U.S. patent application number 09/814687 was filed with the patent office on 2001-09-27 for method of producing superplastic alloys and superplastic alloys produced by the method.
Invention is credited to Crooks, Roy, Starke, Edgar A. JR., Troeger, Lillianne P..
Application Number | 20010023719 09/814687 |
Document ID | / |
Family ID | 26780389 |
Filed Date | 2001-09-27 |
United States Patent
Application |
20010023719 |
Kind Code |
A1 |
Troeger, Lillianne P. ; et
al. |
September 27, 2001 |
Method of producing superplastic alloys and superplastic alloys
produced by the method
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 sufficiently 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 supersaturated 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.
Inventors: |
Troeger, Lillianne P.;
(Norfolk, VA) ; Starke, Edgar A. JR.;
(Charlottesville, VA) ; Crooks, Roy; (Newport
News, VA) |
Correspondence
Address: |
FOLEY & LARDNER
Washington Harbour
3000 K Street, N.W., Suite 500
Washington
DC
20007-5109
US
|
Family ID: |
26780389 |
Appl. No.: |
09/814687 |
Filed: |
March 15, 2001 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
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09814687 |
Mar 15, 2001 |
|
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|
09332736 |
Jun 14, 1999 |
|
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60089239 |
Jun 15, 1998 |
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Current U.S.
Class: |
148/415 |
Current CPC
Class: |
C22C 21/08 20130101;
C22F 1/05 20130101; C22F 1/047 20130101 |
Class at
Publication: |
148/415 |
International
Class: |
C22C 021/00 |
Goverment Interests
[0002] 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.
Claims
What is claimed is:
1. A method for producing a superplastic 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 superplastic forming
temperature of said alloy.
17. A method for producing a superplastic 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.degree. 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.degree. C. and said second heating step is
performed at about 380.degree. 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.degree. C. and said second heating step is
performed at about 450.degree. 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.degree. 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.degree. 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 .mu.m to about
11.6 .mu.m, said grain sizes having a standard deviation in the
range of about 4.7 .mu.m to about 5.6 .mu.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 .mu.m, said
grain sizes having a standard deviation of about 4.7 .mu.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 superplastic 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.
Description
CROSS-REFERENCE To RELATED PROVISIONAL APPLICATION
[0001] The present application claims the benefit of the earlier
filing date of U.S. Provisional Patent Application Ser. No.
60/089,239, filed Jun. 15, 1998, which is incorporated by reference
herein in its entirety.
FIELD OF THE INVENTION
[0003] 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
[0004] The advantages of superplastic properties in metals are well
known, and particularly well employed in the automotive and
aerospace industry. Because of their fine-grained 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.sub.m, where
T.sub.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.
[0005] 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.
[0006] Superplastic behavior in metallic alloys may be described by
the equation
.delta.=k{dot over (.epsilon.)}.sup.m
[0007] where .delta.=flow stress, k=material constant, {dot over
(.epsilon.)}=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.
[0008] 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 .mu.m or less in diameter and preferably about
10 .mu.m 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.
[0009] An extensive amount of research has been conducted in an
effort to discover thermomechanical processes useful for producing
superplastic alloys, including aluminum alloys.
[0010] 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.
[0011] 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 6 xxx 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.
[0012] To date, efforts expended to impart superplasticity to 6xxx
aluminum alloys have not been very successful. U.S. Pat. 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.
[0013] 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 .mu.m (in the rolling plane), as measured using
optical microscopy ("OM") techniques. They obtained a maximum
elongation of 148% at 450.degree. C., due to significant grain
growth occurring at 500.degree. C. and above, within the
superplastic forming temperature range. The Washfold, et al.
process did not achieve superplasticity.
[0014] 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.degree. C. and
570.degree. C. and strain rates of 10.sup.-3 to 10.sup.-6 s.sup.-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.
[0015] 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.sup.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.degree. C., 90% warm rolling at 190.degree. 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 Mg.sub.2Si precipitates.
The process resulted in grains of 12 to 13 .mu.m (measured using
optical microscopy techniques), a strain rate sensitivity of 0.38,
and a maximum elongation of 230% at 520.degree. C. for a strain
rate of 3.times.10.sup.-4 s.sup.-l, 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.
[0016] A similar process to that employed by Chung, et al., but
directed to an altogether different purpose, is described in U.S.
Pat. 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.
[0017] Accordingly, it is an object of the present invention to
provide alloys exhibiting superplasticity, particularly 6xxx alloys
and especially aluminum 6013 and 6111 alloys.
[0018] 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.
[0019] 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.
[0020] 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 .mu.m, and preferably about 10 .mu.m or less.
[0021] 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
[0022] 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.
[0023] 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.
[0024] 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.
[0025] 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.
[0026] 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.
[0027] 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.
[0028] 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.degree. 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.degree. C. for about 24
hours and a second heating step is performed at about 380.degree.
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.
[0029] 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.degree. C. for about 24
hours, and a second heating step at about 450.degree. 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.degree. 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.
[0030] 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.degree. C. and
held there for about five minutes.
[0031] 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 .mu.m to about 11.6 .mu.m, the
grain sizes having a standard deviation in the range of about 4.7
.mu.m to about 5.6 .mu.m. 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.degree. C. for about 24 hours and a second
heating step at about 380.degree. 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.degree. C. for about five minutes, yields an
average grain size of about 9.5 .mu.m (about 4.7 .mu.m 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.degree. C. for a strain rate
range of 2.times.10.sup.-4 s.sup.-1 to 5.times.10.sup.-4 s.sup.-1,
and a maximum elongation of about 375% with a corresponding maximum
stress of approximately 680 psi (4.7 MPa).
[0032] 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
[0033] FIG. 1 is a SEM micrograph (150.times.) of a 6013/6111
alloy, produced in accordance with the method of the present
invention, following solution heat treatment.
[0034] FIG. 2a is a SEM micrograph (500.times.) illustrating banded
deformation structures produced in 30% cold rolled sample E in
accordance with the method of the present invention;
[0035] FIG. 2b is a SEM micrograph (500.times.) illustrating banded
deformation structures produced in 60% cold rolled sample A in
accordance with the method of the present invention;
[0036] FIG. 3a is a SEM micrograph (5000.times.) illustrating
globular or near-spheroid shaped precipitates as produced in sample
A in accordance with the method of the present invention;
[0037] FIG. 3b is a SEM micrograph (5000.times.) illustrating
globular or near-spheroid shaped precipitates as produced in sample
B in accordance with the method of the present invention;
[0038] FIG. 3c is a SEM micrograph (5000.times.) illustrating
globular or near-spheroid shaped precipitates as produced in sample
C in accordance with the method of the present invention;
[0039] FIG. 4a is a TEM micrograph (3300.times.) of sample D after
the aging step of the present invention;
[0040] FIG. 4b is a TEM micrograph (3300.times.) of sample A after
the aging step of the present invention;
[0041] FIG. 5a is a SEM micrograph (1000.times.) illustrating the
distribution of precipitates in an 8% stretched 6013/6111 sample
heated at 380.degree. C. for 17 hours;
[0042] FIG. 5b is a SEM micrograph (500.times.) illustrating the
distribution of precipitates in sample A, produced in accordance
with the method of the present invention;
[0043] FIG. 6 is a SEM micrograph (200 .mu.m width.times.150 .mu.m
height) illustrating the grain size of sample A processed using
optimized downstream processing conditions in accordance with the
method of the present invention;
[0044] FIG. 7 is a 10.degree. misorientation grain boundary map
(200 .mu.m width.times.150 .mu.m height) corresponding to the SEM
micrograph of FIG. 6;
[0045] FIG. 8a is a SEM micrograph (150.times.) illustrating the
recrystallized grain structure of sample A produced in accordance
with the method of the present invention;
[0046] FIG. 8b is a SEM micrograph (500.times.) illustrating the
recrystallized grain structure of sample A produced in accordance
with the method of the present invention;
[0047] FIG. 9a is a SEM micrograph (150.times.) illustrating the
recrystallized grain structure of sample B produced in accordance
with the method of the present invention;
[0048] FIG. 9b is a SEM micrograph (500.times.) illustrating the
recrystallized grain structure of sample B produced in accordance
with the method of the present invention;
[0049] FIG. 10a is a SEM micrograph (150.times.) illustrating the
recrystallized grain structure of sample C produced in accordance
with the method of the present invention;
[0050] FIG. 10b is a SEM micrograph (500.times.) illustrating the
recrystallized grain structure of sample C produced in accordance
with the method of the present invention;
[0051] FIG. 11a is a SEM micrograph (150.times.) illustrating the
recrystallized grain structure of sample E produced in accordance
with the method of the present invention;
[0052] FIG. 11b is a SEM micrograph (500.times.) illustrating the
recrystallized grain structure of sample E produced in accordance
with the method of the present invention;
[0053] 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.degree. C. and 540.degree. C.;
[0054] 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.degree.
C.; and
[0055] 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
[0056] 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
[0057] 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.
[0058] 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, Va. facility. One
half of the ingot was preheated in the conventional manner using a
heat-up rate of about 50.degree. C./hour, a soak temperature of
about 560.degree. C., and a soak time of about four hours. The
other half of the ingot underwent a low-temperature preheat (about
500.degree. C.) using a heat-up rate of about 50.degree. 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.
[0059] 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
[0060] 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.degree. C. The samples were placed in the
furnace for a period of about one hour, excluding heat-up time. A
SEM micrograph (150.times.) of a sample of this material following
solution heat-treatment is shown in FIG. 1.
Rapid Cooling
[0061] 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.
[0062] 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
[0063] 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.
[0064] 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. FIG. 5a shows precipitates that
resulted from 8% stretching, after the material had been heated at
380.degree. 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 FIG.
5a.
[0065] 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.
[0066] 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.
[0067] 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.
[0068] 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.
[0069] 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 JEOLTM JSM-6400
scanning electron microscopy ("SEM") microscope, exhibited banded
deformation structures as shown in FIGS. 2a and 2b.
[0070] 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 FIGS. 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
[0071] 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.
[0072] 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.
[0073] 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.
[0074] 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.
[0075] 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.
1TABLE 1 EXEMPLARY AGING PROCESSES Time at Time at Time at %
300.degree. C. 380.degree. C. 450.degree. C. Samples Cold Rolling
(hours) (hours) (hours) A 60 24 24 0 B 60 24 0 2 C 60 0 0 2 D 0 24
24 0 E 30 24 24 0
[0076] 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.degree. 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 .mu.m in
size.
[0077] 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 .mu.m in diameter,
as shown in FIGS. 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.
[0078] 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 FIG. 4a. A
similar TEM is provided with respect to sample A in FIG. 4b, which
shows the complex precipitate structure present at the end of the
process used to form sample A. A comparison of FIGS. 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.
[0079] A sample that had been stretched by about 8% was subjected
to an aging step at about 380.degree. 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 Sa (stretched sample) and
FIG. 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.
[0080] The dimensional and distribution statistics for samples A,
B, and C are shown in Table 2.
2TABLE 2 AGING STATISTICS Samples D.sub.AVG (.mu.m) .sigma..sub.D
(.mu.m) .lambda. (.mu.m) V.sub.f (%) A 0.70 0.38 10.50 6.00 B 0.66
0.30 11.63 5.62 C 0.92 0.42 13.70 5.30 Where D.sub.AVG = average
particle diameter; .sigma..sub.D = standard deviation of particle
diameters; .lambda. = mean free distance between particles; and
V.sub.f = volume fraction of particles.
[0081] Based on the results shown in Table 2 and FIGS. 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.
[0082] 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.
[0083] It can also be seen that the use of a first, low-temperature
heating step may not result in the preferred globular precipitate
morphology. 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.
[0084] Further analysis was performed to determine whether the
globular precipitate morphology 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.degree.
C. for about 11 days, then about 380.degree. C. for about 29 days.
Examination of this sample by TEM revealed that the large
precipitates still exhibited the same thin, square plate-like
morphologies seen in FIG. 4a. No significant coarsening was
observed, suggesting that the morphology 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
morphologies, whereas this sample exhibits plate-like morphologies,
because sample A was subjected to pre-aging plastic deformation.
The specific reasons for the morphology 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.
[0085] 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. Pat. No. 4,092,181
to Paton, et al., which is incorporated 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.
[0086] 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 .mu.m or less, and most preferably about 10
.mu.m or less) after recrystallization.
[0087] For the example of samples A-C described above,
unidirectional, room temperature rolling was carried out on 8.5"
diameter rolls rotated at 11 rpm. 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.
[0088] 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
.mu.m, with standard deviations increasing with grain size from
about 4.7 to about 5.7 .mu.m. 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
FIG. 6. Its corresponding grain boundary map is shown in FIG. 7,
which illustrates grain boundaries with greater than 10.degree.
misorientation.
Static Recrystallization
[0089] 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.
[0090] 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.
[0091] 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 superplastic forming, in which case
the temperature chosen for recrystallization is the superplastic
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.
[0092] For samples A, B, C and E, plastically deformed as described
above, a recrystallization temperature of about 540.degree. 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.
[0093] 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.
3TABLE 3 GRAIN STATISTICS FOR SAMPLES A-C Sample D.sub.AVG (.mu.m)
.sigma..sub.D (.mu.m) AR .lambda..sub.AR Roundness A 9.48 4.72 1.65
0.55 1.64 B 11.63 5.58 1.85 0.80 1.81 C 10.80 5.63 1.89 0.67 1.80 E
10.82 4.76 1.63 0.58 1.74 Where D.sub.AVG = average grain diameter;
.sigma..sub.D = standard deviation of grain diameters; and AR =
average grain aspect ratio; .sigma..sub.AR = standard deviation of
grain aspect ratios; and Roundness = proximity to circular shape =
(perimeter.sup.2/area .times. 4.sub..pi.).
[0094] 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.
[0095] The data presented in Table 3 shows a fine grain structure
with average grain sizes of about 9.5 .mu.m to about 11.6 .mu.m.
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.
[0096] It is apparent from both a qualitative comparison of FIGS.
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 .mu.m
over the LS planes), equiaxed grain structure. In addition, the
average three-dimensional grain size increased only to about 10.7
.mu.m 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 superplastic alloy.
It is believed that in the alloy of sample A, manganese-bearing
dispersoid particles are responsible for preventing further grain
growth.
4TABLE 4 GRAIN SIZES FOR ALLOY SAMPLE A Plane Soak Time (min)
D.sub.AVG (.mu.m) .sigma..sub.D (.mu.m) LS 5 9.48 4.47 LT 5 10.35
5.50 ST 5 10.90 4.97 LS 60 10.97 4.72 LT 60 10.47 4.54 ST 60 10.72
4.87 Where D.sub.AVG = average grain diameter; and .sigma..sub.D =
standard deviation of grain diameters.
[0097] Orientation Distribution Function and microtexture analyses
indicate a very weak texture on both LT and LS planes.
[0098] 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
[0099] FIG. 12 illustrates the variation of strain rate sensitivity
with strain rate for uniaxial, step strain rate tests at
500.degree. C. and 540.degree. 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.degree. C. for a strain rate range of 2.times.10.sup.-4
s.sup.-1 to 5.times.10.sup.-4 s.sup.-1 (based on initial gage
length). FIG. 13 shows the elongation as a function of strain rate
for a temperature of 540.degree. C. The elongation to fracture
reached 375% with a corresponding maximum stress of approximately
680 psi (4.7 MPa). FIG. 14 shows an undeformed sample alongside
samples deformed to 350 to 375%. Such superplastic elongation
results are superior to any results previously reported for
non-eutectic 6xxx aluminum alloys. Indeed, the marginal
superplastic results of Chung, et al. for a 6013 aluminum alloy, as
discussed previously, yielded only 230% elongation at 520.degree.
C. at a strain rate of 3.times.10.sup.-4 s.sup.-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).
[0100] 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 superplastic forming temperatures.
For many alloys, superior superplastic properties may result.
[0101] In particular, the grain structure characteristics, static
stability and superplastic properties of this superplastic 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
superplastic 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.
[0102] 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.
* * * * *