WO2000000653A1 - Method of producing superplastic alloys and superplastic alloys produced by the method - Google Patents

Abstract

A method for producing new superplastic alloys by inducing in an alloy the formation of precipitates having a sufficient size and homogeneous distribution that a suffiently refined grain structure to produce superplasticity is obtained after subsequent PSN processing. An age-hardenable alloy having at least one dispersoid phase is selected for processing. The alloy is solution heat-treated and cooled to form a supersatured solid solution. The alloy is plastically deformed sufficiently to form a high-energy defect structure useful for the subsequent heterogeneous nucleation of precipitates. The alloy is then aged, preferably by a multi-stage low and high temperature process, and precipitates are formed at the defect sites. The alloy then is subjected to a PSN process comprising plastically deforming the alloy to provide sufficient strain energy in the alloy to ensure recrystallization, and statically recrystallizing the alloy. A grain structure exhibiting new, fine, equiaxed and uniform grains is produced in the alloy. An exemplary 6xxx alloy of the type capable of being produced by the present invention, and which is useful for aerospace, automotive and other applications, is disclosed and claimed. The process is also suitable for processing any age-hardenable aluminum or other alloy.

Description

METHOD OFPRODUCING SUPERPLASTICALLOYS AND SUPERPLASTIC ALLOYS PRODUCED BYTHE METHOD

CROSS-REFERENCE TO RELATED PROVISIONAL APPLICATION The present application claims the benefit of the earlier filing date of U.S.

Provisional Patent Application Serial No. 60/089,239, filed June 15, 1998, which is incorporated by reference herein in its entirety. STATEMENT CONCERNING FEDERALLY SPONSORED RESEARCH

This invention was made with government support under NASA Training Grant No. NGT-1-52117. The U.S. Government has certain rights in the invention. FIELD OF THE INVENTION

The present invention relates to a method for producing fine-grained alloys, particularly fine-grained 6xxx aluminum alloys which exhibit superplasticity, and to the

alloys produced by the method. BACKGROUND OF THE INVENTION

The advantages of superplastic properties in metals are well known, and particularly well employed in the automotive and aerospace industry. Because of their finegrained microstructures, superplastic metals and alloys may exhibit from several hundred percent to several thousand percent elongation without necking when pulled in tension at temperatures exceeding 0.5 T_m, where T_m is the absolute melting temperature of the material. In contrast, non-superplastic metals and alloys typically elongate less than 100% before necking under similar conditions. Accordingly, superplastic metals may be formed into a

multitude of complex shapes not achievable with other metals.

Currently, commercial interest in the aerospace and automotive industries is

focused on superplastic forming ("SPF"). SPF is a manufacturing process which exploits the

phenomenon of superplasticity by using low gas pressures (less than about 1000 psi (7 MPa)), and concomitantly low energies, to form parts having complex shapes. This process reduces part counts and the need for fasteners and connectors, reducing product weight and manufacturing costs. In addition, SPF may be performed using a single surface tool in a single forming operation, thus reducing tooling costs. The advent of SPF therefore increases the potential commercial applications in which superplastic materials may be employed.

Superplastic behavior in metallic alloys may be described by the equation σ = kέ^m where σ = flow stress, k = material constant, έ = strain rate, and m = strain rate sensitivity. In superplastic metals, m usually ranges from about 0.4 to 0.8. "Quasi-superplastic" metals and alloys have m values of around 0.33. Materials having m values less than 0.3 are considered to be non-superplastic.

Most metals and alloys capable of achieving superplasticity must be specially processed for superplasticity. The microstructures of such metals and alloys may be refined through thermomechanical processing to impart such properties to the material. For a material to be superplastic, it is typically refined to possess an equiaxed, fine-grained structure, typically with grains about 20 μm or less in diameter and preferably about 10 μm or less. In addition, for such a material to be commercially useful, it must be statically stable such that its grains do not experience significant growth at superplastic forming temperatures. Where the thermomechanical process for refinement includes static recrystallization, which is a common component of such processes, a weak or random texture and the presence of predominantly high-angle grain boundaries is also required. The development of thermomechanical processes effective for creating alloys having such properties has proven to

be extremely challenging.

An extensive amount of research has been conducted in an effort to discover thermomechanical processes useful for producing superplastic alloys, including aluminum alloys. This work has resulted in the development of several superplastic alloys, but

undoubtedly, many commercially important superplastic alloys have yet to be discovered. In

particular, although several superplastic 2xxx, 5xxx, 7xxx and 8xxx aluminum alloys have

been produced, there has been a significant deficiency in successful research concerning the

grain refinement and superplasticity of 6xxx aluminum alloys. New superplastic 6xxx aluminum alloys would be particularly desirable, because 6xxx alloys are highly weldable,

corrosion resistant, extrudable and low in cost compared with other aluminum alloys. Thus, there is a need for the development of methods for imparting superplastic properties to alloys,

particularly 6xxx aluminum alloys.

Of the 6xxx aluminum alloys, 6061, 6063, 6066, and especially 6013 and

6111, possess substantial promise for extensive use in the aerospace and automotive industries. Indeed, non-superplastic aluminum alloy 6013, a medium strength, age- hardenable alloy developed by ALCOA in the early 1980s, has been selected for use on

Boeing Co.'s state-of-the-art 777 aircraft, as well as for many other automotive and aerospace applications. This is not surprising, given the favorable properties of this alloy and the fact

that it can be processed to develop properties superior to other 6xxx alloys. For example, it

has corrosion resistance superior to that of 2xxx and 7xxx aluminum alloys, which are

heavily used for aerospace applications. The yield strength of 6013-T6 is 12% higher than

that of 2024-T3, it is nearly immune to corrosion that results in exfoliation and stress-

corrosion cracking, and it is 25% stronger than 6061-T6. In addition, the alloy 6013-T4 has

better stretch-forming characteristics than other aerospace aluminum alloys. Accordingly,

there is a need for the development of methods for imparting superplastic properties to 6061,

6063, 6066 alloys, and particularly to 6013 and 6111 aluminum alloys. To date, efforts expended to impart superplasticity to 6xxx aluminum alloys have not been very successful. United States Patent No. 4,092,181 to Paton, et al., which describes what is known in the art as the "Rockwell process," discloses a method for imparting a fine grain structure to aluminum alloys having precipitating constituents. The thermomechanical process of the Paton, et al. method consists of solution heat treating such an alloy, overaging the alloy, then subjecting the alloy to a particle-stimulated nucleation ("PSN") process during which the alloy is mechanically worked and recrystallization is induced. Although the Paton, et al. patent provides several examples of the method described therein, it does not describe the microstructures produced by the method, nor does it suggest that superplastic results were achieved. Indeed, experimental evidence available in the literature indicates that the method disclosed by Paton, et al. is not very useful for imparting superplasticity to 6xxx alloys. This is confirmed by the work performed in connection with

the present invention, as described below.

Similarly, Washfold, et al. attempted to grain refine a 6063 aluminum alloy through PSN in order to induce superplasticity. See Washfold, et al., "Thermomechanical

Processing of an Al-Mg-Si Alloy," Metals Forum (1985) at 56-59. The thermomechanical process used is very different than that employed in the present invention, and consists of a solution heat-treatment followed by slow cooling to an overaging temperature, overaging, slow cooling to room temperature, cold or warm rolling, and static recrystallization with a slow heat-up to the recrystallization temperature. Washfold, et al. produced a microstructure exhibiting a minimum grain diameter of 10.5 μm (in the rolling plane), as measured using optical microscopy ("OM") techniques. They obtained a maximum elongation of 148% at 450°C, due to significant grain growth occurring at 500°C and above, within the superplastic forming temperature range. The Washfold, et al. process did not achieve superplasticity. Kovacs-Csetenyi, et al. attempted to use compositional variation and

thermomechanical processing to refine the grain structure and improve the superplastic performance of aluminum 6066 and three variants of aluminum 6061. See Kovacs-Csetenyi, et al, "Superplasticity of AlMgSi Alloys," Journal of Materials Science 27 (1992) at 6141-45.

The thermomechanical process used consists of solution heat-treatment followed by

overaging, rolling, and static recrystallization, and bears no resemblance to that of the present invention. Kovacs-Csetenyi, et al. report strain rate sensitivity values in the range of 0.4 for

each of the four alloys processed, as studied using temperatures between 500 °C and 570 °C and strain rates of

10^"3 to 10^"6 s^"1, indicating that some degree of superplastic behavior would be expected from the alloys. However, superplasticity was characterized using impression creep tests, and no

uniaxial tensile tests were reported. Thus, it is unclear what amounts of superplastic elongation, if any, were obtained by the processing technique described in this reference.

Chung, et al. also experimented with grain refinement techniques to produce a

superplastic 6013 alloy. See Chung, et al., "Grain Refining and Superplastic Forming of

Aluminum Alloy 6013," The 4^th International Conference on Aluminum Alloys (1994), 434-

42. Chung, et al. employed a thermomechanical process consisting of solution heat-

treatment, 10% cold rolling, overaging at 380?C, 90% warm rolling at 190°C, and

recrystallization. In contrast to the process of the present invention, Chung, et al. employed

mild cold rolling, for the purpose of forming a dislocation network to assist in the

precipitation of what was thought to be Mg₂Si precipitates. The process resulted in grains of

12 to 13 μm (measured using optical microscopy techniques), a strain rate sensitivity of 0.38,

and a maximum elongation of 230% at 520 °C for a strain rate of 3 x 10^"4 s^*1, and at a flow

stress of 972 psi (6.7 MPa). Thus, the product of the Chung, et al. process was only marginally superplastic. Chung, et al. concluded that the size and number of iron-bearing constituents in the alloy needed to be reduced in order to achieve more favorable results.

Chung, et al. clearly were not aware that, as disclosed by the present invention, a significantly

higher energy deformation structure such as a deformation band needed to be imparted to the

material and exploited to form sites for the heterogeneous nucleation of precipitates, enabling

the achievement of a superplastic microstructure.

A similar process to that employed by Chung, et al., but directed to an

altogether different purpose, is described in U.S. Patent No. 3,706,606 to DiRusso, et al. The DiRusso patent addresses the need to develop processes for increasing the mechanical strength of semifinished aluminum alloys. Like Chung, et al., the DiRusso patent describes

using a mild cold or warm rolling between solution heat-treatment and aging steps to provide a dislocation network to assist in precipitation. None of the alloys treated using the process of the DiRusso patent exhibited superplastic properties, as shown by the tensile elongation

tests performed by DiRusso, et al. on such alloys, nor were they intended to do so. Accordingly, it is an object of the present invention to provide alloys exhibiting superplasticity, particularly 6xxx alloys and especially aluminum 6013 and 6111

alloys.

It is another object of the present invention to provide a method for imparting superplastic properties to alloys that is applicable to a wide range of alloys, particularly all

6xxx alloys and especially aluminum 6013 and 6111 alloys.

It is yet another object of the present invention to provide a method for

imparting superplastic properties to alloys that is economical and commercially useful.

It is still another object of the present invention to provide a method for

producing superplastic alloys having an equiaxed, uniform, thermally stable, fine grain structure of less than about 20 μm, and preferably about 10 μm or less.

It is another object of the present invention to provide a method for producing

superplastic alloys having a microstructure with a weak or random texture and a predominance of high-angle grain boundaries. SUMMARY OF THE INVENTION

In accordance with the principles of the present invention, alloys exhibiting

superplasticity and a method for producing the same are provided. The method involves

inducing in an alloy the formation of precipitates having a sufficient size and homogeneous distribution such that, after a subsequent PSN process, a sufficiently refined grain structure to

produce superplasticity results. The process of the present invention differs from previous processes in the particular thermomechanical processing steps required, as well as in the sequence and character of those steps. Because of these differences, the process of the

present invention is capable of imparting to age-hardenable alloys, and particularly to age-

hardenable aluminum alloys, exceptional superplastic characteristics heretofore not obtainable. An exemplary alloy of the type capable of being produced by the present

invention is a superplastic 6xxx alloy which is economically produced and commercially

useful for aerospace, automotive and other applications.

The method for producing a superplastic alloy, as provided by the present invention, comprises providing an age-hardenable alloy for processing which has a matrix

phase and at least two alloying elements, at least one of the alloying elements being, or being

capable of forming, a dispersoid phase substantially insoluble in the matrix phase after basic

ingot processing. The alloy is solution heat-treated, and cooled to form a supersaturated solid

solution. The alloy is then plastically deformed sufficiently to form a high-energy defect

structure, thereby forming nucleation sites useful for the subsequent heterogeneous nucleation of precipitates. The alloy is then aged, forming precipitates at the nucleation sites, and subjected to deforming and recrystallizing through a PSN process.

This process has been shown to effect excellent results in a variant of an aluminum 6013/6111 alloy, but is suitable for processing any age-hardenable alloy.

Aluminum alloys, particularly 6xxx aluminum alloys, and more particularly 6013, 6111,

6061, 6063 and 6066, are particularly good candidates for processing under the present method.

The cooling step following solution heat-treatment may be performed using any mode of rapid cooling. For example, it may be performed by quenching in media such as

water, oil or air. The step of plastically deforming the alloy must be sufficiently severe to form a high-energy defect structure, such as the high-energy defect structures commonly

referred to as "deformation bands," in contrast to lower-energy defect structures such as a

dislocation network. Such severe plastic deformation may be imparted by any means, such as

a rolling, stretching, extrusion, drawing, forging or torsion process at economical temperatures and conditions, and is preferably imparted by cold rolling at room temperature.

The aging process of the present invention may comprise a single heating step in which the alloy is heated at a single temperature for a set period of time, or multiple

heating steps in which the alloy is heated at different temperatures over set time periods.

Preferably, the aging process comprises a first heating step at a first temperature and a second heating step at a second higher temperature. The first heating step may be used to form the

precipitates, which then may be coarsened during the second heating step. Where two or

more heating steps are used, the alloy preferably is cooled after each heating step.

The PSN process preferably includes plastically deforming the alloy to provide

sufficient strain energy in the alloy to ensure recrystallization, and statically recrystallizing the alloy. The plastic deformation step of the PSN process may include any mode of plastic

deformation, but preferably comprises cold rolling the alloy at room temperature. The static recrystallization step of the PSN process preferably includes rapidly heating the alloy to a

temperature at which recrystallization occurs and at which recovery is minimized. In one

embodiment, such rapid heating is provided by selecting a recystallization temperature in the

range of the solution heat-treatment temperature for the alloy. In another embodiment, rapid

heating is provided by heating the alloy to the superplastic forming temperature of the alloy.

One of the alloys which may be processed to exhibit exceptional superplastic properties using the method of the present invention is a 6013/6111 aluminum alloy having

the approximate composition 97.3 wt % Al - 0.8 wt % Mg - 0.7 wt % Si - 0.8 wt % Cu - 0.3

wt % Mn - 0.1 wt % Fe. In one embodiment of the present invention, the solution heat- treating step is performed by heating this alloy at a temperature of about 540 °C for about one

hour, excluding heat-up time. The solution heat-treated alloy is then rapidly cooled,

preferably by cold water quenching. The alloy is then plastically deformed to a sufficient

degree to form the required deformation bands or other high-energy defect structures in the material. This may be done, for example, by cold rolling at room temperature by about 30% or more. Most preferably, the plastic deformation is performed such that, after subsequent

aging, the alloy will exhibit a uniform distribution of globular or near-spheroid shaped

precipitates. Aging may be performed using any combination of aging steps, but preferably is

performed using a two-step aging process. In one embodiment of the invention, a first

heating step is performed at about 300°C for about 24 hours and a second heating step is

performed at about 380°C for about 24 hours, with the alloy being cooled after the each of

the heating steps. Precipitates preferably are formed during the first heating step and

coarsened during the second heating step. According to another exemplary embodiment, the 6013/6111 superplastic aluminum alloy of the present invention may be aged using a first heating step at about 300°C

for about 24 hours, and a second heating step at about 450 °C for about 2 hours. Under yet

another exemplary embodiment, the alloy may be aged using a single heating step, at a

temperature of about 450 °C for about 2 hours. Although the microstructure of this single- heating step alloy may be somewhat less ideal than those of the alloys produced using the dual heating steps of the other exemplary embodiments, such a low temperature/short heating

time process may be preferred for commercial applications where energy consumption and time are important factors.

After aging, the 6013/6111 aluminum alloy of the present invention is plastically deformed to provide sufficient strain energy in the alloy to ensure recrystallization.

In one embodiment of the invention, the alloy is cold rolled at room temperature by about

80% or more. In particular, cold rolling at room temperature by about 80%, 87% and 92% has produced exceptional results. Smaller amounts of plastic deformation may also be employed. The alloy is then recrystallized. In connection with the recrystallization step, the

alloy should be rapidly heated to the temperature at which recrystallization occurs to minimize recovery within the deformation zones around the precipitates and to activate the

largest number of recrystallized nuclei. In one embodiment of the invention, the alloy is

rapidly heated to about 540°C and held there for about five minutes.

Processing the 6013/6111 aluminum alloy as discussed yields a superplastic

alloy with a microstructure having a fine average grain size in the range of about 9.5 μm to

about 11.6 μm, the grain sizes having a standard deviation in the range of about 4.7 μm to

about 5.6 μm. In addition, the microstructure of the alloy has a low average grain aspect ratio

(i.e., ratio of major axis to minor axis) in the range of about 1.6 to about 1.9, the grain aspect ratios having a standard deviation in the range of about 0.6 to about 0.8. The alloy also has a grain roundness in the range of about 1.6 to about 1.8, a maximum strain rate sensitivity of at

least about 0.5, and a maximum elongation capability of at least about 350%, preferably

375% or more. Specifically, in one embodiment, processing the 6013/6011 alloy using a first heating step at about 300°C for about 24 hours and a second heating step at about 380°C for

about 24 hours, with the alloy being cooled after the each heating step, and subsequently cold

rolling the aged alloy by about 87% and recrystallizing the alloy at about 540 °C for about five

minutes, yields an average grain size of about 9.5 μm (about 4.7 μm standard deviation), and

an average grain aspect ratio of about 1.6 (about 0.6 standard deviation). The resulting alloy

has a maximum strain rate sensitivity of about 0.5 at 540 °C for a strain rate range of 2x10^"4 s^"1 to 5X10^"4 s^"1, and a maximum elongation of about 375% with a corresponding maximum stress of approximately 680 psi (4.7 MPa).

The foregoing and other features, objects and advantages of the present invention will be apparent from the following detailed description, taken in connection with

the accompanying figures, the scope of the invention being set forth in the appended claims.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a SEM micrograph (150x) of a 6013/6111 alloy, produced in

accordance with the method of the present invention, following solution heat treatment.

FIG. 2a is a SEM micrograph (500x) illustrating banded deformation

structures produced in 30% cold rolled sample E in accordance with the method of the

present invention;

FIG. 2b is a SEM micrograph (500x) illustrating banded deformation

structures produced in 60% cold rolled sample A in accordance with the method of the present invention;

FIG. 3a is a SEM micrograph (5000x) illustrating globular or near-spheroid

shaped precipitates as produced in sample A in accordance with the method of the present invention;

FIG. 3b is a SEM micrograph (5000x) illustrating globular or near-spheroid

shaped precipitates as produced in sample B in accordance with the method of the present invention;

FIG. 3c is a SEM micrograph (5000x) illustrating globular or near-spheroid shaped precipitates as produced in sample C in accordance with the method of the present invention;

FIG. 4a is a TEM micrograph (3300x) of sample D after the aging step of the present invention;

FIG. 4b is a TEM micrograph (3300x) of sample A after the aging step of the

present invention; FIG. 5a is a SEM micrograph (lOOOx) illustrating the distribution of

precipitates in an 8% stretched 6013/6111 sample heated at 380° C for 17 hours;

FIG. 5b is a SEM micrograph (500x) illustrating the distribution of

precipitates in sample A, produced in accordance with the method of the present invention;

FIG. 6 is a SEM micrograph (200 μm width x 150 μm height) illustrating the

grain size of sample A processed using optimized downstream processing conditions in

accordance with the method of the present invention;

FIG. 7 is a 10° misorientation grain boundary map (200 μm width x 150 μm

height) corresponding to the SEM micrograph of FIG. 6;

FIG. 8a is a SEM micrograph (150x) illustrating the recrystallized grain structure of sample A produced in accordance with the method of the present invention;

FIG. 8b is a SEM micrograph (500x) illustrating the recrystallized grain

structure of sample A produced in accordance with the method of the present invention;

FIG. 9a is a SEM micrograph (150x) illustrating the recrystallized grain structure of sample B produced in accordance with the method of the present invention;

FIG. 9b is a SEM micrograph (500x) illustrating the recrystallized grain structure of sample B produced in accordance with the method of the present invention;

FIG. 10a is a SEM micrograph (150x) illustrating the recrystallized grain structure of sample C produced in accordance with the method of the present invention;

FIG. 10b is a SEM micrograph (500x) illustrating the recrystallized grain structure of sample C produced in accordance with the method of the present invention;

FIG. 1 la is a SEM micrograph (150x) illustrating the recrystallized grain

structure of sample E produced in accordance with the method of the present invention;

FIG. 1 lb is a SEM micrograph (500x) illustrating the recrystallized grain structure of sample E produced in accordance with the method of the present invention;

FIG. 12 is a graph illustrating the variation of strain rate sensitivity with strain rate for uniaxial, step strain rate tests of sample A, produced in accordance with the method

of the present invention, at 500°C and 540°C;

FIG. 13 is a graph illustrating the variation of elongation with strain rate for

sample A, produced in accordance with the method of the present invention, at a temperature

of 540°C ; and

FIG. 14 is a photograph of an undeformed sample alongside samples deformed

to 350 to 375%, each of the samples representing sample A, produced in accordance with the

method of the present invention. DETAILED DESCRIPTION OF THE INVENTION

The preferred embodiments of the method of the present invention, and the

alloys produced in accordance with the present invention, will now be described.

Providing An Alloy According to the method of the present invention, an alloy must be provided for processing. Any age-hardenable alloy, such as a 2xxx, 6xxx, 7xxx and some 8xxx aluminum alloy, conceivably is a candidate for processing in accordance with this invention.

The alloy must include a matrix phase and at least two alloying elements, at least one of the

alloying elements being, or being capable of forming, an insoluble dispersoid phase present as particles typically less than one micron in diameter which are substantially insoluble in the matrix phase of the alloy. The dispersoids are utilized by the present invention during

recrystallization to help retain a fine grain structure by limiting grain growth.

Although the process does not require any particular alloy composition, it has been demonstrated to work particularly well for a variant of an aluminum 6013/6111 alloy having the approximate composition 97.3 wt % Al - 0.8 wt % Mg - 0.7 wt % Si - 0.8 wt % Cu

- 0.3 wt % Mn - 0.1 wt % Fe. The alloy was cast and ingot-processed by Reynolds Metals Company at Reynolds' Richmond, Virginia facility. One half of the ingot was preheated in

the conventional manner using a heat-up rate of about 50°C/hour, a soak temperature of

about 560°C, and a soak time of about four hours. The other half of the ingot underwent a

low-temperature preheat (about 500°C) using a heat-up rate of about 50°C/hour and a soak

time of about eight hours, to achieve a finer size distribution and slightly higher volume

fraction of dispersoids than that obtained using the conventional preheat. Each ingot was

then rolled to form an approximately 1 " thick plate.

It should be noted that the terms "about" or "approximately," as used in the present application, are intended to encompass values within ± 25% of the stated value.

Solution Heat-Treatment The alloy selected for processing is solution heat-treated in the conventional manner. It will be readily appreciated that the temperature and heating time of this step

depend upon the type and thickness of the alloy being processed, and that for standard alloys,

these parameters may be readily ascertained from the alloy's manufacturer or material data sheet. In any event, the alloy should be heated to a temperature below that at which melting

begins, and the heating time should be sufficient to achieve the dissolution of all normally soluble phases. For the 1 " thick plate samples discussed above, an air furnace was preheated

to a temperature of about 540 °C. The samples were placed in the furnace for a period of about one hour, excluding heat-up time. A SEM micrograph (150x) of a sample of this

material following solution heat- treatment is shown in Figure 1.

Rapid Cooling Following solution heat-treatment, the alloy must be cooled to form a supersaturated solid solution. Although the mode of cooling is not critical, rapidly cooling

the alloy to a temperature at which the diffusion rate of any of the elements in the alloy is not appreciable, and the formation of precipitates prevented, ensures the retention of the

equilibrium number of atomic vacancies (or as many of such vacancies as practicable) from

solution heat-treatment. This will assist in the diffusion and nucleation of precipitates during

the aging step of the present invention, which is discussed in detail below. Rapid cooling will

also serve to trap as much solute in solid solution as possible, making the maximum amount

of solute available for the subsequent formation of precipitates during aging. The rapid

cooling may be accomplished, for example, by quenching in a medium such as water, oil or

air, or any other known rapid cooling mechanism. The alloy forming the 1 " thick plates discussed in the example above was particularly sensitive to the speed of the cooling process. Accordingly, the plates were quenched using room temperature water.

Plastic Deformation In accordance with the method of the present invention, once solution heat- treatment is complete, the alloy must be sufficiently plastically deformed to produce high-

energy defect structures, such as the high-energy defect structures commonly referred to as

"deformation bands." Such high-energy defect structures may be exploited to promote a more uniform distribution of heterogeneously nucleated precipitate particles after aging than would otherwise be obtainable.

In contrast, attempts recently have been made to achieve such a favorable distribution of precipitates by imparting deformation to the material sufficient to induce a dislocation network, a lower-energy defect structure than contemplated by the present

invention. As discussed previously, Chung, et al. attempted to obtain such a dislocation

network by cold rolling, but with marginal results. In fact, the inventors hereof attempted to improve upon Chung, et al.'s efforts by stretching the subject material, since stretching would

be expected to impart a more uniform deformation across the thickness of the material. This

effort, too, was unsuccessful. Figure 5a shows precipitates that resulted from 8% stretching,

after the material had been heated at 380°C for 17 hours. Amounts of stretching from about

0% to 8% and heating times of about 2 to 17 hours resulted in precipitates having a similar

appearance to those shown in Figure 5a.

The inventors hereof have found that, instead of dislocation networks,

substantially higher-energy defects such as deformation bands must be formed. Deformation

bands provide nucleation sites at the interfaces of the bands which may be exploited to homogenize the precipitate distribution as needed for producing the fine-grained structure necessary for inducing superplasticity. Deformation bands are just one type of high-energy

defect structure that may be useful in the process of the present invention, however, and it is

not intended that the present invention be limited to the use of deformation bands. For example, other high-energy defect structures known as microbands, kink bands and bands of secondary slip may be used to equal effect.

Deformation bands or other high-energy defect structures useful under the present invention may be obtained by severely plastically deforming the solution heat-treated

alloy. Many processes for plastically deforming a material are known to those skilled in the art, such as rolling, stretching, extrusion, drawing, forging, and torsion processes, among

others. It is anticipated that any mode of plastic deformation may be used, so long as it is sufficiently severe to produce the required high-energy defect structure in the grains of the material. Preferably, the amount of reduction per pass and number of passes is such that the

deformation fully penetrates the alloy. It is also preferable that the deformation be uniform

throughout the thickness of the alloy.

The deformation of the solution heat-treated alloy preferably is carried out at

room temperature, although this temperature will vary with alloy composition, since some

alloying additions, such as magnesium in solid solution, are known to lower the dynamic

recovery rate. This step also may be carried out at other temperatures. Most preferably, the

deformation is performed at whatever temperature is most convenient and economical,

provided that sufficient energy is retained in the alloy for the formation of a high-energy

defect structure.

It is well-known that some alloying elements enhance the work hardening

behavior of alloys when such alloying elements are present in solid solution. For example, magnesium is known to have this effect in aluminum alloys, and makes possible the high strengths developed in wrought 5xxx alloys. Indeed, aluminum alloys containing Mg in solid

solution, such as the 6013/6111 alloy formed in accordance with the process of the present invention, may develop greater stored strain energy for a given amount of deformation than

alloys not containing Mg. Accordingly, the high-energy defect structures required for the

process of the present invention may be more readily attainable for alloys containing one or

more strength-enhancing alloying elements such as Mg than for alloys not having such alloying elements.

For the example of the 1" thick plate described above (standard preheat), unidirectional, room temperature rolling was carried out on 8.5" diameter rolls rotated at 11 rpm. The plate was reduced in thickness by about 10% per pass for a total of 9 passes. The microstructures of two such samples (rolling reduction of about 30% and 60%) were examined using SEM micrographs obtained using the electron channeling contrast technique

in a JEOL™ JSM-6400 scanning electron microscopy ("SEM") microscope, exhibited

banded deformation structures as shown in Figures 2a and 2b.

Following the aging step discussed below, a homogeneous precipitate distribution was observed. In addition, an unexpected and surprising effect of the severe

deformation step also was observed. Each of the samples subjected to aging after being

plastically deformed in accordance with the present invention exhibited precipitates that were

globular or near-spheroid in shape, as can be seen in Figures 3a, 3b and 3c. Such

morphologies are believed to be preferable over precipitates having other shapes, such as the

thin, square, plate-like morphologies that form in the absence of the severe deformation

disclosed herein, because spheroid or near-spheroid precipitates should be able to store strain

more uniformly. The formation of such globular precipitates is therefore believed to be a significant synergistic advance presented by the present invention.

Aging Once the alloy has been plastically deformed, it is aged to induce the nucleation and growth of precipitates. The preferred times and temperatures for the aging

process are dependent upon the type of alloy used, and are well known in the art (or may be obtained from the alloy manufacturer) for standard alloys. Where a unique alloy is being

processed with respect to which such times and temperatures have not been established, the

known times and temperatures for analogous alloys will provide a highly useful reference

point. As is well known, low aging temperatures require longer aging periods, whereas high

aging temperatures require shorter aging periods to achieve the same effect.

The aging process is preferably accomplished using more than one heating step, such that a relatively low temperature aging step may be used to form a fine distribution of precipitates, while one or more subsequent higher heating steps may be used to increase the speed of coarsening once precipitates have been formed in order to provide sufficiently coarse

particles to stimulate recrystallization. Beginning the aging process with a relatively lower

temperature increases the driving force for precipitation, thereby increasing the number density of precipitates, and continuing the aging process with a relatively higher temperature

decreases the aging time and enhances the economy of the process.

As will be appreciated, a single step aging process involving the use of a single

low or high-temperature aging step may also be used to form the desired distribution of

precipitates. As is explained in connection with the example discussed below, however, it is

possible that the preferred globular or near-spheroid precipitate morphology will not be

obtained where a single low-temperature aging step is used. Alternatively, the use of a single,

high-temperature step may be adequate to provide the preferred precipitate morphology, but may not provide as favorable a precipitate distribution. It has been found that by utilizing a low-temperature aging step followed by a high-temperature aging step, both the preferred morphology and distribution of precipitates may be realized.

Regardless of how many aging steps are used, the alloy may be cooled after

each aging step, preferably by air cooling. Air cooling should result in a larger volume

fraction of precipitates because the degree of supersaturation of the matrix is increased as the sample cools, while there is still enough thermal energy available for the diffusion of solute

atoms to the precipitate interfaces. Air cooling is also easier and less-costly to implement than other cooling methods such as quenching.

Exemplary samples of plastically deformed plates of the type discussed previously (identified below as samples A through E) were processed using single and dual precipitation heating steps, as shown in Table 1.

TABLE 1

EXEMPLARY AGING PROCESSES

The temperatures and heating times of the samples were varied in an attempt

to optimize the size, shape and distribution of the precipitates. With respect to the approximately 60%-rolled samples A and B, the presence of precipitates along parallel

deformation bands was apparent after only about one minute of heating at about 300 °C. For the approximately 60%-rolled samples, the precipitated zone was wider than for the

approximately 30%-rolled samples. After additional heating, precipitation between the deformation bands was visible, resulting in a fairly homogeneous distribution of precipitates

less than 1 μm in size.

The 60%-rolled samples A, B and C were analyzed further, and each of the

three samples exhibited a generally uniform distribution of globular precipitates about 1 -3 μm in diameter, as shown in Figures 3a, 3b and 3c, respectively. As noted previously, globular or

low aspect ratio precipitates are believed to be preferable over precipitates having other shapes, because spheroid or near-spheroid precipitates are able to store deformation more uniformly.

Sample D, which had not been subjected to a plastic deformation step

preceding the aging step, was also processed using an aging step in accordance with the

present invention. However, this sample exhibited a markedly less favorable precipitate

distribution and morphology when compared to those of the other samples. The result of the aging step on sample D is shown in the transmission electron microscopy ("TEM")

micrograph of Figure 4a. A similar TEM is provided with respect to sample A in Figure 4b,

which shows the complex precipitate structure present at the end of the process used to form

sample A. A comparison of Figures 4a and 4b illustrates that, with respect to the large

precipitates, a profound morphology change has resulted in sample A. The large particles

present in sample D are thin, square plates, while those present in sample A are finer and

more equiaxed with globular shapes, sometimes with facets. Further SEM analysis (not

shown) also revealed that the process used to form sample D, which did not include a pre- aging plastic deformation step, results in an extremely non-uniform distribution of the plate- shaped precipitates.

A sample that had been stretched by about 8% was subjected to an aging step at about 380°C for about 17 hours. The stretched sample exhibited large globular precipitates

and needle-like intragranular precipitates. It has been. shown that the grain boundary particles

coarsen while the intragranular particles resist coarsening. A comparison of Figure 5a

(stretched sample) and Figure 5b (cold-rolled sample A) shows that the distribution of

precipitates in sample A is extremely uniform compared to that produced in the stretched sample. It is believed that a dislocation network, instead of one of the desired higher-energy

defect structures, was produced in the stretched sample. Thus, plastic deformation such as

that applied to sample A by rolling is believed to be preferred over that applied by stretching, although stretching may still be an adequate mode of deformation where it is possible to impart sufficiently severe deformation to the material to produce a high-energy defect

structure without inducing fracture. The dimensional and distribution statistics for samples A, B, and C are shown

in Table 2.

TABLE 2

AGING STATISTICS

Where D_AVG = average particle diameter;

σ_D = standard deviation of particle diameters;

λ = mean free distance between particles; and

V_f = volume fraction of particles.

Based on the results shown in Table 2 and Figures 3a, 3b and 3c, process A

appeared to produce the best microstructural candidate for the PSN process. This was confirmed after a PSN process was applied to the material, as is discussed further below. It

will be appreciated, however, that sample B or C may be commercially preferable over

sample A despite their less ideal microstructures, in light of the fact that they require

significantly shorter time periods for aging than sample A.

It can be seen from these examples that a process utilizing a relatively low-

temperature aging step followed by a relatively high-temperature aging step (samples A and

B) provides a more uniform precipitate distribution than that utilizing a single, high-

temperature aging step (sample C). Specifically, although the precipitate distributions are

similar for samples A and B, the distribution resulting from process C consisted of a lower

number density of larger particles. This is probably due to the decreased driving force for

nucleation of precipitates for sample C as compared with samples A and B, since sample C (in contrast to samples A and B) was not processed using an initial low-temperature aging step.

It can also be seen that the use of a first, low-temperature heating step may not

result in the preferred globular precipitate moφhology. Specifically, after being subjected to such a heating step, sample A comprised only needle and rod/lath-shaped precipitates. The

globular-shaped precipitates appeared only during the second aging step.

Further analysis was performed to determine whether the globular precipitate

moφhology of sample A was the result of the plastic deformation imparted to the material

prior to aging. As part of this analysis, a sample was subjected to about 300°C for about 11 days, then about 380°C for about 29 days. Examination of this sample by TEM revealed that

the large precipitates still exhibited the same thin, square plate-like moφhologies seen in Figure 4a. No significant coarsening was observed, suggesting that the moφhology difference between micron-sized precipitates from sample A and this sample was not the

result of accelerated coarsening in sample A. Accordingly, it is believed that sample A exhibits generally globular precipitate moφhologies, whereas this sample exhibits plate-like

moφhologies, because sample A was subjected to pre-aging plastic deformation. The

specific reasons for the moφhology change are not known, although there is evidence that it

is due to different nucleation and growth conditions for the precipitates, or due to

simultaneous precipitation and/or phase change and recrystallization within the deformation

bands during aging.

Plastic Deformation

Once the aging step is completed, the alloy is subjected to a PSN process, the

general parameters of which are well known in the art. See, e.g.. U.S. Patent No. 4,092,181

to Paton, et al., which is incoφorated by reference herein in its entirety. The first step of this process is to plastically deform the material to form areas of strain, referred to as deformation

zones, around the precipitates. Each deformation zone provides favorable sites for nucleation

of recrystallized grains. As in the prior severe plastic deformation step, any mode of plastic deformation may be used, so long as it generally uniformly and completely penetrates the material. Also as in the severe plastic deformation step, the deformation of the present step

may be carried out at room temperature or at other lower or higher temperatures, but

preferably is performed at the temperature at or below the recrystallization temperature at

which the greatest amount deformation is stored around the particles.

The number of passes and the amount of deformation applied per pass will depend upon the alloy being worked, as well as the size of the precipitates. In any event, the

deformation stored in the alloy must be sufficient to ensure recrystallization through PSN. Preferably, it will be sufficient to produce fine grain sizes (preferably about 20 μm or less,

and most preferably about 10 μm or less) after recrystallization.

For the example of samples A-C described above, unidirectional, room

temperature rolling was carried out on 8.5" diameter rolls rotated at 11 φm. The plates were

reduced in thickness by a total of about 80% and 87% by applying 20% reductions. Sample E required a larger subsequent rolling reduction (about 92%) to attain the same final thickness

as samples A-C reduced about 87%. This produced excellent results, as discussed in detail

below in connection with the 87% reduction. It is contemplated that for some alloys, rolling

reductions even less than about 80% will produce sufficient deformation to yield satisfactory

results.

Sample A was further studied to optimize the effects of roll speed, reductions-

per-pass and total rolling reduction on the final grain size and shape. For the six

combinations of parameters obtainable from these three variables, average grain sizes (on LS sections at midthickness) ranged from about 9.5 to about 11.6 μm, with standard deviations

increasing with grain size from about 4.7 to about 5.7 μm. The finest grain size corresponded to the slower roll speed, higher total rolling reduction, and larger number of reductions-per-

pass is shown in the SEM micrograph of Figure 6. Its corresponding grain boundary map is

shown in Figure 7, which illustrates grain boundaries with greater than 10 ° misorientation.

Static Recrystallization The next step of the PSN process is to subject the alloy to a conventional static

recrystallization process to recrystallize to a fine grain structure. During the recrystallization

step, the highly strained regions of the deformation zones or other high-energy defect structures have a significant effect in encouraging nucleation of recrystallization. The

recrystallized grains grow to consume the deformation zones until the grains impinge on one another or until the drag force exerted on them by dispersoid particles balances the driving force for grain growth. Thus, important to controlling grain growth in this process is the use

of insoluble dispersoids present in the alloy.

As persons having skill in the art will recognize, the parameters of the

recrystallization process will depend upon the composition of the particular alloy being

processed and the amount of deformation stored in the material. Preferably, however, the

heat-up rate to the temperature at which recrystallization occurs is sufficiently rapid that no

recovery occurs in the deformation zones, which would effect a reduction in the driving force

for nucleation of recrystallization. Indeed, when PSN is exploited for grain-size control, an

increased heating rate during recrystallization has been shown to increase the number of

activated recrystallization nuclei. Thus, the heat-up rate preferably is as high as possible.

The heating time should only be as long as necessary to achieve complete recrystallization.

The temperature chosen for recrystallization must be equal to or greater than the critical recrystallization temperature for the material at which recrystallization occurs and

recovery is minimized. In one embodiment of the present invention, recrystallization occurs

during supeφlastic forming, in which case the temperature chosen for recrystallization is the

supeφlastic forming temperature. Regardless of the recrystallization temperature used, care must be taken to rapidly cool the alloy once recrystallization is complete. Accordingly, cold

water quenching or its equivalent is preferred.

For samples A, B, C and E, plastically deformed as described above, a

recrystallization temperature of about 540 °C was used, which is approximately the same

temperature as that used for solution heat treating. An air furnace was first fully preheated to

this temperature. The alloy samples were placed in the heated furnace and allowed to soak for about five minutes, after which they were quenched using room temperature water.

The recrystallized grain structures are characterized in Table 3, which contains

statistics related to average grain diameters and aspect ratios (measured on LS planes, at

midwidth and midthickness) for samples A, B, C and E.

TABLE 3

GRAIN STATISTICS FOR SAMPLES A-C

Where D_Δ average grain diameter;

^: standard deviation of grain diameters; and

AR = average grain aspect ratio;

σ₄ standard deviation of grain aspect ratios; and

Roundness = proximity to circular shape = (perimeters/area x 4π). The grain sizes, aspect ratios and size distributions represented in Table 3 were determined using quantitative image analysis of grain boundary maps generated from

microtexture data, which minimizes the influence of subgrain size on the average grain size.

Thus, as will be appreciated by persons skilled in the art, this technique provides a much

more rigorous and conservative evaluation of grain size statistics than do the optical microscopy techniques used in several of the background studies discussed previously in this

application. Indeed, unlike the technique employed in connection with the results presented

here, optical microscopy techniques do not permit one to easily distinguish between subgrain

boundaries and grain boundaries, making it virtually impossible to properly and accurately

limit grain sizes to areas bounded by high-angle grain boundaries.

The data presented in Table 3 shows a fine grain structure with average grain sizes of about 9.5 μm to about 11.6 μm. The grains are nearly equiaxed, having average

aspect ratios of about 1.6 to about 1.9. The size and aspect ratio distributions are narrow,

indicating a high degree of uniformity of this grain structure. Compared with commercial 6xxx aluminum alloys of similar composition (Al-Mg-Si or Al-Mg-Si-Cu), the average grain

sizes, aspect ratios, distributions and roundness of these samples are statistically superior.

It is apparent from both a qualitative comparison of Figures 8-11 and a quantitative comparison of the average grain diameters shown in Table 3 that the process

used to produce sample A yielded the finest, most equiaxed and uniform grain structure.

Table 4 shows the results of grain boundary map analysis taken from the LS, LT and ST

planes of the recrystallized sample A produced using the optimized downstream rolling and recrystallization conditions previously discussed. As illustrated by Table 4, the result of the

present process is a fine (average grain size of about 10.3 μm over the LS planes), equiaxed grain structure. In addition, the average three-dimensional grain size increased only to about

10.7 μm after a one hour exposure to the same temperature, demonstrating that the grain size

is statically stable, a critical property if the material is to be useful as a supeφlastic alloy. It is believed that in the alloy of sample A, manganese-bearing dispersoid particles are responsible

for preventing further grain growth.

TABLE 4

GRAIN SIZES FOR ALLOY SAMPLE A

Where D_AVG = average grain diameter; and

σ_D = standard deviation of grain diameters.

Orientation Distribution Function and microtexture analyses indicate a very

weak texture on both LT and LS planes. A sample A alloy made from the ingot subjected to a low-temperature preheat,

as described previously, also was 87% cold rolled at room temperature and statically recrystallized. The resultant grains were finer but less equiaxed than the grains of the sample

A described in Table 4, the statistics for which were derived from the ingot subjected to the

conventional preheat. Thus, it was concluded that ingot subjected to the standard preheat is

preferable to ingot subjected to the low-temperature preheat for processing using the method

of the present invention.

Superplastic Results of the Present Invention Figure 12 illustrates the variation of strain rate sensitivity with strain rate for

uniaxial, step strain rate tests at 500°C and 540°C, performed on the version of sample A

produced using the optimized downstream rolling and recrystallization conditions previously discussed. The material exhibited a maximum strain rate sensitivity of 0.5, which occurred at

540 °C for a strain rate range of 2x10^"4 s^"1 to 5X10^" s^"! (based on initial gage length). Figure

13 shows the elongation as a function of strain rate for a temperature of 540 °C. The elongation to fracture reached 375% with a corresponding maximum stress of approximately

680 psi (4.7 MPa). Figure 14 shows an undeformed sample alongside samples deformed to

350 to 375%. Such supeφlastic elongation results are superior to any results previously reported for non-eutectic 6xxx aluminum alloys. Indeed, the marginal supeφlastic results of

Chung, et al. for a 6013 aluminum alloy, as discussed previously, yielded only 230% elongation at 520°C at a strain rate of 3 x 10^-4 s^"1 and a flow stress of 972 psi (6.7 MPa).

Chung, et al. also obtained a maximum strain rate sensitivity of only 0.38. It also may be

noted for comparison that a baseline, commercially available 6013 -T4 sheet tested under the same conditions as sample A fractured after about 120% elongation with a maximum stress

of approximately 860 psi (5.9 MPa).

Accordingly, the results of the process of the present invention, as exemplified

by sample A, illustrate that the distribution of precipitates in an alloy may be significantly

homogenized by creating and exploiting deformation bands or other high-energy defect

structures as heterogeneous nucleation sites for precipitation. This approach, preferably

coupled with a multi-step low and high temperature aging process, produces the uniform

distribution of micron-size precipitates necessary for the subsequent development of a fine,

equiaxed grain structure following PSN that is stable at supeφlastic forming temperatures.

For many alloys, superior supeφlastic properties may result.

In particular, the grain structure characteristics, static stability and supeφlastic

properties of this supeφlastic alloy are exceptional. Indeed, the 6013/6111 alloy produced

using the preferred process of the present invention is markedly superior to those reported previously for other 6xxx aluminum alloys claiming supeφlastic properties. Given its superior characteristics, and the relatively energy efficient and rapid process by which it is

produced, this alloy is potentially useful for many commercial applications, including many

conceivable applications in the aerospace and automotive industries. In addition, the process of the present invention is expected to be similarly useful for many other alloys, including

aluminum 6061, 6063 and 6066 alloys, as well as many other age-hardenable aluminum

alloys, and including magnesium, iron, titanium, nickel and other alloy systems.

It is believed that the many advantages of the present invention will now be apparent to those skilled in the art. It will also be apparent that a number of variations and

modifications may be made thereto without departing from its spirit and scope. Accordingly,

the foregoing description is to be construed as illustrative only, rather than limiting. The

present invention is limited only by the scope of the following claims.

Claims

What is claimed is:

1. A method for producing a supeφlastic alloy, said method comprising:

providing an alloy for processing, said alloy comprising a matrix phase and at least

two alloying elements, at least one of said alloying elements being, or being capable of forming, a dispersoid phase substantially insoluble in said matrix phase; solution heat treating said alloy;

cooling the alloy to form a supersaturated solid solution; plastically deforming said alloy in a first deformation step sufficiently to form a high- energy defect structure, thereby forming nucleation sites useful for the subsequent nucleation of precipitates;

aging said alloy, thereby forming precipitates at said nucleation sites; and plastically deforming said alloy in a second deformation step, and statically recrystallizing said alloy, through a particle-stimulated nucleation process.

2. The method of claim 1 , wherein said step of providing an alloy for processing

comprises providing an aluminum alloy.

3. The method of claim 2, wherein said aluminum alloy is selected from the group

consisting of aluminum alloys 6013, 6111, 6061 , 6063 , and 6066.

4. The method of claim 1 , wherein said cooling step comprises quenching.

5. The method of claim 1 , wherein said first deformation step comprises plastically

deforming said alloy sufficiently to form deformation bands.

6. The method of claim 1 , wherein said first deformation step comprises cold rolling said alloy.

7. The method of claim 6, wherein said first deformation step further comprises cold rolling said alloy at room temperature.

8. The method of claim 7, wherein said first deformation step further comprises cold rolling said alloy to a reduction of at least about 30%.

9. The method of claim 1 , wherein said aging step comprises a first heating step at a first

temperature and a second heating step at a second higher temperature.

10. The method of claim 9, wherein said precipitates are formed during said first heating

step and coarsened during said second heating step.

11. The method of claim 9, wherein said alloy is cooled after said first heating step and

after said second heating step.

12. The method of claim 1 , wherein said second deformation step comprises cold rolling

said alloy.

13. The method of claim 12, wherein said second deformation step further comprises cold

rolling said alloy at room temperature.

14. The method of claim 1, wherein said static recrystallization step comprises rapidly

heating said alloy to a temperature at which recrystallization occurs.

15. The method of claim 1 , wherein said static recrystallization step comprises heating said alloy to a temperature in the range of a solution heat-treatment temperature for said alloy.

16. The method of claim 1 , wherein said static recrystallization step comprises heating said alloy to a supeφlastic forming temperature of said alloy.

17. A method for producing a supeφlastic aluminum alloy, said method comprising:

providing an alloy for processing, said alloy being a 6013/6111 alloy;

solution heat treating said alloy; cooling the alloy to form a supersaturated solid solution; plastically deforming said alloy in a first deformation step sufficiently to form a high- energy defect structure, thereby forming nucleation sites useful for the subsequent nucleation

of precipitates;

aging said alloy, thereby forming precipitates at said nucleation sites;

plastically deforming said alloy in a second deformation step to provide sufficient

strain energy in said alloy to ensure recrystallization; and

statically recrystallizing said alloy.

18. The method of claim 17, wherein said 6013/6111 alloy has the approximate

composition 97.3 wt % Al - 0.8 wt % Mg - 0.7 wt % Si - 0.8 wt % Cu - 0.3 wt % Mn - 0.1 wt % Fe.

19. The method of claim 17, wherein said solution heat treating step is performed at a temperature of about 540 ┬░C for about one hour.

20. The method of claim 17, wherein said cooling step comprises quenching.

21. The method of claim 17, wherein said first deformation step comprises cold rolling said alloy.

22. The method of claim 21, wherein said first deformation step comprises cold rolling said alloy to a reduction of at least about 30%.

23. The method of claim 22, wherein said first deformation step comprises cold rolling said alloy to a reduction of at least about 60%.

24. The method of claim 17, wherein said first deformation step comprises plastically

deforming said alloy sufficiently to form deformation bands.

25. The method of claim 17, wherein said first deformation step is performed such that,

after subsequent aging, said alloy exhibits globular or near-spheroid shaped precipitates.

26. The method of claim 17, wherein said aging step comprises a first heating step at a

first temperature and a second heating step at a second higher temperature.

27. The method of claim 26, wherein said precipitates are formed during said first heating step and coarsened during said second heating step.

28. The method of claim 26, wherein said alloy is cooled after said first heating step and after said second heating step.

29. The method of claim 26, wherein said first heating step is performed at about 300┬░C

and said second heating step is performed at about 380┬░C.

30. The method of claim 29, wherein the duration of said first heating step is about 24

hours, and the duration of said second heating step is about 24 hours.

31. The method of claim 26, wherein said first heating step is performed at about 300┬░C

and said second heating step is performed at about 450┬░C.

32. The method of claim 31 , wherein the duration of said first heating step is about 24

hours, and the duration of said second heating step is about 2 hours.

33. The method of claim 17, wherein said aging step comprises heating said alloy at a

temperature of about 450┬░C for about 2 hours.

34. The method of claim 17, wherein said second deformation step comprises cold rolling

said alloy.

35. The method of claim 34, wherein said second deformation step comprises cold rolling said alloy to a reduction of at least about 80%.

36. The method of claim 35, wherein said second deformation step comprises cold rolling said alloy to a reduction of at least about 87%.

37. The method of claim 36, wherein said second deformation step comprises cold rolling

said alloy to a reduction of at least about 92%.

38. The method of claim 17, wherein said static recrystallization step comprises rapidly

heating said alloy to a temperature at which recrystallization occurs.

39. The method of claim 38, wherein said static recrystallization step comprises heating

said alloy to a temperature of about 540 ┬░C for about 5 minutes.

40. A 6xxx aluminum alloy having a microstructure comprising grains having an average

size in the range of about 9.5 ╬╝m to about 11.6 ╬╝m, said grain sizes having a standard

deviation in the range of about 4.7 ╬╝m to about 5.6 ╬╝m, and said grains further having an

average grain aspect ratio in the range of about 1.6 to about 1.9, said grain aspect ratios

having a standard deviation in the range of about 0.6 to about 0.8.

41. The 6xxx aluminum alloy of claim 40, wherein said alloy has a grain roundness in the

range of about 1.6 to about 1.8.

42. The 6xxx aluminum alloy of claim 40, wherein said alloy comprises grains having an

average size of about 9.5 ╬╝m, said grain sizes having a standard deviation of about 4.7 ╬╝m,

and said grains further having an average grain aspect ratio of about 1.6, said grain aspect

ratios having a standard deviation of about 0.6.

43. The 6xxx aluminum alloy of claim 40, wherein said alloy exhibits a maximum elongation of at least about 350%.

44. The 6xxx aluminum alloy of claim 40, wherein said alloy exhibits a maximum strain rate sensitivity of at least about 0.5.

45. The 6xxx aluminum alloy of claim 44, wherein said alloy exhibits a maximum

elongation of at least about 375%.

46. A supeφlastic 6xxx aluminum alloy having a maximum strain rate sensitivity of at

least about 0.5 and exhibiting a maximum elongation of at least about 350%, and being

produced according to a process comprising: providing an alloy for processing, said alloy comprising a matrix phase and at least

two alloying elements, at least one of said alloying elements being, or being capable of

forming, a dispersoid phase substantially insoluble in said matrix phase;

solution heat-treating said alloy;

cooling the alloy to form a supersaturated solid solution;

plastically deforming said alloy in a first deformation step sufficiently to form a high- energy defect structure, thereby forming nucleation sites useful for the subsequent nucleation of precipitates;

aging said alloy, thereby forming precipitates at said nucleation sites; and plastically deforming said alloy in a second deformation step, and statically

recrystallizing said alloy, through a particle-stimulated nucleation process.