching-induced defect formation(i.e., part distortion). The step of aging can include subjecting the composition to atemperature of 230-240 °C for 5 hrs. The step of preconditioning can be subjectingthe composition to a temperature of 300 °C for at least 100 hrs. Other method stepsare possible. 13 168773228.1

[0051] Microstructural evolution and creep response were investigated in thecast Al-5.0Cu-0.3Mn-0.2Zr (wt.%) alloy with and without addition of slow-diffusing, intermetallic-forming elements Fe, Ni, or Co. Baseline Al-5.0Cu-0.3Mn-0.2Zr alloy exhibits high creep resistance at 300 °C, up to ~ 75 MPa, which is attributed to high-aspect-ratio, intragranular -Al2Cu precipitates that effectively suppress dislocationclimb. -Al2Cu precipitate-free zones form along grain boundaries uponheat treatment, whose extent is amplified during subsequent creep. Such weak regions experience dislocation creep, leading to strain localization and acceleration of grain-boundary sliding. Adding Ni and Co, individually or in combination, leads to the formation of grain-boundary precipitates (Al9Co2, Al3Ni2) which are resistant to -Al2Cu precipitate-free zones. This microstructure provides high creep resistance at stresses up to 75-80 MPa, withstrain rates much lower than in unmodified Al-5.0Cu-0.3Mn-0.2Zr. Adding Fe, whichresults in extensive decoration of grain boundaries with coarsening-resistant Al7Cu2Fe, and then increasing the Cu content to compensate for the Cu loss to thisnew phase, leads to a new Al-7.4Cu-1.6Fe-0.3Mn-0.2Zr alloy with creep resistanceat 300 ºC that surpasses presently known cast aluminum alloys. Adding Fe toimprove the creep resistance of Al-Cu alloys is both cost-effective and sustainable.These methods offer guidelines applicable to various alloy systems on controlling the-Al2Cu precipitate-free zones and its ensuing effects on creep deformation.

[0052] Coarsening-resistant grain boundary intermetallic particles enhance thecreep resistance of -stabilized Al-Cu-Mn-Zr alloys. Elemental candidates for forming new coarsening-resistant grain-boundary intermetallic phases must have the following characteristics: (i) slowly diffusing elements with low solubility in aluminum that do not interfere with the -stabilization mechanism and (ii) elements forming intermetallic phases with liquid Al in the last stage of solidification, therefore remaining at the grain boundaries. A series of Al-Cu-Mn-Zr alloys that formcoarsening-resistant grain-boundary precipitates were prepared. These alloyscontain individual and combined Ni and Co additions or individual Fe additions. The effect of these three alloying elements on the microstructure evolution duringsolidification and after subsequent heat treatments was assessed, and the impact ofthese microstructures on the compressive creep properties of Al-Cu-Mn-Zr alloy at 14 168773228.1300 °C was evaluated. A Co-and Ni-free Al-7.4Cu-1.57Fe-0.34Mn-0.15Zr (wt.%)alloy was prepared by increasing the Fe concentration in the Al-Cu-Mn-Zr alloy toform thermally stable Al7Cu2Fe (instead of -Al2Cu) to decorate grain boundaries extensively, and controlling / increasing the Cu concentration so that this Al7Cu2Fe formation does not compromise the fraction of intragranular -Al2Cu. This alloyexhibits excellent creep resistance surpassing that of existing cast aluminum alloysat 300oC, including Ni / Co-containing commercial RR350.

[0053] Cast Al-Cu-Mn-Zr alloys, generally referred to as ACMZ alloys, werefurther alloyed with three intermetallic-forming, slow-diffusing transition metals: Ni,Co, and Fe. The chemical composition of the alloys was measured using inductivelycoupled plasma optical emission spectroscopy (ICP-OES) at Dirats Laboratories(Westfield, MA, USA), and are listed in Table 1. Pure Al, Al-50Si, Al-50Fe, Al-50Cu,Al-60Mn, Al-10Zr, Al-10Co, and Al-20Ni master alloys, were melted at appropriate levels in an electric arc furnace under Ar atmosphere to obtain the six alloys with the compositions listed in Table 1. Other master alloys could be used. Rod-shaped billets (70 mm long and 20 mm in diameter) were obtained by pouring the liquid metal into a water-cooled copper mold. ACMZ-Co / Ni alloy listed in Table 1 is compositionally similar to the RR350 alloy, except for the additional small amounts of Sb and Ti in the latter alloy.

[0054] Table 1. Chemical composition of studied alloys (wt.%).Cu Mn Zr Si Fe Ni Co AlACMZ 5.0 0.29 0.16 0.07 0.08 - - Bal.ACMZ-Co 5.0 0.28 0.14 0.08 0.07 - 1.33 Bal.

[0055] The alloys were initially subjected to a T7 heat treatment, withsolutionizing conducted at 540 °C for 5 h, quenching in water at 80 °C, and artificial aging at 240 °C for 5 h. The alloys were then preconditioned at 300 °C for 200 h, a second aging treatment that stabilizes the microstructure prior to creep at 300 ºC. Heat treatments were performed in a muffle furnace with ± 3 °C temperature control accuracy. 15 168773228.1

[0056] Microstructural observations were carried out using a scanning electronmicroscope (SEM, FEI Quanta 650) equipped with energy-dispersive spectroscopy (EDS) and electron backscattered diffraction (EBSD) on the surfaces parallel to the longitudinal axis of the cast rods or creep-loading direction. Specimens were extracted from the center region of the rods, which were then mounted, ground, and polished down to ~ 0.06 μm, with the last polishing step performed using colloidalsilica. EBSD was performed at 20 kV, and the step size was set to 3 and 0.15 μmfor grain structure and phase contrast / orientation observations, respectively. While EBSD maps showing grain structures were post-processed using AztecCrystal software to remove non-indexed spots for better visualization, the EBSD maps showing phase contrast and orientation were not subjected to post-processing. A - Al2Cu precipitates in the base ACMZ alloy. TEM foils were prepared by electropolishing of disks with 3 mm diameter and ~50 μm thickness in an electrolyte consisting of 35% nitric acid plus 65% methanol, maintained within the negative 20- 30 ºC temperature range at 20 kV.

[0057] X-ray diffraction (XRD) measurements identified phases present invarious alloys. A Rigaku Smartlab Diffractometer equipped with a sealed tube Cu source and operating at 40 kV and 35 mA was used for measurements in Bragg- Brentano geometry. radiation. HyPix 3000 area detector operating in 1D mode collected the scattered intensity.

[0058] Creep tests were conducted in compression at 300 ±3 °C andmonitored using a K-type thermocouple attached to the specimen. The creep specimens, sized 10 mm in diameter and 20 mm in length, were machined from the cast rods. Each specimen produced several steady-state strain rates corresponding to different stresses. Testing was performed at constant load until a well-defined steady-state strain rate corresponding to the applied stress was reached. A series of stress jumps were made, each lasting until a steady-state strain rate was achieved, until the specimen accumulated ~10% total strain. The creep strain rate was measured using a linear variable differential transformer (LVDT), with ~ 10 μm accuracy, installed at the cold end of the creep extensometer. 16 168773228.1

[0059] The Vickers hardness of alloys after preconditioning treatment wasmeasured using a 5 kg load applied for 10 s, averaged over at least six indentations located in separate grains.

[0060] Compositional variations impact the grain structure - an essentialcharacteristic in determining alloy creep performance. Fig. 1(A) is an EBSD all-Eulerangle map showing the typical grain structure of base ACMZ alloy, with the grainstructures of the other alloys provided in FIGS. 12(B-F). Figure 12 shows EBSD all-Euler angle maps showing the grain structures of: FIG.12(A) ACMZ; FIG. 12(B)ACMZ-Fe; FIG.12(C) ACMZ-Cu / Fe; FIG.12(D) ACMZ-Ni; FIG.12(E) ACMZ-Co and FIG.12(F) ACMZ-Co / Ni.

[0061] Grains are equiaxed for all alloys. Fig. 1(B) is a plot of average graindiameter measured for all alloys. As displayed in Fig. 1(B), all alloys have grains withan equivalent size of ~50 μm, except for ACMZ-Co / Ni, which has coarser ~80 μm grains. EBSD scans were made on alloys after heat treatment. Black lines indicate high-angle grain boundaries with misorientations higher than 15°.

[0062] Figure 2 shows XRD spectra for the six ACMZ alloys (base andmodified) showing diffraction peaks and the associated phases. The peak at ~ 44 ° isdue to W (K-alpha) deposited on the Cu anode. -Al (FCC) and -Al2Cu (tetragonal I4 / mmm, a = 4.040 Å, c = 5.800 Å) exist in all alloys, whereas thepeaks related to -Al2Cu (tetragonal I4 / mcm, a = 6.067 Å, c = 4.877 Å) are found inACMZ, ACMZ-Co, ACMZ-Cu / Fe and ACMZ-Fe, but not in Ni-containing alloys, i.e.,ACMZ-Ni and ACMZ-Co / Ni. The Al7Cu2Fe phase (tetragonal P4 / mnc, a = 6.336 Å, c= 14.870 Å) is observed in ACMZ, ACMZ-Co, ACMZ-Fe, and ACMZ-Cu / Fe alloys, existing at considerably higher volume fraction in the latter three alloys relative toACMZ. The Al9Co2 (monoclinic P121 / a1, a = 8.565 Å, b = 6.290 Å, c = 6.213 Å)phase is present in alloys containing Co, i.e., ACMZ-Co and ACMZ-Co / Ni. XRD spectra from Ni-containing alloys (ACMZ-Ni and ACMZ-Co / Ni) reveal additionalpeaks exhibiting an excellent match with the Al3Ni2 phase (trigonal P-3 / m1, a = 4.053Å, c = 4.903 Å).

[0063] Figure 3 summarizes the typical microstructures of the alloys, and theresults of EDS measurements of various phases are provided in Table 2. Figure 3(A) shows a Bright-field TEM micrograph, acquired in two-beam mode near

[0001] Al 17 168773228.1-Al2Cu precipitates at grain interiors of base ACMZ alloy. Figures 3(B,C,E,F,G,H) show backscattered electron SEM micrographs showing grain-boundary microstructures for FIG.3(B) ACMZ; FIG.3(C) ACMZ-Co; FIG.3(E) ACMZ-Ni; FIG.3(F) ACMZ-Co / Ni; FIG.3(G) ACMZ-Fe; and FIG.3(H)ACMZ-Cu / Fe. Figure 3(D) is a phase-contrast EBSD map showing the distribution ofintermetallic phases indexed as Al7Cu2Fe and Al9Co2in the ACMZ-Co alloy.

[0064] The TEM micrograph in FIG. 3(A) shows -Al2Cuprecipitates in the grain interiors in the ACMZ alloy, which is representative of allinvestigated alloys, except for some variation in the precipitate fraction among -Al2Cu precipitates up to 350 °C. However, as discussed next, the grain-boundary microstructure varies between alloys.

[0065] In ACMZ, inter-dendritic / granular Al7Cu2Fe and -Al2Cu precipitatesare observed at a low (~0.9 vol.%) volume fraction (FIG. 3(B) and Table 3). Fine -Al2Cu precipitates are also visible in the SEM micrograph (FIG.3(B)), which are --PFZ) developedalong grain boundaries (inset in FIG.3(B)).

[0066] In ACMZ-Co, the grain boundaries display mostly Al9Co2 precipitates,which form on solidification with various morphologies, some being coarse and platelet-shaped, while others being relatively finer and irregularly shaped (FIG.3(C)). Such morphological variation results from their precipitation during various stages of solidification. In addition to Al9Co2, a minor grain-boundary phase is identified as Al7Cu2(Fe,Co), based on the XRD and EDS measurements (FIG.2 and Table 2,respectively), which is also formed on solidification (see FIG. 3(C)). This phase isisostructural to Al7Cu2Fe observed in ACMZ, yet it forms at a much higher fraction in ACMZ-Co than in ACMZ because a significant amount of Co also incorporates into this phase (Table 2). A higher fraction of insoluble Al7Cu2(Fe,Co) scavenges moreCu, which is -Al matrix of ACMZ-Co has lesser Cu compared to ACMZ(1.81 vs 2.21 at.%, see Table 2). Both ACMZ and ACMZ-Co have similar hardness, -Al matrix after solutionizing has no significant influence on the fraction of the -Al2Cu formed onsubsequent aging, as this strengthening phase largely determines alloy hardness.ACMZ-Co has much higher grain-boundary intermetallic phase fraction compared to 18 168773228.1ACMZ (~6.5 vs ~0.9%, see Table 3), thus enabling more extensive decoration ofinter- -PFZ formation during heattreatment (FIG.3(C)). A phase-contrast map, acquired by EBSD measurements, shows that Al7Cu2Fe precipitates reside mainly in the vicinity of Al9Co2phase. In contrast, the latter phase also appears at grain interiors and along grain boundaries.

[0067] Figure 13(A) is a phase-contrast map and FIG. 13(B) is an inverse polefigure EBSD map, acquired from ACMZ-Co alloy after preconditioning heat treatment. Figure 13(A) shows the distribution of Al9Co2 and Al7Cu2Fe precipitates in the microstructure. Figure 13(B) show the orientation variation for Al9Co2and Al7 -Al grains.

[0068] The IPF map in FIG. 13(B) indicates that numerous Al9Co2 andAl7Cu2Fe precipitates residing within a single grain exhibit a non-fixed orientation -Al, and these intermetallic phases are expected to exhibit -Al. Figures 14 (A-D) show a comparison of the typicalhigh-magnification microstructures of FIG.14(A) baseline ACMZ; FIG.14(B) ACMZ-Ni; FIG. 14(C) ACMZ-Fe; and FIG. 14(D) ACMZ- -Almatrix, measured by SEM-EDS, is provided for each alloy for easy comparison. Theprecipitates residing at the grain boundaries, which have interfaces with two differently oriented grains (FIGS.14 (A-D)), are also expected to remain incoherent, at least along most of their interfaces.

[0069] ACMZ-Ni exhibits grain-boundary precipitates that are predominantlycomposed of the Cu / Ni-rich intermetallic phase (~4.8 vol%, Table 3), which, similar to those in ACMZ-Fe and ACMZ-Co alloys, extensively decorate inter- -PFZ evolution (FIG.3(E)). XRD measurement reveals the presence of the Al3Ni2 phase in this alloy (FIG.2), which is identified as Al3(Ni,Cu)2, also based on the EDS measurement of this phase (Table 2). -Al grain interiors of ACMZ-Ni contain ~1.5 at.% Cu (Table 2 andFIG. 14(B)), which is between values obtained for ACMZ and ACMZ-Fe; thisvariation is consistent with the hardness evolution tendency among these alloys (Table 4). Turning to ACMZ-Co / Ni, its grain-boundary microstructure consists ofAl9(Co,Ni)2 and Al3(Cu,Ni)2 phases (FIG. 3(F) and Table 2), with a total volumefraction of ~6.2% (Table 3). As both phases consume a significant amount of Cu -Al matrix is ~ 2 at.%, close to those of ACMZ and 19 168773228.1ACMZ-Co, which is why the hardness of these three alloys is the same, withinmeasurement error (Table 4).

[0070] The microstructure of ACMZ-Fe (FIG. 3(G)) is characterized by arelatively high volume fraction of non- intermetallic phase (~ 7%, Table 3), which isidentified as Al7Cu2Fe based on the XRD and EDS measurements (FIG.2 and Table 2, respectively). These interconnected and irregular-shaped intermetallic precipitates are distributed both along grain boundaries and interdendritic regions, thus almost suppressing all -PFZ formation (FIG.3(D)). However, a high volumefraction of insoluble Cu-rich intermetallic precipitates reduces Cu supersaturation in-Al matrix during solutionizing. Hence, less Cu is available for precipitationduring aging, as evidenced by much lower EDS- -Al (~1at.%, Table 2), and much lower hardness obtained for this alloy as compared toACMZ (63.3 HV vs. 93.3 HV, see Table 4), which is also consistent with lower-Al2Cu in ACMZ-Fe relative to ACMZ (FIG. 3(G) vs. FIG. 3(B) and FIG.14(C) vs. FIG. 14(A)).

[0071] To obtain a high fraction of Al7Cu2Fe -Al2Cufraction, the Cu content of the alloy was increased such that excess Cu (beyond the amount needed for the Al7Cu2 -Al at solutionizing temperature. The resulting ACMZ-Cu / Fe alloy shows a similar microstructure as ACMZ-Fe regarding the types of phases and their morphologies. However, the fraction of -Al2Cu precipitates is much higher in the former than in thelatter (FIG.3(H) vs. FIG. 3(G) and FIG. 14(D) vs. FIG. 14(C)). The higher Cu-matrix and the corresponding higher hardness observed in this alloy -Al2Cu.

[0072] The volume fraction of Al7Cu2Fe is higher in ACMZ-Cu / Fe than inACMZ-Fe (7.6 vs.6.9 vol.%, Table 3) due to the higher Fe content of the former alloy (Table 1).

[0073] Table 2 shows EDS measurements of the constituent phases (exceptfor -Al2Cu (I4 / mmm)) observed in various alloys' microstructures (FIG.3) after preconditioning heat treatment. Results are in at%. Phases were identified basedon the XRD measurements reported in FIG. 2, and the EDS data provides aqualitative measure of the solubility of various elements in the constituent phases. 20 168773228.1Table 2. Alloys Phase Cu Mn Zr Si F Co Ni AlACMZ -Al (FCC) 2.2 0.1 0.0 0.0 0. - - BaAl7Cu2Fe (P4 / mnc) 19. 1.7 - 0.2 4. - - Baintermetallic phases quantified for various alloys after preconditioning heat treatment. Table 3 Alloys ACMZ ACMZ-Co ACMZ-Fe ACMZ-Cu / Fe ACMZ-Ni ACMZ-Co / Ni

[0075] Table 4 shows the Vickers hardness (HV5) of various alloys afterpreconditioning heat treatment. Table 4 Alloys ACMZ ACMZ- ACMZ- ACMZ- ACMZ- ACMZ-93.3 ± 3.6 92.6 ± 63.3 ± 85.0 ± 1.5 79.3 ± 88.0 ±

[0076] Figure 4 shows double-logarithmic plots of steady-state strain rateversus compressive stress at 300 °C for the six ACMZ alloys (base and modified). The volume fraction (ƒ) of non- comparison. Figure 4 displays the evolution of minimum creep strain rates as a21 168773228.1function of applied compressive stress at 300 °C, plotted in accordance withNorton’s power-law for creep:=(1)where B is a material constant containing temperature and activation energy, and is the apparent stress exponent. The ACMZ alloy, for the studied range of applied stresses, exhibits two distinct creep regimes, with =2.8 at < 90 MPa and =20-30 at higher stresses. The creep curves of ACMZ-Ni, ACMZ-Co, andACMZ-Co / Ni point to three different creep regimes: (i) ~ 1 at low stresses (< 40MPa), (ii) ~ 3-4 at intermediate stresses (< 70-80 MPa) and (iii) ~ 20-30 athigh stresses (> 70-80 MPa) (FIG.4). These alloys display roughly an order of magnitude lower creep strain rates relative to ACMZ at low and intermediate stresses, whereas at high stresses, their creep strain rates are largely comparable (FIG.4). Among Ni- and / or Co-containing alloys, ACMZ-Co / Ni alloy exhibits a higher creep resistance than ACMZ-Ni or ACMZ-Co and this difference is more pronounced in the ~ 1 regime (FIG.4).

[0077] The creep response of ACMZ-Fe differs significantly from all otheralloys. While ACMZ-Fe exhibits comparable creep strain rates as ACMZ-Ni or ACMZ-Co up to 30 MPa, at higher stresses, its creep resistance deteriorates drastically, even becoming much inferior to ACMZ at > 40 MPa. The ~ 1 regime observed at < 35 MPa changes directly to ~ 20-30 at > 35 MPa, without a ~ 3-4 regime at intermediate stresses. ACMZ-Cu / Fe, by contrast, displays comparable creep response as ACMZ-Ni or ACMZ-Co, with three clearly-defined distinct creep regimes: ~ 1, ~ 3-4 and ~ 20-30. Finally, similar to Co- and / or Ni-containingCu / Fe exhibits almost an order of magnitude lower strain rates than ACMZ at low and intermediate stresses. ACMZ-Cu / Fe alloy displays a similar creep response to all other alloys (except ACMZ-Fe) at high stresses.

[0078] Figures 5(A-F) show SEM backscattered electron micrographs showingtypical post-creep microstructures of five ACMZ alloys (base and modified): FIGS. 5(A,B) ACMZ; FIG.5(C) ACMZ-Fe; FIG.5(D) ACMZ-Cu / Fe, FIG.5(E) ACMZ-Co and FIG.5(F) ACMZ-Co / Ni. Loading direction is vertical for all micrographs. Dashed 22 168773228.1arrows point to the cracked grain-boundary intermetallic precipitates. Figure 5shows typical post-creep microstructures of these alloys. A common observation relevant for all alloys is the presence of the -Al2Cu precipitates in their grain interiors. This observation is consistent with the stabilizing effects of Mn and Zr present in all investigated alloys (Table 1). However, ACMZ-Fe has fewer -Al2Cu than all other alloys (FIG.5(C)). Grain-boundary microstructure, which differs significantly between alloys prior to creep (FIG.3), evolves at distinct rates for each alloy during creep.

[0079] In ACMZ alloy, inter- -Al2Cu precipitates are coarserthan those prior to creep (FIGS.5(A,B) vs. FIG.3(B)), indicating a marked coarsening during creep. -Al2Cu coarsening occurs at the expense of nearby -Al2Cu precipitates, resulting in much wider -PFZ (FIGS.6(A,B)). -Al2Cu precipitates at triple junctions appear coarser than those at grain boundaries, while those at interdendritic regions underwent less coarsening during creep (FIG.5(A)). The total creep time for this specimen was ~ 386 h, almost twice that of the priorpreconditioning heat -PFZ-PFZ andthe concomitant grain- -PFZevolution during creep.

[0080] Figures 6(A-F) show EBSD maps acquired from base ACMZ alloybefore (FIGS.6(A,B,C)) and after (FIGS.6(D,E,F)) creep deformation (with ~ 10% total strain accumulated within the applied stress range of 20-133 MPa during the stress-jump creep test): (FIGS.6(A,D)) inverse pole figure (IPF) maps showingdistribution of low- -(FIGS.6(B,E)) kernel average misorientation (KAM) maps, calculated using the circular-shaped kernel of 5x5 matrix with a threshold angle of 5°, showing local misorientation distribution, calculated using the center point of the kernel as a reference orientation, and (FIGS.6(C,F)) grain reference orientation deviation (GROD) maps showing misorientation distribution within the grains, calculated using the averaged grain orientation as a reference orientation. High GROD values obtained for grains indicated by arrows (FIGS.6(C,F)) are not reflective of high strain accumulation in these grains, but rather are due to the presence of a grain 23 168773228.1comprising several dendritic arms with apparent misorientation between each other. The loading direction is vertical for the crept condition.

[0081] The inverse pole figure (IPF) maps clearly show that creep deformationcauses a significant increase in the quantity of low-angle grain boundaries (FIG. 6(A,D)). In contrast, kernel average misorientation (KAM) maps (FIG.6B,E)) demonstrate geometrically necessary dislocations (GNDs)-induced misorientationconcentrations near the grain boundaries, that is -PFZ, suggesting strainlocalization in these regions, which is consistent with our previous reports. Here, we also demonstrate, by creating the grain reference orientation deviation (GROD)maps (FIG. 6(C,F) -PFZ along grain boundaries accommodate much higherstrain compared to grain interiors, as evidenced by significantly increased GROD -PFZ along grain boundaries after creep deformation relative to thoseprior to creep (FIG. 6(F) vs. FIG. 6(C)). KAM maps mirror the actual dislocationstructure within the local region, that is, within the kernel, and so, they are sensitiveto the effective plastic strain. The GROD maps, in contrast, are better suited to correlate with the degree of total plastic strain and creep damage as GROD maps display intragranular misorientation distribution calculated using the averaged grain orientation as its reference orientation. Therefore, a significant increase in GROD -PFZ along grain boundaries after creep deformation (FIG. 6(F)) indicateshigher creep strain accumulation in these regions than grain interiors.

[0082] Turning to Fe-modified alloys, both ACMZ-Fe and ACMZ-Cu / Fe exhibitalmost no -PFZ along grain boundaries decorated with Al7Cu2Fe (FIG. 5(C) andFIG.5(D), respectively), with the only difference being a higher number density of - Al2Cu in the latter alloy than in the former. In ACMZ-Co, -PFZ can be seen at some grain boundaries (FIG.5(E)), but most grain boundaries, which are decorated with Al9Co2 and Al7Cu2 -PFZ. Although this alloy has a high volume fraction of intermetallic phase (6.5%), not all grain boundaries are well decorated with intermetallic phases because some Al9Co2 likely formed in the earlier stages of solidification; intermetallic-free grain boundaries seem -PFZ evolution during creep. ACMZ-Co / Ni exhibits well-decoratedgrain boundaries with intermetallic precipitates and without any -PFZ (FIG. 6(F)).The same observation applies to ACMZ-Ni (FIG. 15). Figure 15 is a SEMbackscattered electron micrographs showing the typical post-creep microstructure of24 168773228.1ACMZ- -PFZ isobserved at / along the grain boundaries of this alloy.

[0083] Cracked intermetallic precipitates are also frequently observed in post-creep microstructures, as indicated by the dashed arrows in FIG.5.

[0084] Prior research has addressed the creep behavior of cast Al-Cu-Mn-Zr(ACMZ) alloys differing from the base ACMZ (with ~5 wt.% Cu) alloy studied here by (i) the additional presence of 0.1 wt.% Ti and (ii) higher Cu concentration, between~6 and ~9 wt.%. Figure 7 compares the creep response of all these ACMZ alloys-Al2Cuprecipitates but varying fractions of grain- -Al2Cu precipitates dependingon the Cu content of the alloys. The creep resistance of our base ACMZ alloy is comparable to other ACMZ alloys (FIG.7), confirming that their creep resistance isprimarily controlled by -precipitates within grains, with no significantstrengthening from grain-boundary -Al2Cu. Creep mechanisms that operate in these alloys at various stresses are discussed in the following. Figure 7 shows double-logarithmic plots of steady-state strain rate versus compressive stress at 300 °C for the three unmodified ACMZ alloys with various Cu concentrations, andtherefore, with different fraction ƒ of grain- -Al2Cu precipitates; also shownare the data for ACMZ-Cu / Fe, RR350, and Al-1Mn-0.3Zr-0.3Er-0.05Si (wt.%).Volume fraction (ƒ) of grain-boundary intermetallic precipitates in each alloy is provided for comparison. Creep data for RR350 and ACMZ (5.7% Cu) are reproduced from (J.U. Rakhmonov, S. Bahl, A. Shyam, D.C. Dunand, Cavitation- -strengthened Al-Cu based alloys, Acta Mater.228 (2022) 117788.13), and those for ACMZ (6% Cu) and ACMZ (9% Cu) are obtained from (S. Bahl, J.U. Rakhmonov, C. Kenel, D.C.Dunand, A. Shyam, Effect of grain- -Al2Cu precipitates on tensile andcompressive creep properties of cast Al–Cu–Mn–Zr alloys, Materials Science and Engineering: A 840 (2022) 142946.11). Creep data for cast Al-1Mn-0.3Zr-0.3Er-0.05Si strengthened by intragranular L12- -AlMnSiprecipitates are reproduced from Ref. (A.R. Farkoosh, D.C. Dunand, D.N. Seidman, Solute-induced strengthening during creep of an aged-hardened Al-Mn-Zr alloy, Acta Mater.219 (2021) 117268.22). The dashed and dotted lines (which are parallel to 25 168773228.1the n=1 slope) indicate strain rates estimated using Coble and Nabarro-Herring (N- H) models for a grain size of 55 μm.

[0085] The diffusional creep rates for both Coble and Nabarro-Herring creepfor a grain size of 55 μm are estimated using equations and materials parameters forpure Al by Frost and Ashby (H.J. Frost, M.F. Ashby, Deformation-mechanism maps:the plasticity and creep of metals and ceramics, Pergamon Press, New York, 1982), as plotted in FIG.7. At this grain size, both models predict similar creep rates, with Coble model producing slightly higher creep rates as compared to Nabarro-Herring creep (FIG.7). For ACMZ alloys, the measured and estimated strain rates up to 20- 25 MPa applied stress, i.e., in the ~ 1 regime, are comparable (FIG.1). Milliganet al. (B.K. Milligan, S. Roy, C.S. Hawkins, L.F. Allard, A. Shyam, Impact ofmicrostructural stability on the creep behavior of cast Al–Cu alloys, MaterialsScience and Engineering: A 772 (2020) 138697) found that the grain size exponentthat is sensitive to the diffusional creep mechanism is close to that of Coble creep for the ACMZ (with ~ 7% Cu) alloy. Coble creep as the dominant diffusional creep mechanism is reasonable given the limited strengthening effect of grain-boundary - Al2 -PFZ evolution along grain boundaries that accelerates vacancy flow, particularly, with the increased level of strain localization. For all other alloys, except for ACMZ-Fe, the diffusional creep dominates up to ~ 40 MPa stress (FIG.4), indicating the delayed onset of ~ 3 regime. Also, the creep rates for all these alloys are much lower than ACMZ alloy. For instance, a thousandfold lower strain rate is obtained for all other alloys than for ACMZ alloy at 20 MPa, which reflects the beneficial effects of the coarsening-resistant intermetallic phase extensively decorating grain boundaries in terms of (i) reducing the vacancyflow along grain boundaries due to near- -PFZ and (ii)restricting GBS needed to accommodate the diffusional creep.

[0086] The ~ 3 regime (with =2.8 for ACMZ and ~ 3-4 for otheralloys, except for ACMZ-Fe) reflects creep deformation occurring over a wide rangeof stresses, particularly for the base ACMZ. In previous work with ACMZ (with 7%Cu) and RR350 alloys (J.U. Rakhmonov, B. Milligan, S. Bahl, D. Ma, A. Shyam, D.C.Dunand, Progression of creep deformation from grain boundaries to grain interior inAl-Cu-Mn-Zr alloys, Acta Mater. 250 (2023) 118886), which involved in situ neutrondiffraction measurements during creep and ex situ TEM observations of 26 168773228.1microstructures before and after creep, it was demonstrated that dislocation creep inthe grain interiors remains limited at stresses corresponding to ~ 3. Given much higher strain rates measured for these alloys, including ACMZ alloy, than the predicted diffusional creep rate (FIG.7) and the limited intragranular dislocation creep in the ~ 3 regime, a GBS mechanism is likely the rate-limiting mechanism in the ~ 3 regime. Since GBS also occurs in the ~ 1 regime to accommodate diffusional creep, as noted earlier, the question arises as to how GBS can significantly increase the strain rate in the ~ 3 regime relative to theestimated diffusional creep rate. It is believed that the stresses in the ~ 3regime activate Rachinger sliding, which requires intragranular dislocation movement as an accommodation mechanism for GBS, as opposed to Lifshitz sliding needed to accommodate Coble or Nabarro-Herring creep which requires vacancy movement. -PFZ where strain localizes (FIG.6) and the localized dislocation -PFZ occurring at even low stresses corresponding to the ~ 3 regime can create favorable conditions for Rachinger sliding to dominate in the ~ 3 regime. Rachinger sliding strongly depends on grain size (Q. Zhang, Y. Zhu, X. Gao, Y. Wu, C. Hutchinson, Training high-strength aluminum alloys to withstand fatigue, Nature Communications 11(1) (2020) 5198), where alloys with finer grains exhibit higher strain rates than those with coarser ones at a given stress.

[0087] Figure 8 (A,B) shows SEM backscattered electron micrographsshowing the typical post-creep microstructures of base ACMZ, with FIG.8(B) showing a higher-magnification view of an area within the dotted rectangle in FIG. 8(A). Figure 8(C) is a schematic diagram illustrating the main microstructural factors controlling the creep behavior of ACMZ alloy and the presence of increased GND -PFZ, as determined by EBSD mapping provided in FIGS.6(d- -PFZ distribution is non-uniform, with some g -PFZ,-PFZ. The arrow in FIG.8 (A,B) indicates the direction in which GBS likely occurs relatively easily, resulting in the formation of the denuded zones along the boundaries of this grain, thus contributing -PFZ. The dotted circle in FIG.8(B) indicates the region of the grain that likely experiences increased stress concentrations near a triple junction. Given the irregular morphology of the grains, some have physical constraints for GBS to proceed as indicated for one such grain in FIG.8(A). The 27 168773228.1schematic in FIG.8(C) illustrates that, although dislocation movement in the graincores is inhibited in the n_app=1- -PFZ enablesRachinger-type GBS to occur, consistent with the core-mantle model by Gifkins (R.C. Gifkins, Grain-boundary sliding and its accommodation during creep and superplasticity, Metall. Trans. A 7(8) (1976) 1225-1232).

[0088] As shown in FIGS. 8(A,B) -PFZ evolves extensively along certaingrain boundaries, particularly those that can easily experience GBS and the regions that experience high local stresses, that is, triple junctions. Given the irregular morphology of the grains, not all grains exhibit favorable boundaries for sliding to occur, with some boundaries remaining physically constrained to slide, as indicatedin FIG. 8(A) -PFZ. The schematic in FIG.8(C)c illustrates, for the ACMZ alloy, the post-creep microstructure and the dislocation structure (created based on the experimental observations of the microstructures (FIGS.8(A,B)) and the distribution of GNDs (FIGS.6(E,F)). Although dislocation movement is expected to be inhibited in the grain interiors in the -PFZ enables Rachinger sliding to proceed, consistent with the core-mantle model by Gifkins, according to which GBS and its accommodation are only limited to dislocation movements within the peripheral “mantles”.

[0089] The ~ 20-30 regime has been convincingly attributed to-precipitates within grains, in previous reports on ACMZ (with 7% Cu) and RR350 alloys. The Orowan stress is predominantly dictated by the spacing (which is a function of a precipitate size and -Al2Cu precipitates. M -Al2Cu precipitates in ACMZ-Fe is responsible for a significant shift in the ~ 20- 30 regime to lower stresses (or higher strain rates). Given the high aspect ratio of - Al2Cu precipitates and their resistance to shearing, dislocation climb remains inactive, and dislocation-controlled creep in the grain cores / interiors only activates upon reaching the Orowan stress of the alloy. Although Orowan looping becomesthe rate-limiting mechanism, other creep mechanisms, such as diffusional creep andGBS also remain active.28 168773228.1

[0090] The formation of PFZs is common for many precipitation-hardenedaluminum alloys, with narrow PFZs in some systems forming upon aging. Although --PFZbecomes extensive during creep (FIGS. 5(A,B) and FIGS. 8(A,B)). The optimization-PFZ width in the alloy cannot suppress its further evolution during creep if grain boundaries are not effectively strengthened by coarsening-resistant intermetallics, as discussed further in the next section. Grain- -Al2Cu precipitates in ACMZ experience growth and coarsening, -Al2Cu precipitates resulting in an -PFZ (FIGS. 5(A,B) and FIGS. 8(A,B)). -Al2Cu coarsening and theresultant growth of -PFZ is much more pronounced at the triple junctions (FIGS.5(A,B) and FIGS. 8(A,B)), due likely to the enhanced local diffusion and the GBS-induced high local stress concentrations near triple junctions. The GBS resulting in the stress concentrations near the triple junctions and the concomitant widening of -PFZ are schematically illustrated in FIG. 8(C). Not only high local stresses andaccelerated diffusion through short-circuit paths, that is, grain boundaries anddislocation cores that accelerate -to-with respect to their neighbors can result in the formation of denuded zones, which is -PFZ volume during creep.

[0091] -PFZ volume increases the creep-PFZ even in the ~ 1 and ~3 regimes. Given the creep-induced large volume of widened and networkedlocalized dislocation creep likely occurs continuously during creep, contributing to the overall creep strain rate in ACMZ. GROD angle distribution maps before and after creep (FIG.6) prove the accumulation of much higher plastic strain -PFZ than in the grain interiors of base ACMZ.

[0092] The Ni- and / or Co-modified ACMZ alloys, labeled ACMZ-Ni, ACMZ-Co,and ACMZ-Ni / Co, exhibit comparable creep resistance to RR350 (FIGS.4 and 7), despite RR350 being compositionally more complex than our alloys, as it also contains Sb and Ti (which are expected to improve alloy castability and refine grains,respectively). Grain-boundary microstructures of ACMZ-Ni, ACMZ-Co, and ACMZ-Ni / Co differ from each other by the type and morphology of grain-boundary precipitates (FIG.3 and Table 2), but a common observation for these alloys is the 29 168773228.1extensive decoration of grain boundaries by these precipitates and the effective -PFZ formation observed in unmodified ACMZ. Although both individual and combined Co and Ni additions result in extensive grain- boundary decoration, the microstructure of ACMZ-Co also reveals some coarse Al9Co2 platelets residing in grain interiors and boundaries (FIGS.3(C,D)); this is attributed to the broad temperature range for the formation of Co-rich precipitates,some of which form as pr -Al during solidification, asexpected from the Calphad-based simulation of the alloy displayed in FIG.9(A). Therefore, selection of the alloying element / s to strengthen grain boundaries should also consider the temperature range for precipitate phase formation to maximize their effectiveness in terms of decorating grain boundaries while also exhibiting more desirable morphologies.

[0093] ACMZ-Co / Ni exhibits a higher creep resistance than RR350 at low andintermediate stresses despite their comparable fraction of grain-boundary phases (~ 6 vol%). However, ACMZ-Ni / Co has a coarser grain size compared to RR350 (~ 75 vs.57 μm), which explains why the former alloy exhibits lower creep strain rates in the ~ 1 and ~ 3 regimes where grain-boundary microstructure plays a crucial role in the alloy creep behavior. Also, different casting procedures employed for RR350 and for our alloys produced different cooling rates during solidification, as noted from much finer secondary dendrite arm spacing (SDAS) measured for ACMZ-Ni / Co compared to RR350 (~ 19 vs. 35 μm, FIG. 16). Figure 16 provides acomparison of the typical microstructures of ACMZ-Co / Ni and RR350 alloys after preconditioning heat treatment. Intermetallic precipitates are much finer in ACMZ- Co / Ni than in RR350. Finer secondary dendrite arm spacing (SDAS) is measured for ACMZ-Co / Ni as compared to RR350 alloy (~ 19 vs.35 μm), indicating a higher cooling rate during solidification of the former alloy; lower cooling rate duringsolidification of RR350 is due to a different casting method used for this alloy. Earlierin-situ neutron diffraction measurements during creep of RR350 and ACMZ,revealed that Ni / Co-rich precipitates are more effective in sharing load with -Almatrix compared to -Al2 - -Al2Cu accelerates plastic relaxation mechanisms, reducing the effectiveness of load -Al2Cu at grain boundaries. The load-transfer effectiveness of Ni / Co-rich second phases at grain boundaries is expected to increase further with their 30 168773228.1refinement, because a higher stress for activating plastic relaxation is needed for finer particles, as compared to coarser ones.

[0094] The Fe-modified ACMZ alloys, or ACMZ-Fe, display high creepresistance below 30 MPa, with almost an order of magnitude lower creep strain rates when compared to ACMZ; however, above 40 MPa, this alloy becomes much lesscreep resistant than ACMZ (FIG. 4). The apparent stress exponent above 40 MPa ishigh ( = 20-30) and almost the same as those obtained for the other alloysabove 70-75 MPa, where Orowan looping is the rate-limiting mechanism for dislocation motion in the grain interiors. The poor creep resistance above 40 MPa is -Al2Cu precipitates, as -Al2Cu and increased interprecipitate spacing in ACMZ-Fe compared to ACMZ (FIG.5(C) vs. FIG.5(B)). Lower hardness -Al matrix (Table 2) in ACMZ-Fe compared to -Al2Cu precipitates obtained in ACMZ-Fe compared to other alloys. The high fraction of Al7Cu2Fe the grains during subsequent aging treatment. By contrast, ACMZ-Cu / Fe, for which -Al2Cu and Al7Cu2Fe, exhibits improved creep resistance up to 70-75 MPa. This alloy outperforms RR350 at low and intermediate stresses (FIG.7) due to mostly higher fraction of grain-boundaryprecipitates (~7.5 vs. 6.1 vol%) and their refined size (FIG. 3(H) vs. FIG. 16(B)),which further restrict GBS, and possibly provide higher load-transfer strengthening effect.

[0095] Figure 9(A,B) are plots of equilibrium fraction of phases vs.temperature, as calculated using Thermo-Calc (TCal 8 database) for FIG.9 (a) ACMZ-Co alloy and FIG.9 (b) ACMZ-Cu / Fe compositions. For ACMZ-Cu / Fe composition, Fe- and Mn-rich phases (e.g., Al13Fe4 and Al6Mn) other than Al7Cu2Fe were excluded from the simulation due to their absence in its microstructure. Unlike Al9Co2 which forms as a primary phase with platelet morphology during the earlystages of solidification in ACMZ-Co (FIGS. 3(C,D) and FIG. 9(A)), the Al7Cu2Fephase in ACMZ-Cu / Fe forms at the later stages of solidification (FIG.9(B)), leading to an extensive decoration of grain boundaries. The high coarsening resistance of Al7Cu2Fe is provided by the slow diffusivity and low solid solubility of Fe in Al, which 31 168773228.1-PFZ formation, thereby restricting both diffusional creep and Rachinger sliding and rendering the alloy highly creep resistant in the ~ 1and ~ 3 regimes.

[0096] Designing aluminum alloys with improved creep resistance at 300 ºCopens the door to widespread application in various sectors. Existing Al-Si(-Mg)(-Cu) alloys used widely in the automotive industry are not suitable for use above 250 °C due to their intrinsically inferior high-temperature strength. Also, the strategy of -Al2Cu up to 350 °C via interfacial segregation ofslow diffusers, achieved in Al-Cu-based alloys, cannot be easily introduced into Al-Sialloys because hig -Al2Cu and-Al2Cu upon exposure above 250 °C. Aluminum alloys strengthened with coarsening-resistant L12-Al3 - Al(FeMn)Si display moderate creep resistance in the 300-400 °C range. As shown in FIG.7, a recently introduced L12 -strengthened Al-1Mn-0.3Zr-0.3Er-0.05Si (wt.%), exhibits creep resistance at 300 ºC much lower than those of even base ACMZ alloy. Low solid solubilities of transition elements, such as Zr, Fe, and Mn, do not allow for a significant increase in the precipitate fraction formed on aging. Moreover, due to the equiaxed shape of precipitates, dislocation motion in the graininteriors occurs at stresses well below the yield strength of the alloy via thethermally-activated dislocation climb mechanism.

[0097] -strengthened Al-Cu alloys exhibit unique microstructures,providing pathways to improve their creep resistance significantly. Figure 10 schematically illustrates the main microstructural constituents in ACMZ - with and without Ni, Co, or Fe addition, and the creep damage evolution in such alloys with its concomitant effect on alloy creep resistance. platelet morphology with high aspect ratio, can effectively restrict dislocation climb in Al-Cu alloy, delaying extensive dislocation creep in the grain interiors / cores until the Orowan stress of the alloy is reached, as these precipitates also resist shearing. In the presence of grain-boundary precipitates with low coarsening resistance, such as -Al2Cu, these precipitates grow and coarsen during heat treatments and -PFZ, where localized dislocation creep occurs at stresses well below the Orowan stress within the grains, thereby increasing the overall strain rate of the alloy. Decorating the 32 168773228.1grain boundaries with coarsening-resistant precipitates, such as Al7Cu2Fe, Al3(Cu,Ni)2, and Al9(Co,Fe)2, as done in the present alloys, effectively suppresses -PFZ and slows the GBS rate, making grain boundaries morecreep-resistant. The Al-7.3Cu-1Fe-0.4Mn-0.2Zr (ACMZ-Cu / Fe) alloy exhibitssuperior creep resistance to RR350 and offers potential for its use at high temperatures given that such high creep resistance is achievable only through the control of Fe and Cu contents of ACMZ, making the alloy (i) cost-effective due to the absence of costly Ni and Co alloying additions and (ii) compatible with thesustainable alloy design efforts since both Fe and Cu – common impurities inrecycled aluminum – are turned from undesirable impurities into beneficial alloyingadditions.

[0098] Figure 10 shows schematic illustration of the primary microstructuralconstituents before and after the compressive creep in baseline and Fe / Ni / Co- modified ACMZ alloys. Figure 11 shows the differences in their creep performance along with the proposed creep mechanisms and the underlying main microstructuralfactors. The light solid lines point to grain boundaries, and the dashed lines indicate-PFZ. The Coble creep and its dependence on grain size, thehigh- -Al2Cu precipitates effectively suppressing dislocation climb, andthe dislocations by -Al2Cu precipitates by Orowan looping in ACMZ havebeen established (B. Milligan, D. Ma, L. Allard, A. Clarke, A. Shyam, Dislocation- (Al2Cu) interactions during creep deformation of an Al-Cu alloy, Scripta Mater.217 (2022) 114739; B.K. Milligan, S. Roy, C.S. Hawkins, L.F. Allard, A. Shyam, Impact of microstructural stability on the creep behavior of cast Al–Cu alloys, Materials Science and Engineering: A 772 (2020) 138697; J.U. Rakhmonov, S. Bahl, A. Shyam, D.C. Dunand, Cavitation-resistant intergranular precipitates enhance creep -strengthened Al-Cu based alloys, Acta Mater.228 (2022) 117788; J.U. Rakhmonov, B. Milligan, S. Bahl, D. Ma, A. Shyam, D.C. Dunand, Progression of creep deformation from grain boundaries to grain interior in Al-Cu-Mn-Zr alloys, Acta Mater.250 (2023) 118886). The napp ~ 3 regime where the dislocation creep in -strengthened grain interiors remains inactive

[0014] and the higher strain rate was measured experimentally than that estimated theoretically when considering thegrain-boundary or lattice diffusion (FIG. 7), the Rachinger sliding is proposed to also-PFZ stemming from the localized 33 168773228.1dislocation creep likely activate the Rachinger-type sliding, consistent with the core- mantle model by Gifkins.

[0099] Microstructural evolution during heat treatment and creep deformationin Al-5.0Cu-0.3Mn-0.2Zr (wt.%) alloy, with and without addition of slow-diffusing, intermetallic-forming Ni and / or Co or Fe, and its effect on alloy creep resistance at300 °C were investigated. Base Al-Cu-Mn-Zr alloy with intragranular -Al2Cu andfew grain-boundary -Al2Cu / Al7Cu2Fe precipitates showed high creep resistance(e.g., ~ 2x10-9s-1at 20 MPa), with a ~ 3 regime prevailing at stresses between ~ 20 and ~ 85 MPa; this high creep resistance is primarily controlled byhigh-aspect- -Al2Cu precipitates which effectively suppress dislocation climb.At higher stresses, ~ 20 prevails, consistent with the activation of dislocation movements within the grain interiors / cores via Orowan looping. Base Al-Cu-Mn-Zr prior to creep exposure exhibits thin -PFZ developed along grain boundaries during --Al2Cuprecipitate-free zones increased significantly during creep, with triple junctions -Al2Cuprecipitate-free zones. Strain localization at / -Al2Cu precipitate-free zones,along with the possible grain-boundary sliding, make the base Al-Cu-Mn-Zr alloy less creep resistant in the ~ 3 regime as compared to Al-5Cu-0.2Mn-0.17Zr-0.25Co- 1.5Ni-0.21Ti-0.15Sb (wt.%). This alloy, commercially known as RR350, exhibits a -dominated intragranular microstructure but a different grain-boundary -Al2Cu precipitate-free zones following creep.

[0100] Adding Ni and Co individually (1.3 wt.% each) or in combination (1.1and 0.16 wt.%, respectively) to Al-Cu-Mn-Zr caused significant microstructural changes only at grain boundaries relative to the base Al-Cu-Mn-Zr alloy. While Ni (or Ni+Co) additions led to an extensive decoration of grain boundaries, individual Co addition exhibited relatively lower grain-boundary decoration efficiency. This is because high Co addition caused some Co-rich intermetallic precipitates to form as a primary phase in the very early stages of solidification, which were found incorporated in the grain interiors upon solidification. The -Al2Cuprecipitate-free zones remained negligible, enabling improved creep resistance inthe ~ 3 regime for all three alloy variants, with combined Ni and Co addition34 168773228.1exhibiting the highest creep resistance, surpassing RR350 at stresses below 50 MPa.

[0101] Adding Fe (1.2 wt.%) to Al-Cu-Mn-Zr produced notable microstructuralchanges in grain boundaries and interiors. While grain interiors exhibited a much- - Al2Cu, as also well reflected in its much lower room- temperature hardness relative to that of a base Al-Cu-Mn-Zr (63 vs.93 HV) alloy, grain boundaries were found well decorated with the coarsening-resistant, irregular- shaped Al7Cu2 -Al2Cu precipitate-free zones. Such microstructure rendered the alloy highly creep resistant, well-surpassing baseAl-Cu-Mn-Zr and comparable to RR350 at stresses up to 40 MPa; however, at higherstresses, its creep resistance deteriorated, producing a ~ 20 regime, which was observed for other alloys at much higher stresses ( > 75-80 MPa) because a much -Al2Cu in Fe-modified alloy led to the activation of Orowan looping at low stress.

[0102] Further increasing the Cu content of the Fe-modified Al-Cu-Mn-Zr, tobalance the Cu consumed in the formation of the Al7Cu2Fe phase, produced acombination of -rich intragranular microstructure, similar to that of base Al-Cu-Mn-Zr alloy, and intermetallic-decorated grain-boundary microstructure. Such-Al2Cu precipitate-free zones during creep exposure, making the alloy highly creep resistant, surpassing the creep performance of RR350 alloy.

[0103] The invention as shown in the drawings and described in detail hereindisclose arrangements of elements of particular construction and configuration for illustrating preferred embodiments of structure and method of operation of the present invention. It is to be understood however, that elements of different construction and configuration and other arrangements thereof, other than those illustrated and described may be employed in accordance with the spirit of the invention, and such changes, alternations and modifications as would occur to those skilled in the art are considered to be within the scope of this invention as broadly defined in the appended claims. In addition, it is to be understood that the phraseology and terminology employed herein are for the purpose of description and 35 168773228.1should not be regarded as limiting. All references cited herein are incorporated fullyby reference. 36 168773228.1

Claims

CLAIMS We claim:

1. An alloy composition comprising:4.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; balance Al, wherein the alloy with the average grain diameters ranging between 10 and 1000 μm produces minimum creep rates from 10-10to 3x10-9s-1at stresses up to 30 MPa, and from 10-10s-1to 2x10-8s-1at stresses up to 70 MPa at 300oC.

2. The alloy composition of claim 1, wherein total amount of impurities is presentat less than 0.5 wt. % based on total weight of the alloy.

3. The alloy composition of claim 1, wherein the impurity is Ag and is present atless than 0.05 wt. %.

4. The alloy composition of claim 1, wherein each impurity element other than Agis present at less than 0.1 wt. % based on the total weight of the alloy.

5. The alloy composition of claim 1, wherein the alloy comprises in weight percent-Al2- Al2Cu phase.

6. The alloy composition of claim 1, comprising 0.05-0.15 wt.% Si based on thetotal weight of the composition. 37 168773228.

17. The alloy composition of claim 1, wherein x= Fe + Ni + Co, and x = 0.5 to 3 wt. % based on the total weight of the alloy.

8. The alloy composition of claim 7, wherein the ratio of Cu / x is from 5:1 to 8:1 in wt. % based on the total weight of the alloy.

9. The alloy composition of claim 8, wherein the concentration of intragranular Cu (Cu content in the grain interiors) is from 2.5 wt. % to 6 wt. %, based on the total weight of the alloy.

10. The alloy composition of claim 1, wherein the volume fraction of intragranular-Al2Cu is within 0.005-0.04 (0.5-4%) based on the total volume of the alloy.

11. -Al2Cu phaseprecipitates and intergranular Al7Cu2(Fe,Co,Ni), Al9(Fe,Co,Ni)2 and Al3(Fe,Co,Ni)2 precipitates.

12. The alloy of claim 1, wherein a linear fraction of grain boundary being coveredby intergranular Al7Cu2(Fe,Co,Ni), Al9(Fe,Co,Ni)2and Al3(Fe,Co,Ni)2precipitates is at least 20%.

13. The alloy composition of claim 1, wherein the alloy consists essentially of:4.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; balance Al.

14. The alloy composition of claim 1, wherein the alloy consists of:38 168773228.14.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; balance Al.

15. A method of making an alloy, comprising the steps of: combining 4.0-24.0 wt. % Cu; 0.5-3.0 wt. % Fe; 0.1-0.5 wt. % Mn; 0.1-0.3 wt. % Zr; 0-0.15 wt.% Si; 0-0.3 wt. % Mg; 0-0.3 wt. % Ti; and balance Al; melting the composition; casting and solidifying the alloy to create a solidified alloy; and,heat-treating the solidified alloy, the heat treating comprising at least one selected from the group consisting of solutionizing heat treatment, quenching in hot water, artificial aging and preconditioning heat treatment.

16. The method of claim 15, wherein the heat-treating comprises a solutionizingheat treatment.

17. The method of claim 16, wherein the solutionizing heat treatment comprisesheating to 530-545oC for at least 4 hours.

18. The method of claim 15, wherein the heat-treating step comprises quenchingin hot water having a temperature of 60-85oC. 39 168773228.

119. The method of claim 15, wherein the heat-treating step comprises aging at230-240oC for at least 5 hours.

20. The method of claim 15, wherein the heat-treating step comprisespreconditioning at 300oC for at least 100 hours. 40 168773228.1

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