Al-Cu-Si-Ge alloys

Abstract

An Al—Cu—Si—Ge quaternary alloy uses Si—Ge additions to provide a reasonably dense and homogeneous distribution of precipitates. Heterogeneous nucleation on this Si—Ge template is used to enhance both strength and thermal stability. These precipitates are used as a template for heterogeneous precipitation of other hardening phases, particularly the •′ (Al2Cu) phase. Thus, a Si—Ge addition is used to provide a dense template of heterogeneous nucleation sites for subsequent precipitation of Al—Cu precipitates.

Description

BACKGROUND OF THE INVENTION

[0003] The invention relates to aluminum alloys, particularly aluminum-copper alloys, and more particularly to thermally stable precipitate hardened aluminum alloys.

[0004] Advanced aerospace and automotive systems need light-weight aluminum alloys that develop high strength and retain it on exposure at moderately high temperature. Aluminum alloy 2219 is considered the leading high temperature stability structural alloy. However, for many applications it does not possess sufficient strength (or hardness which is an indicator of strength). It therefore would be advantageous to develop new precipitation-hardening schemes that satisfy those needs.

[0005] The essential factor that controls the properties of precipitation hardened structural alloys is the type, size and distribution of the strengthening precipitates in the metal matrix. To meet a sufficient level of strength it is important to achieve an adequate volume fraction and number density of strengthening precipitates which are resistant to dislocation shearing.

[0006] Research on the mechanisms of precipitation hardening has identified several criteria that need to be satisfied for optimal hardening. Among them, (1) the precipitate phase should be sufficiently different from the matrix phase that the precipitates become impenetrable to dislocations at very small size, (2) the precipitate species should have sufficient high-temperature solubility to produce a dense distribution of precipitates, and (3) the precipitate size distribution should be as tightly peaked as possible to minimize the spread in particle diameter. Moreover, the precipitate distribution must coarsen relatively slowly if strength is to be retained on exposure to high temperature.

[0007] The first two of these criteria are, ordinarily, conflicting. Precipitates that differ dramatically in structure from the parent phase usually involve species that have low solubility and precipitate heterogeneously in a sparse distribution. The most common precipitation hardening systems employ coherent precipitate phases that are often metastable, but form dense, homogeneous precipitate distributions. Prominent examples in Al alloy systems include θ′ (Al2Cu) and δ′ (Al3Li).

[0008] In commercial, Al—Cu based alloys, the elements Si, Mn, Be, Ge, Sn, Ag, and Cd have all been used to modify the dominant precipitation reaction. One mechanism by which the precipitation reaction is altered is the preferential precipitation of alloy modifiers, producing heterogeneous sites that nucleate Al—Cu or Al—Cu—Mg precipitates. It has been found that trace additions of Sn, Cd, or In promote a dense, homogeneous distribution of fine •′ precipitates (a metastable form of the equilibrium • phase, Al2Cu). In the case of Sn, transmission electron microscopy (TEM) studies show that the Sn precipitates first, providing heterogeneous sites for nucleation of •′. Si behaves in a similar way.

[0009] Si and Ge produce precipitates in Al that have the diamond cubic structure and are impenetrable at very small sizes. However, because of their high lattice strains they do not form dense distributions and are, therefore, not particularly useful for hardening. A number of researchers have attempted to overcome this problem, and have had some success in creating more dense precipitate distributions by using roughly equiatomic combinations of Si and Ge. The partial success of this approach is attributed to the mutual strain compensation of Si and Ge in the Al lattice, which increases their solubility and may promote the formation of Si—Ge clusters that facilitate nucleation. However, the strengths of Al—Ge—Si alloys remain relatively low.

[0010] Since it has been known for some time that alloying Al with a combination of Si and Ge produces a moderately dense distribution of fine, compact precipitates, the Al—Si—Ge system has been proposed as a possible source of thermally stable Al alloys with moderate strength. A reinvestigation of aging behavior in this alloy system shows, however, that it is difficult to achieve precipitate densities that are high enough to make a meaningful contribution to the strength.

[0011] Precipitation hardening in Al—Ge—Si

[0012] The hardness as a function of time for the alloy Al-1Si-1Ge (atomic percent) aged at 160° C. (the temperature previously reported to yield peak hardness) increases slowly to a peak of about 33 HB after 20 hrs. aging. This corresponds to a yield strength of roughly 150 MPa.

[0013] The slow aging kinetics of Al—Ge—Si are surprising, since Si—Ge precipitates have the diamond cubic structure and are impenetrable by dislocations when they are only a few nanometers in diameter. High-resolution transmission electron microscopic studies show why hardening is sluggish. The volume fraction of precipitates increases very slowly with the consequence that the precipitate density is low until the particles have coarsened to significant size. The principal reason for the low precipitation rate is the slow rate of nucleation, though size-dependent solubility via the Gibbs-Thomson effect may also play a role. The need for heterogeneous nucleation, the low nucleation rate, and the size-dependent solubility are all at least partly consequences of the high interfacial tension of the diamond-cubic Si—Ge precipitate in the Al matrix.

[0014] Precipitation hardening in Al—Cu and Al—Si—Cu

[0015] The hardening scheme that is widely used in Al—Cu alloys is the precipitation of the metastable θ′ (Al2Cu) phase. While θ′ is a coherent precipitate that is cut by dislocations when its size is small, the relatively high solubility of Cu and the coherency of the intermetallic θ′ phase have the consequence that a dense, homogeneous θ′ distribution forms with relative ease. A drawback with the θ′ hardening precipitate is its relatively rapid coarsening, which limits its use at higher temperatures. Like almost all coherent precipitates, θ′ is elastically strained, and can minimize its strain energy by spreading out into a wide, thin plate. The decrease of elastic energy as the precipitate thins adds to the driving force for coarsening and accelerates the process.

[0016] The evolution of the θ′ aspect ratio with size in Al2Cu reflects two simultaneous processes. The first is kinetically controlled, and particularly affects very thin precipitates. The θ′ plates are coherent on their broad faces, which lie in {100}Al planes. The plates grow via the nucleation and propagation of discrete ledges. Since ledge nucleation is difficult when the plate is small and thin, some of the precipitates grow into thin plates with very high aspect ratios. The second process is thermodynamically controlled, and particularly governs the shapes of larger plates. The broader the plate, the higher the value of the aspect ratio that minimizes the combination of surface energy and elastic energy. It has been shown that the preferred value of the aspect ratio (k=L/d) increases (approximately) with the square root of the plate diameter (L). The overall evolution of aspect ratio with particle size follows this prediction.

[0017] It has recently been shown that the precipitation of θ′ can be catalyzed and its growth stabilized by the addition of Si to the Al—Cu base alloy. θ′ precipitates form on preexisting Si particles in Al—Cu—Si. The precipitates are more equiaxed than those in the binary and are much slower to coarsen. The apparent reason for the beneficial effect on precipitate morphology and coarsening is strain compensation. The net volume change on forming a Si precipitate is strongly positive, while that associated with θ′ is negative. The elastic energy is minimized when the two precipitates are proximate to one another, catalyzing precipitation. Strain compensation is lost when θ′ grows beyond the strain field of Si. Strain compensation thus lowers the driving force for particle coarsening when the particle size is small, decreasing the coarsening rate.

SUMMARY OF THE INVENTION

[0018] The present invention is an Al—Cu—Si—Ge quaternary alloy, with up to 1 at. % each of Si and Ge, and up to 3.5 at. % of Cu, in an Al metal matrix. While Si—Ge additions to Al alloys do not themselves lead to exceptionally high strengths, they do provide a reasonably dense and homogeneous distribution of precipitates. Since the Al—Si—Ge system provides a finer and more dense distribution of precipitates than the Al—Si binary, heterogeneous nucleation on this Si—Ge template is used to enhance both strength and thermal stability. The present invention uses these precipitates as a template for heterogeneous precipitation of other hardening phases. A particular phase is θ′ (Al2Cu), which is known to be catalyzed by Si precipitates. Thus, a Si—Ge addition is used to provide a dense template of heterogeneous nucleation sites for subsequent precipitation of Al—Cu precipitates.

[0019] The resulting Al—Cu—Si—Ge alloys have excellent combinations of strength and thermal stability in the aged condition. In particular Al—Cu—Si—Ge quaternary alloys with nominal compositions Al-2Cu-1Si-1Ge and Al-2.5Cu-0.5Si-0.5Ge were tested. On aging, both alloys rapidly develop dense distributions of •′. Particle-by-particle analyses show that these precipitates nucleate heterogeneously on pre-existing Si—Ge precipitates. The precipitate distributions exhibit good thermal stability, coarsening slowly.

BRIEF DESCRIPTION OF THE DRAWINGS

[0020] FIG. 1A is a bright-field image near the [110]Al zone axis for Al-2 at. pct. Cu-1 at. pct. Si-1 at. pct. Ge alloy aged for 3 h at 190° C.

[0021] FIGS. 1B, C are dark-field images of θ′ precipitates obtained using 112•′ and 112•′, respectively, with Si—Ge nucleation sites for θ′ indicated by arrows.

[0022] FIG. 2 shows the hardening curves for the two alloys, Al-2Cu-1Si-1Ge and Al-2.5Cu-0.5Si-0.5Ge, together with alloy 2219 (Al-6.3 wt. % Cu-0.3Mn-0.18Zr-0.1V-0.06Ti, or Al-2.8 at. % Cu-0.15Mn-0.05Zr-0.055V-0.035Ti) and alloy 2014 (Al-4.4 wt. % Cu-0.8Si-0.8Mn-0.5Mg or Al-1.94 at. % Cu-0.8Si-0.4Mn-0.6 Mg).

[0023] FIG. 3 shows a high resolution images of the homogeneous nucleation of θ′ precipitate on a Si—Ge particle in Al-2Cu-1Si-1Ge alloy.

[0024] FIG. 4 shows bright field and dark field micrographs of θ′ precipitates in different orientations, after 1 and 3 hours aging at 190° C.

[0025] FIG. 5 shows dark field images of θ′ precipitates tilted such that the regions where they are attached to the Si—Ge can be clearly seen.

[0026] FIG. 6 is a plot of the solubility of Cu in Al, alone and in the presence of Si, as a function of temperature.

DETAILED DESCRIPTION OF THE INVENTION

[0027] The invention is a quaternary aluminum-copper-silicon-germanium (Al—Cu—Si—Ge) alloy having the general composition Al-xCu-ySi-zGe where x, y, z are the atomic percentages. In general, the atomic percentage of Cu is greater than zero and ranges up to about 3.5%. The atomic percentages of Si and Ge are each greater than zero and each range up to about 1%; the atomic percentages of Si and Ge are more preferably up to about 0.5%.

[0028] The strength of a dispersion hardened alloy scales as the inverse of the particle spacing. Thus, an optimal high strength microstructure consists of a dense dispersion of small unshearable particles. The Al—Cu—Si—Ge system relies on an extremely numerically dense distribution of ultra-fine Si—Ge precipitates as a template for heterogeneous nucleation of other strengthening precipitates. In this case, Si—Ge particles are used as a template for the formation of θ′ precipitates that are also are very fine and densely distributed.

[0029] Bulk alloys of composition Al-2 at. % Cu-1 at. % Si-1 at. % Ge and Al-2.5 at. % Cu-0.5 at. % Si-0.5 at. % Ge, were made by arc melting, 99.999 (wt. %) Si, 99.9999 (wt. %) Ge, 99.999 (wt. %) Cu and 99.99 (wt. %) Al. The samples then were cold swaged to achieve 10 to 15% plastic deformation. They were then encapsulated in a sealed quartz glass tube that was back-filled with argon and annealed for 24 hours at 500° C. and quenched into ice water. The final shape of the bulk alloy was roughly cylindrical, approximately 20 mm in length and 10 mm in diameter. The cylindrical ingots were sliced into discs 0.5 mm in thickness, cut normal to the cylinder axis.

[0030] TEM samples were electrochemically polished in 75% methanol-25% HN03 solution at a temperature of −25° C. with a polishing voltage of around 20 V.

[0031] Conventional TEM was performed using a JEOL 200 CX at 200 kV. Energy dispersive X-ray spectroscopy (EDX) was done using a JEOL 2000X and Philips CM200-FEG Analytical Transmission Electron Microscopes, equipped with light element detectors and operated at 200 kV. High Resolution TEM was performed on the JEOL ARM operated at 800 kV and a Philips CM300-FEG operated at 300kV.

[0032] Al—Cu—Si—Ge samples were aged for varying times at 190° C. This temperature corresponds to the standard heat treatment given to 2000 series (Al—Cu based) alloys.

[0033] After aging the Al—Cu—Si—Ge samples were indented using the Rockwell Hardness B ({fraction (1/16)}th inch steel ball, 100 kg load) scale. Rockwell Hardness B values were converted to HB using standard ASTM tables for aluminum.

[0034] FIGS. 1A-C show the microstructure of Al-2Cu-1Si-1Ge, after aging for 3 hours at 190° C. FIG. 1A is a bright-field image of the microstructure near the [110]Al zone axis. Visible are both plates, identified in dark field in FIG. 1B as edge-on θ′ and spherical Si—Ge particles. From FIGS. 1A, B, it can be observed that both phases are densely distributed and are relatively fine and uniform in size. FIG. 1C shows a dark-field image of the θ′ precipitates oriented approximately 35 deg. to the foil normal. They are imaged in the [110]θ′ zone axis, which is oriented 10 deg. away from the [110]Al zone axis tilted along the 200 Kikuchi lines. The θ′ precipitates are growing around the Si—Ge particles, giving the appearance in dark field of the θ′ containing holes. The most dramatic examples of this are indicated by arrows.

[0035] FIG. 2 shows the hardening behavior for two quaternary Al—Cu—Si—Ge alloys of the invention, Al-2Cu-1Si-1Ge and Al-2.5Cu-0.5Si-0.5Ge, together with alloy 2219 (Al-6.3 wt. % Cu-0.3Mn-0.18Zr-0.1V-0.06Ti, or Al-2.8 at. % Cu-0.15Mn-0.05Zr-0.055V-0.035Ti) and alloy 2014 (Al-4.4 wt. % Cu-0.8Si0.8Mn-0.5Mg or Al-1.94 at. % Cu-0.8Si-0.4Mn-0.6Mg). The quaternary alloys harden more quickly than 2014 or 2219, and maintain high strength for long times. The rapid aging is due to the catalyzed precipitation. The thermal stability appears to be due to strain compensation.

[0036] Alloys 2219 and 2014 are used for comparison since they represent two Al—Cu based alloys known to exhibit high strength in their T-6 (solutionized, quenched and artificially aged) condition. The most promising feature of the Al—Cu—Si—Ge alloys is their combination of fast aging response, high hardness and good high temperature stability. It is apparent that the alloy with the higher Cu content exhibits higher peak hardness than the alloy with the higher Si and Ge content. This is expected because the hardening potential of these alloys is related to the amount and precipitation kinetics of θ′ and not to the amount of the diamond cubic Si—Ge precipitates whose presence does not result in appreciable improvement in mechanical properties.

[0037] At 190° C. alloy 2014, which is not known to display very good high temperature stability, quickly deteriorates. It should be noted that by aging 2014 at 160° C. for 13 hours it is possible to obtain a Brinell Hardness of 122. However, the hardness then quickly drops off with prolonged aging.

[0038] Compared to 2219 both Al-2Cu-1Si-1Ge and Al-2.5Cu-0.5Si-0.5Ge possess a higher peak hardness. Both alloys reach maximum hardness after only 3 hours, instead of the 8 hours necessary for 2219. With prolonged aging time, the Al—Cu—Si—Ge alloys overage at a rate similar to 2219, and after approximately 400 hours at elevated temperature the hardness of all three alloys decreases asymptotically to approximately 86 HB. This is a very promising result since alloy 2219 is known for its resistance to averaging.

[0039] FIG. 3 is a high resolution electron microscope (HREM) image of the edge-on θ′ precipitate in contact with a multiply twinned Si—Ge particle. The sample was aged for 1 hour, and the image was taken in the [001]Al zone axis. The segments B and D of the Si—Ge particle have the Baker-Nutting orientation relationship with aluminum the matrix.

[0040] FIG. 4 shows bright field (BF) and dark field (DF) micrographs of •′ precipitates in different orientations, after 1 hour aging (left side) and 3 hours aging (right side) at 190° C. for the Al-2Cu-1Si-1Ge alloy. It is important to note the low aspect ratio and the fine size of the •′ imaged edge on the two central dark field images.

[0041] FIG. 5 shows dark field images of •′ precipitates tilted such that the regions where they are attached to the Si—Ge can be clearly seen (arrows), the Si—Ge particles being visible as holes in the •′. The Al-2Cu-1Si-1Ge alloy is aged for 1 h (left side) and 3 h (right side) at 190° C. Since the aspect ratio (length/thickness) varies as the square root of precipitate length, because of their small size the •′ are short and thick. This makes them resistant to being sheared during deformation and to coarsening at high temperatures, both of which would degrade mechanical properties.

[0042] Another advantage of this alloy system is the enhanced solubility of Cu in Al in the presence of Si. FIG. 6 is a plot of the solubility of Cu in Al as a function of temperature. At the solutionizing temperature of 500° C., 1 at. % Cu is soluble in Al for the binary alloy, while 2.5 at. % Cu is soluble in Al for the ternary. The higher solubility results in a larger volume fraction of precipitates during aging, further improving alloy strength.

[0043] Al—Cu—Si—Ge displays a uniquely fast aging response, a high peak hardness and a good stability during prolonged aging. The high hardness of the Cu containing alloy is due to the dense and uniform distribution of fine θ′ precipitates (metastable Al2Cu) which are heterogeneously nucleated on the Si—Ge particles. High resolution TEM demonstrates that all the Si—Ge precipitates start out, and remain multiply twinned throughout the aging treatment. Since the twinned section of the precipitate does not maintain a low index interface with the matrix, the Si—Ge precipitates are equiaxed in morphology.

[0044] Al—Cu—Si—Ge alloys display a far superior peak hardness compared to alloy 2219, while having equal, if not better stability after extended aging at high temperatures. Additionally, Al—Cu—Si—Ge requires less aging time to achieve maximum hardness than 2219, making this new class of alloys less expensive to heat treat. The fast initial hardening response will also make this alloy very attractive for applications where the structure will undergo multiple pass welds, since each subsequent weld will harden the heat effected zone.

[0045] This class of alloys could be used in any structural application requiring good peak hardness and exceptional high temperature stability. Two classic applications are aerospace and automotive. Specific applications include, but are not limited to, heavy duty forgings, plate and extrusions for aircraft fittings, wheels, space booster tankage and structure, truck frame and suspension components, rivets, screw machine products, aircraft and automotive pistons, aircraft engine cylinder heads, jet engine impellers, compressor rings, aircraft skin and other structural applications.

[0046] Changes and modifications in the specifically described embodiments can be carried out without departing from the scope of the invention which is intended to be limited only by the scope of the appended claims.

Claims

1. A composition of matter, comprising a quaternary aluminum-copper-silicon-germanium alloy.

2. The composition of claim 1 wherein the alloy comprises aluminum-copper precipitates heterogeneously nucleated on silicon-germanium precipitates in an aluminum metal matrix.

3. The composition of claim 2 wherein the aluminum-copper precipitates are θ′ (Al2Cu) precipitates.

4. The composition of claim 1 having the general formula Al-xCu-ySi-zGe wherein x, y, z are the atomic percentages of the constituents.

5. The composition of claim 4 wherein y and z are each up to about 1 at. %, and x is up to about 3.5 at. %.

6. The composition of claim 5 wherein y and z are each up to about 0.5 at. %.

7. The composition of claim 4 wherein the alloy comprises aluminum-copper precipitates heterogeneously nucleated on silicon-germanium precipitates in an aluminum metal matrix.

8. The composition of claim 7 wherein the aluminum-copper precipitates are θ′ (Al2Cu) precipitates.

9. The composition of claim 5 wherein the alloy comprises aluminum-copper precipitates heterogeneously nucleated on silicon-germanium precipitates in an aluminum metal matrix.

10. The composition of claim 9 wherein the aluminum-copper precipitates are θ′ (Al2Cu) precipitates.

11. A method of forming a precipitate hardened aluminum alloy, comprising melting aluminum, copper, silicon, and germanium together to first form silicon-germanium precipitates which serve as nucleation sites, and then form aluminum-copper precipitates on the silicon-germanium nucleation sites.

12. The method of claim 11 wherein the atomic percentage of Cu is greater than zero and ranges up to about 3.5%, and the atomic percentages of Si and Ge are each greater than zero and each range up to about 1%.

13. The method of claim 12 wherein the atomic percentages of Si and Ge are each up to about 0.5%.

14. The method of claim 11 further comprising annealing the alloy.

15. The method of claim 11 comprising annealing at about 500° C. for about 24 hours.

16. The method of claim 14 further comprising aging the annealed alloy.

17. The method of claim 16 comprising aging at about 190° C. for about 1-3 hours.

18. A precipitate hardened aluminum alloy formed by the method of claim 11.

19. A precipitate hardened aluminum alloy formed by the method of claim 12.

20. A precipitate hardened aluminum alloy formed by the method of claim 17.