U.S. patent application number 12/398712 was filed with the patent office on 2010-09-09 for high strength l12 aluminum alloys produced by cryomilling.
This patent application is currently assigned to UNITED TECHNOLOGIES CORPORATION. Invention is credited to Awadh B. Pandey.
Application Number | 20100226817 12/398712 |
Document ID | / |
Family ID | 42678419 |
Filed Date | 2010-09-09 |
United States Patent
Application |
20100226817 |
Kind Code |
A1 |
Pandey; Awadh B. |
September 9, 2010 |
HIGH STRENGTH L12 ALUMINUM ALLOYS PRODUCED BY CRYOMILLING
Abstract
A method and apparatus produces high strength aluminum alloys
from a cryomilled powder containing L1.sub.2 intermetallic
dispersoids. The cryomilled powder is degassed, sealed under vacuum
in a container, heated, consolidated by vacuum hot pressing, and
extruded.
Inventors: |
Pandey; Awadh B.; (Jupiter,
FL) |
Correspondence
Address: |
KINNEY & LANGE, P.A.
THE KINNEY & LANGE BUILDING, 312 SOUTH THIRD STREET
MINNEAPOLIS
MN
55415-1002
US
|
Assignee: |
UNITED TECHNOLOGIES
CORPORATION
Hartford
CT
|
Family ID: |
42678419 |
Appl. No.: |
12/398712 |
Filed: |
March 5, 2009 |
Current U.S.
Class: |
420/532 ;
419/23 |
Current CPC
Class: |
B22F 2999/00 20130101;
B22F 2999/00 20130101; C22C 21/06 20130101; C22C 21/12 20130101;
B22F 9/04 20130101; C22C 1/0416 20130101; C22C 21/02 20130101; B22F
2009/041 20130101; B22F 9/04 20130101; B22F 3/14 20130101; B22F
3/1216 20130101; B22F 3/1216 20130101; B22F 3/20 20130101; B22F
9/082 20130101; B22F 2201/20 20130101; B22F 2202/03 20130101; B22F
2998/10 20130101; C22C 21/003 20130101; B22F 2003/145 20130101;
B22F 2999/00 20130101; B22F 9/04 20130101; C22C 21/10 20130101;
B22F 2998/10 20130101 |
Class at
Publication: |
420/532 ;
419/23 |
International
Class: |
C22C 1/04 20060101
C22C001/04; C22C 21/02 20060101 C22C021/02 |
Claims
1. A method for producing a high strength aluminum alloy billet
containing L1.sub.2 dispersoids, comprising the steps of:
cryomilling a quantity of an aluminum alloy powder containing an
L1.sub.2 dispersoid therein to produce a mesh size of less than
-325 mesh; placing the cryomilled powder in a container; vacuum
degassing the powder at a temperature of about 500.degree. F.
(260.degree. C.) to about 900.degree. F. (482.degree. C.) for about
12 hours to about 8 days; sealing the degassed powder in the
container under vacuum; heating the sealed container at about
700.degree. F. (371.degree. C.) to about 900.degree. F.
(482.degree. C.) for about one to eight hours; vacuum hot pressing
the heated container to form a billet; and removing the container
from the formed billet.
2. The method of claim 1, wherein the container is aluminum having
a central axis, and vacuum hot pressing is done along the axis
while restraining radial movement of the container.
3. The method of claim 1, wherein the vacuum hot pressing includes
blind die compaction for about 1 hour to about 8 hours at a
temperature of 700.degree. F. (371.degree. C.) to about 900.degree.
F. (482.degree. C.) under uni-axial pressure of about 20 ksi (138
Mpa) to about 100 ksi (690 MPa).
4. The method of claim 1, wherein the vacuum hot pressing produces
a billet of the aluminum alloy powder having a theoretical density
of about 100 percent.
5. The method of claim 1, wherein the degassing includes rotating
the aluminum alloy powder to heat and expose all the powder to
vacuum.
6. The method of claim 1, wherein the thus formed billet is
extruded at a pressure of about 20 ksi (138 Mpa) to about 100 ksi
(690 MPa).
7. The method of claim 6, wherein the extrusion temperature is
about 300.degree. F. (149.degree. C.) to about 850.degree. F.
(454.degree. C.), the billet soak time is about 0.5 hours to about
8 hours at a rate of about 0.2 inch per minute (0.51 cm per minute)
to about 10 inch per minute (25.4 cm per minute), and an extrusion
ratio of about 2:1 to about 40:1.
8. The method of claim 1, wherein the L1.sub.2 dispersoids comprise
Al.sub.3X dispersoids wherein X is at least one first element
selected from the group comprising: about 0.1 to about 4.0 weight
percent scandium, about 0.1 to about 20.0 weight percent erbium,
about 0.1 to about 15.0 weight percent thulium, about 0.1 to about
25.0 weight percent ytterbium, and about 0.1 to about 25.0 weight
percent lutetium; at least one second element selected from the
group comprising about 0.1 to about 20.0 weight percent gadolinium,
about 0.1 to about 20.0 weight percent yttrium, about 0.05 to about
4.0 weight percent zirconium, about 0.05 to about 10.0 weight
percent titanium, about 0.05 to about 10.0 weight percent hafnium,
and about 0.05 to about 5.0 weight percent niobium; and the balance
substantially aluminum.
9. The method of claim 8, wherein the aluminum alloy powder
contains at least one third element selected from the group
consisting of silicon, magnesium, lithium, copper, zinc, and
nickel.
10. The method of claim 9, wherein the third element comprises at
least one of about 4 to about 25 weight percent silicon, about 1 to
about 8 weight percent magnesium, about 0.5 to about 3 weight
percent lithium, about 0.2 to about 3 weight percent copper, about
3 to about 12 weight percent zinc, about 1 to about 12 weight
percent nickel.
11. The method of claim 8, wherein the L1.sub.2 dispersoids
comprise Al.sub.3X dispersoids wherein X is at least one first
element selected from the group comprising: about 0.1 to about 0.5
weight percent scandium, about 0.1 to about 6.0 weight percent
erbium, about 0.1 to about 10.0 weight percent thulium, about 0.1
to about 15.0 weight percent ytterbium, and about 0.1 to about 12.0
weight percent lutetium; at least one second element selected from
the group comprising about 0.1 to about 4.0 weight percent
gadolinium, about 0.1 to about 4.0 weight percent yttrium, about
0.05 to about 1.0 weight percent zirconium, about 0.05 to about 2.0
weight percent titanium, about 0.05 to about 2.0 weight percent
hafnium, and about 0.05 to about 1.0 weight percent niobium; and
the balance substantially aluminum.
12. A method for producing a high strength aluminum alloy billet
containing L1.sub.2 dispersoids, comprising the steps of: sieving a
quantity of an aluminum alloy powder containing an L1.sub.2
dispersoid therein to produce a particle size of less than 100
mesh; blending the sieved aluminum alloy powder to homogenize the
particle size distribution; cryomilling the blended aluminum alloy
powder containing an L1.sub.2 dispersoid therein to produce a mesh
size of less than -325 mesh; and compacting the cryomilled powder
to form a billet.
13. The method of claim 12, wherein the cryomilled powder is placed
in a container, vacuum degassed at a temperature of about
500.degree. F. (260.degree. C.) to about 900.degree. F.
(482.degree. C.) for about 12 hours to about 8 days; followed by
sealing the degassed powder in the container under vacuum and
heating the sealed container at about 700.degree. F. (371.degree.
C.) to about 900.degree. F. (482.degree. C.) for about one to eight
hours.
14. The method of claim 13, which further includes the step of
vacuum hot pressing the heated container to form a billet; and
removing the container from the formed billet.
15. The method of claim 14, wherein the vacuum hot pressing
produces a billet of the aluminum alloy powder having a density of
about 100 percent.
16. The method of claim 13, wherein the degassing includes rotating
the aluminum alloy powder to heat and exposing to vacuum all the
powder.
17. The method of claim 12, wherein the formed billet is extruded
at a pressure of about 20 ksi (138 MPA) to about 100 ksi (690
MPa).
18. The method of claim 17, wherein the extrusion temperature is
about 300.degree. F. (149.degree. C.) to about 850.degree. F.
(454.degree. C.), the billet soak time is about 0.5 hours to about
8 hours at a rate of about 0.2 inch per minute (0.51 cm per minute)
to about 10 inch per minute (25.4 cm per minute), and an extrusion
ratio of about 2:1 to about 40:1.
19. The method of claim 12, wherein the L1.sub.2 dispersoids
comprise Al.sub.3X dispersoids wherein X is at least one first
element selected from the group comprising: about 0.1 to about 4.0
weight percent scandium, about 0.1 to about 20.0 weight percent
erbium, about 0.1 to about 15.0 weight percent thulium, about 0.1
to about 25.0 weight percent ytterbium, and about 0.1 to about 25.0
weight percent lutetium; at least one second element selected from
the group comprising about 0.1 to about 20.0 weight percent
gadolinium, about 0.1 to about 20.0 weight percent yttrium, about
0.05 to about 4.0 weight percent zirconium, about 0.05 to about
10.0 weight percent titanium, about 0.05 to about 10.0 weight
percent hafnium, and about 0.05 to about 5.0 weight percent
niobium; and the balance substantially aluminum.
20. The method of claim 19, wherein the third element comprises at
least one of about 4 to about 25 weight percent silicon, about 1 to
about 8 weight percent magnesium, about 0.5 to about 3 weight
percent lithium, about 0.2 to about 3 weight percent copper, about
3 to about 12 weight percent zinc, about 1 to about 12 weight
percent nickel.
21. A high strength aluminum alloy billet comprising: aluminum
alloy matrix; and dispersoids within the aluminum matrix wherein X
is at least one first element selected from the group consisting
of: about 0.1 to about 4.0 weight percent scandium, about 0.1 to
about 20.0 weight percent erbium, about 0.1 to about 15.0 weight
percent thulium, about 0.1 to about 25.0 weight percent ytterbium,
and about 0.1 to about 25.0 weight percent lutetium; at least one
second element selected from the group consisting of about 0.1 to
about 20.0 weight percent gadolinium, about 0.1 to about 20.0
weight percent yttrium, about 0.05 to about 4.0 weight percent
zirconium, about 0.05 to about 10.0 weight percent titanium, about
0.05 to about 10.0 weight percent hafnium, and about 0.05 to about
5.0 weight percent niobium; and at least one selected from the
group consisting of about 4 to about 25 weight percent silicon,
about 1 to about 8 weight percent magnesium, about 0.5 to about 3
weight percent lithium, about 0.2 to about 3 weight percent copper,
about 3 to about 12 weight percent zinc, about 1 to about 12 weight
percent nickel; wherein the billet dorsile strength of at least 100
ksi (690 MPa) and a yield strength of at least 95 ksi (655 MPa).
Description
CROSS-REFERENCE TO RELATED APPLICATION(S)
[0001] This application is related to the following co-pending
applications that were filed on Dec. 9, 2008 herewith and are
assigned to the same assignee: CONVERSION PROCESS FOR HEAT
TREATABLE L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/316,020; A METHOD
FOR FORMING HIGH STRENGTH ALUMINUM ALLOYS CONTAINING L1.sub.2
INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,046; and A METHOD FOR
PRODUCING HIGH STRENGTH ALUMINUM ALLOY POWDER CONTAINING L1.sub.2
INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,047.
[0002] This application is also related to the following co-pending
applications that were filed on Apr. 18, 2008, and are assigned to
the same assignee: L1.sub.2 ALUMINUM ALLOYS WITH BIMODAL AND
TRIMODAL DISTRIBUTION, Ser. No. 12/148,395; DISPERSION STRENGTHENED
L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/148,432; HEAT TREATABLE
L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/148,383; HIGH STRENGTH
L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/148,394; HIGH STRENGTH
L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/148,382; HEAT TREATABLE
L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/148,396; HIGH STRENGTH
L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/148,387; HIGH STRENGTH
ALUMINUM ALLOYS WITH L1.sub.2 PRECIPITATES, Ser. No. 12/148,426;
HIGH STRENGTH L1.sub.2 ALUMINUM ALLOYS, Ser. No. 12/148,459; and
L1.sub.2 STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, Ser. No.
12/148,458.
BACKGROUND
[0003] The present invention relates generally to aluminum alloys
and more specifically to a method for forming high strength
aluminum alloy powder having L1.sub.2 dispersoids therein.
[0004] The combination of high strength, ductility, and fracture
toughness, as well as low density, make aluminum alloys natural
candidates for aerospace and space applications. However, their use
is typically limited to temperatures below about 300.degree. F.
(149.degree. C.) since most aluminum alloys start to lose strength
in that temperature range as a result of coarsening of
strengthening precipitates.
[0005] The development of aluminum alloys with improved elevated
temperature mechanical properties is a continuing process. Some
attempts have included aluminum-iron and aluminum-chromium based
alloys such as Al--Fe--Ce, Al--Fe--V--Si, Al--Fe--Ce--W, and
Al--Cr--Zr--Mn that contain incoherent dispersoids. These alloys,
however, also lose strength at elevated temperatures due to
particle coarsening. In addition, these alloys exhibit ductility
and fracture toughness values lower than other commercially
available aluminum alloys.
[0006] Other attempts have included the development of mechanically
alloyed Al--Mg and Al--Ti alloys containing ceramic dispersoids.
These alloys exhibit improved high temperature strength due to the
particle dispersion, but the ductility and fracture toughness are
not improved.
[0007] U.S. Pat. No. 6,248,453 owned by the assignee of the present
invention discloses aluminum alloys strengthened by dispersed
Al.sub.3X L1.sub.2 intermetallic phases where X is selected from
the group consisting of Sc, Er, Lu, Yb, Tm, and Lu. The Al.sub.3X
particles are coherent with the aluminum alloy matrix and are
resistant to coarsening at elevated temperatures. The improved
mechanical properties of the disclosed dispersion strengthened
L1.sub.2 aluminum alloys are stable up to 572.degree. F.
(300.degree. C.). U.S. Patent Application Publication No.
2006/0269437 A1, also commonly owned, discloses a high strength
aluminum alloy that contains scandium and other elements that is
strengthened by L1.sub.2 dispersoids.
[0008] L1.sub.2 strengthened aluminum alloys have high strength and
improved fatigue properties compared to commercially available
aluminum alloys. Fine grain size results in improved mechanical
properties of materials. Hall-Petch strengthening has been known
for decades where strength increases as grain size decreases. An
optimum grain size for optimum strength is in the nano range of
about 30 to 100 nm. These alloys also have lower ductility.
SUMMARY
[0009] The present invention is a method for consolidating
cryomilled aluminum alloy powders into useful components with high
temperature strength and fracture toughness. In embodiments,
powders include an aluminum alloy having coherent L1.sub.2
Al.sub.3X dispersoids where X is at least one first element
selected from scandium, erbium, thulium, ytterbium, and lutetium,
and at least one second element selected from gadolinium, yttrium,
zirconium, titanium, hafnium, and niobium. The balance is
substantially aluminum containing at least one alloying element
selected from silicon, magnesium, lithium, copper, zinc, and
nickel.
[0010] The powders are classified by sieving and blended to improve
homogeneity. Cryomilling is an essential step in the manufacturing
process. The cryomilled powders are then vacuum degassed in a
container that is then sealed. The sealed container (i.e. can) is
compressed by vacuum hot pressing, hot isostatic pressing or blind
die compaction to densify the powder charge. The can is removed and
the billet is extruded, forged and/or rolled into useful shapes
with high temperature strength and fracture toughness.
BRIEF DESCRIPTION OF THE DRAWINGS
[0011] FIG. 1 is an aluminum scandium phase diagram.
[0012] FIG. 2 is an aluminum erbium phase diagram.
[0013] FIG. 3 is an aluminum thulium phase diagram.
[0014] FIG. 4 is an aluminum ytterbium phase diagram.
[0015] FIG. 5 is an aluminum lutetium phase diagram.
[0016] FIG. 6A and 6B are SEM photos of the gas atomized inventive
L1.sub.2 aluminum alloy powder.
[0017] FIG. 7A and 7B are photomicrographs of cross-sections
showing the cellular microstructure of the gas atomized inventive
L1.sub.2 aluminum alloy powder.
[0018] FIG. 8A and 8B are photomicrographs of cryomilled powder of
the inventive L1.sub.2 aluminum alloy powder.
[0019] FIG. 9A and 9B are photomicrographs of cross-sections of
cryomilled powder of the inventive L1.sub.2 aluminum alloy
powder.
[0020] FIG. 10 is a diagram showing the processing steps to
consolidate L1.sub.2 aluminum alloy powder.
[0021] FIG. 11 is a photo of a 3-inch diameter copper jacketed
L1.sub.2 aluminum alloy billet.
[0022] FIG. 12 is a photo of extrusion dies for 3-inch diameter
billet.
[0023] FIG. 13 is a photo of extruded L1.sub.2 aluminum alloy rods
from 3-inch diameter billets.
[0024] FIG. 14 is a photo of machined L1.sub.2 aluminum alloy
billets.
[0025] FIG. 15 is a photo of a machined three-piece L1.sub.2
aluminum alloy billet assembly for 6-inch copper jacketed extrusion
billet.
[0026] FIG. 16 is a photo of extruded L1.sub.2 aluminum alloy rods
from 6-inch diameter billets.
DETAILED DESCRIPTION
1. L1.sub.2 Aluminum Alloys
[0027] Alloy powders refined by this invention are formed from
aluminum based alloys with high strength and fracture toughness for
applications at temperatures from about -420.degree. F.
(-251.degree. C.) up to about 650.degree. F. (343.degree. C.). The
aluminum alloy comprises a solid solution of aluminum and at least
one element selected from silicon, magnesium, lithium, copper,
zinc, and nickel strengthened by L1.sub.2 Al.sub.3X coherent
precipitates where X is at least one first element selected from
scandium, erbium, thulium, ytterbium, and lutetium, and at least
one second element selected from gadolinium, yttrium, zirconium,
titanium, hafnium, and niobium.
[0028] The aluminum silicon system is a simple eutectic alloy
system with a eutectic reaction at 12.5 weight percent silicon and
1077.degree. F. (577.degree. C.). There is little solubility of
silicon in aluminum at temperatures up to 930.degree. F.
(500.degree. C.) and none of aluminum in silicon. However, the
solubility can be extended significantly by utilizing rapid
solidification techniques
[0029] The binary aluminum magnesium system is a simple eutectic at
36 weight percent magnesium and 842.degree. F. (450.degree. C.).
There is complete solubility of magnesium and aluminum in the
rapidly solidified inventive alloys discussed herein
[0030] The binary aluminum lithium system is a simple eutectic at 8
weight percent lithium and 1105.degree. (596.degree. C.). The
equilibrium solubility of 4 weight percent lithium can be extended
significantly by rapid solidification techniques. There can be
complete solubility of lithium in the rapid solidified inventive
alloys discussed herein.
[0031] The binary aluminum copper system is a simple eutectic at 32
weight percent copper and 1018.degree. F. (548.degree. C.). There
can be complete solubility of copper in the rapidly solidified
inventive alloys discussed herein.
[0032] The aluminum zinc binary system is a eutectic alloy system
involving a monotectoid reaction and a miscibility gap in the solid
state. There is a eutectic reaction at 94 weight percent zinc and
718.degree. F. (381.degree. C.). Zinc has maximum solid solubility
of 83.1 weight percent in aluminum at 717.8.degree. F. (381.degree.
C.), which can be extended by rapid solidification processes.
Decomposition of the super saturated solid solution of zinc in
aluminum gives rise to spherical and ellipsoidal GP zones, which
are coherent with the matrix and act to strengthen the alloy.
[0033] The aluminum nickel binary system is a simple eutectic at
5.7 weight percent nickel and 1183.8.degree. F. (639.9.degree. C.).
There is little solubility of nickel in aluminum. However, the
solubility can be extended significantly by utilizing rapid
solidification processes. The equilibrium phase in the aluminum
nickel eutectic system is L1.sub.2 intermetallic Al.sub.3Ni.
[0034] In the aluminum based alloys disclosed herein, scandium,
erbium, thulium, ytterbium, and lutetium are potent strengtheners
that have low diffusivity and low solubility in aluminum. All these
elements form equilibrium Al.sub.3X intermetallic dispersoids where
X is at least one of scandium, erbium, thulium, ytterbium, and
lutetium, that have an L1.sub.2 structure that is an ordered face
centered cubic structure with the X atoms located at the corners
and aluminum atoms located on the cube faces of the unit cell.
[0035] Scandium forms Al.sub.3Sc dispersoids that are fine and
coherent with the aluminum matrix. Lattice parameters of aluminum
and Al.sub.3Sc are very close (0.405 nm and 0.410 nm respectively),
indicating that there is minimal or no driving force for causing
growth of the Al.sub.3Sc dispersoids. This low interfacial energy
makes the Al.sub.3Sc dispersoids thermally stable and resistant to
coarsening up to temperatures as high as about 842.degree. F.
(450.degree. C.). Additions of magnesium in aluminum increase the
lattice parameter of the aluminum matrix, and decrease the lattice
parameter mismatch further increasing the resistance of the
Al.sub.3Sc to coarsening. Additions of zinc, copper, lithium,
silicon, and nickel provide solid solution and precipitation
strengthening in the aluminum alloys. These Al.sub.3Sc dispersoids
are made stronger and more resistant to coarsening at elevated
temperatures by adding suitable alloying elements such as
gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or
combinations thereof, that enter Al.sub.3Sc in solution.
[0036] Erbium forms Al.sub.3Er dispersoids in the aluminum matrix
that are fine and coherent with the aluminum matrix. The lattice
parameters of aluminum and Al.sub.3Er are close (0.405 nm and 0.417
nm respectively), indicating there is minimal driving force for
causing growth of the Al.sub.3Er dispersoids. This low interfacial
energy makes the Al.sub.3Er dispersoids thermally stable and
resistant to coarsening up to temperatures as high as about
842.degree. F. (450.degree. C.). Additions of magnesium in aluminum
increase the lattice parameter of the aluminum matrix, and decrease
the lattice parameter mismatch further increasing the resistance of
the Al.sub.3Er to coarsening. Additions of zinc, copper, lithium,
silicon, and nickel provide solid solution and precipitation
strengthening in the aluminum alloys. These Al.sub.3Er dispersoids
are made stronger and more resistant to coarsening at elevated
temperatures by adding suitable alloying elements such as
gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or
combinations thereof that enter Al.sub.3Er in solution.
[0037] Thulium forms metastable Al.sub.3Tm dispersoids in the
aluminum matrix that are fine and coherent with the aluminum
matrix. The lattice parameters of aluminum and Al.sub.3Tm are close
(0.405 nm and 0.420 nm respectively), indicating there is minimal
driving force for causing growth of the Al.sub.3Tm dispersoids.
This low interfacial energy makes the Al.sub.3Tm dispersoids
thermally stable and resistant to coarsening up to temperatures as
high as about 842.degree. F. (450.degree. C.). Additions of
magnesium in aluminum increase the lattice parameter of the
aluminum matrix, and decrease the lattice parameter mismatch
further increasing the resistance of the Al.sub.3Tm to coarsening.
Additions of zinc, copper, lithium, silicon, and nickel provide
solid solution and precipitation strengthening in the aluminum
alloys. These Al.sub.3Tm dispersoids are made stronger and more
resistant to coarsening at elevated temperatures by adding suitable
alloying elements such as gadolinium, yttrium, zirconium, titanium,
hafnium, niobium, or combinations thereof that enter Al.sub.3Tm in
solution.
[0038] Ytterbium forms Al.sub.3Yb dispersoids in the aluminum
matrix that are fine and coherent with the aluminum matrix. The
lattice parameters of Al and Al.sub.3Yb are close (0.405 nm and
0.420 nm respectively), indicating there is minimal driving force
for causing growth of the Al.sub.3Yb dispersoids. This low
interfacial energy makes the Al.sub.3Yb dispersoids thermally
stable and resistant to coarsening up to temperatures as high as
about 842.degree. F. (450.degree. C.). Additions of magnesium in
aluminum increase the lattice parameter of the aluminum matrix, and
decrease the lattice parameter mismatch further increasing the
resistance of the Al.sub.3Yb to coarsening. Additions of zinc,
copper, lithium, silicon, and nickel provide solid solution and
precipitation strengthening in the aluminum alloys. These
Al.sub.3Yb dispersoids are made stronger and more resistant to
coarsening at elevated temperatures by adding suitable alloying
elements such as gadolinium, yttrium, zirconium, titanium, hafnium,
niobium, or combinations thereof that enter Al.sub.3Yb in
solution.
[0039] Lutetium forms Al.sub.3Lu dispersoids in the aluminum matrix
that are fine and coherent with the aluminum matrix. The lattice
parameters of Al and Al.sub.3Lu are close (0.405 nm and 0.419 nm
respectively), indicating there is minimal driving force for
causing growth of the Al.sub.3Lu dispersoids. This low interfacial
energy makes the Al.sub.3Lu dispersoids thermally stable and
resistant to coarsening up to temperatures as high as about
842.degree. F. (450.degree. C.). Additions of magnesium in aluminum
increase the lattice parameter of the aluminum matrix, and decrease
the lattice parameter mismatch further increasing the resistance of
the Al.sub.3Lu to coarsening. Additions of zinc, copper, lithium,
silicon, and nickel provide solid solution and precipitation
strengthening in the aluminum alloys. These Al.sub.3Lu dispersoids
are made stronger and more resistant to coarsening at elevated
temperatures by adding suitable alloying elements such as
gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or
mixtures thereof that enter Al.sub.3Lu in solution.
[0040] Gadolinium forms metastable Al.sub.3Gd dispersoids in the
aluminum matrix that are stable up to temperatures as high as about
842.degree. F. (450.degree. C.) due to their low diffusivity in
aluminum. The Al.sub.3Gd dispersoids have a D0.sub.19 structure in
the equilibrium condition. Despite its large atomic size,
gadolinium has fairly high solubility in the Al.sub.3X
intermetallic dispersoids (where X is scandium, erbium, thulium,
ytterbium or lutetium). Gadolinium can substitute for the X atoms
in Al.sub.3X intermetallic, thereby forming an ordered L1.sub.2
phase which results in improved thermal and structural
stability.
[0041] Yttrium forms metastable Al.sub.3Y dispersoids in the
aluminum matrix that have an L1.sub.2 structure in the metastable
condition and a D0.sub.19 structure in the equilibrium condition.
The metastable Al.sub.3Y dispersoids have a low diffusion
coefficient, which makes them thermally stable and highly resistant
to coarsening. Yttrium has a high solubility in the Al.sub.3X
intermetallic dispersoids allowing large amounts of yttrium to
substitute for X in the Al.sub.3X L1.sub.2 dispersoids, which
results in improved thermal and structural stability.
[0042] Zirconium forms Al.sub.3Zr dispersoids in the aluminum
matrix that have an L1.sub.2 structure in the metastable condition
and D0.sub.23 structure in the equilibrium condition. The
metastable Al.sub.3Zr dispersoids have a low diffusion coefficient,
which makes them thermally stable and highly resistant to
coarsening. Zirconium has a high solubility in the Al.sub.3X
dispersoids allowing large amounts of zirconium to substitute for X
in the Al.sub.3X dispersoids, which results in improved thermal and
structural stability.
[0043] Titanium forms Al.sub.3Ti dispersoids in the aluminum matrix
that have an L1.sub.2 structure in the metastable condition and
DO.sub.22 structure in the equilibrium condition. The metastable
Al.sub.3Ti dispersoids have a low diffusion coefficient, which
makes them thermally stable and highly resistant to coarsening.
Titanium has a high solubility in the Al.sub.3X dispersoids
allowing large amounts of titanium to substitute for X in the
Al.sub.3X dispersoids, which result in improved thermal and
structural stability.
[0044] Hafnium forms metastable Al.sub.3Hf dispersoids in the
aluminum matrix that have an L1.sub.2 structure in the metastable
condition and a D0.sub.23 structure in the equilibrium condition.
The Al.sub.3Hf dispersoids have a low diffusion coefficient, which
makes them thermally stable and highly resistant to coarsening.
Hafnium has a high solubility in the Al.sub.3X dispersoids allowing
large amounts of hafnium to substitute for scandium, erbium,
thulium, ytterbium, and lutetium in the above-mentioned Al.sub.3X
dispersoids, which results in stronger and more thermally stable
dispersoids.
[0045] Niobium forms metastable Al.sub.3Nb dispersoids in the
aluminum matrix that have an L1.sub.2 structure in the metastable
condition and a D0.sub.22 structure in the equilibrium condition.
Niobium has a lower solubility in the Al.sub.3X dispersoids than
hafnium or yttrium, allowing relatively lower amounts of niobium
than hafnium or yttrium to substitute for X in the Al.sub.3X
dispersoids. Nonetheless, niobium can be very effective in slowing
down the coarsening kinetics of the Al.sub.3X dispersoids because
the Al.sub.3Nb dispersoids are thermally stable. The substitution
of niobium for X in the above mentioned Al.sub.3X dispersoids
results in stronger and more thermally stable dispersoids.
[0046] Al.sub.3X L1.sub.2 precipitates improve elevated temperature
mechanical properties in aluminum alloys for two reasons. First,
the precipitates are ordered intermetallic compounds. As a result,
when the particles are sheared by glide dislocations during
deformation, the dislocations separate into two partial
dislocations separated by an anti-phase boundary on the glide
plane. The energy to create the anti-phase boundary is the origin
of the strengthening. Second, the cubic L1.sub.2 crystal structure
and lattice parameter of the precipitates are closely matched to
the aluminum solid solution matrix. This results in a lattice
coherency at the precipitate/matrix boundary that resists
coarsening. The lack of an interphase boundary results in a low
driving force for particle growth and resulting elevated
temperature stability. Alloying elements in solid solution in the
dispersed strengthening particles and in the aluminum matrix that
tend to decrease the lattice mismatch between the matrix and
particles will tend to increase the strengthening and elevated
temperature stability of the alloy.
[0047] L1.sub.2 phase strengthened aluminum alloys are important
structural materials because of their excellent mechanical
properties and the stability of these properties at elevated
temperature due to the resistance of the coherent dispersoids in
the microstructure to particle coarsening. The mechanical
properties are optimized by maintaining a high volume fraction of
L1.sub.2 dispersoids in the microstructure. The L1.sub.2 dispersoid
concentration following aging scales as the amount of L1.sub.2
phase forming elements in solid solution in the aluminum alloy
following quenching. Examples of L1.sub.2 phase forming elements
include but are not limited to Sc, Er, Th, Yb, and Lu. The
concentration of alloying elements in solid solution in alloys
cooled from the melt is directly proportional to the cooling
rate.
[0048] Exemplary aluminum alloys for the bimodal system alloys of
this invention include, but are not limited to (in weight percent
unless otherwise specified):
[0049] about Al-M-(0.1-4)Sc-(0.1-20)Gd;
[0050] about Al-M-(0.1-20)Er-(0.1-20)Gd;
[0051] about Al-M-(0.1-15)Tm-(0.1-20)Gd;
[0052] about Al-M-(0.1-25)Yb-(0.1-20)Gd;
[0053] about Al-M-(0.1-25)Lu-(0.1-20)Gd;
[0054] about Al-M-(0.1-4)Sc-(0.1-20)Y;
[0055] about Al-M-(0.1-20)Er-(0.1-20)Y;
[0056] about Al-M-(0.1-15)Tm-(0.1-20)Y;
[0057] about Al-M-(0.1-25)Yb-(0.1-20)Y;
[0058] about Al-M-(0.1-25)Lu-(0.1-20)Y;
[0059] about Al-M-(0.1-4)Sc-(0.05-4)Zr;
[0060] about Al-M-(0.1-20)Er-(0.05-4)Zr;
[0061] about Al-M-(0.1-15)Tm-(0.05-4)Zr;
[0062] about Al-M-(0.1-25)Yb-(0.05-4)Zr;
[0063] about Al-M-(0.1-25)Lu-(0.05-4)Zr;
[0064] about Al-M-(0.1-4)Sc-(0.05-10)Ti;
[0065] about Al-M-(0.1-20)Er-(0.05-10)Ti;
[0066] about Al-M-(0.1-15)Tm-(0.05-10)Ti;
[0067] about Al-M-(0.1-25)Yb-(0.05-10)Ti;
[0068] about Al-M-(0.1-25)Lu-(0.05-10)Ti;
[0069] about Al-M-(0.1-4)Sc-(0.05-10)Hf;
[0070] about Al-M-(0.1-20)Er-(0.05-10)Hf;
[0071] about Al-M-(0.1-15)Tm-(0.05-10)Hf;
[0072] about Al-M-(0.1-25)Yb-(0.05-10)Hf;
[0073] about Al-M-(0.1-25)Lu-(0.05-10)Hf;
[0074] about Al-M-(0.1-4)Sc-(0.05-5)Nb;
[0075] about Al-M-(0.1-20)Er-(0.05-5)Nb;
[0076] about Al-M-(0.1-15)Tm-(0.05-5)Nb;
[0077] about Al-M-(0.1-25)Yb-(0.05-5)Nb; and
[0078] about Al-M-(0.1-25)Lu-(0.05-5)Nb.
[0079] M is at least one of about (4-25) weight percent silicon,
(1-8) weight percent magnesium, (0.5-3) weight percent lithium,
(0.2-3) weight percent copper, (3-12) weight percent zinc, and
(1-12) weight percent nickel.
[0080] The amount of silicon present in the fine grain matrix, if
any, may vary from about 4 to about 25 weight percent, more
preferably from about 4 to about 18 weight percent, and even more
preferably from about 5 to about 11 weight percent.
[0081] The amount of magnesium present in the fine grain matrix, if
any, may vary from about 1 to about 8 weight percent, more
preferably from about 3 to about 7.5 weight percent, and even more
preferably from about 4 to about 6.5 weight percent.
[0082] The amount of lithium present in the fine grain matrix, if
any, may vary from about 0.5 to about 3 weight percent, more
preferably from about 1 to about 2.5 weight percent, and even more
preferably from about 1 to about 2 weight percent.
[0083] The amount of copper present in the fine grain matrix, if
any, may vary from about 0.2 to about 3 weight percent, more
preferably from about 0.5 to about 2.5 weight percent, and even
more preferably from about 1 to about 2.5 weight percent.
[0084] The amount of zinc present in the fine grain matrix, if any,
may vary from about 3 to about 12 weight percent, more preferably
from about 4 to about 10 weight percent, and even more preferably
from about 5 to about 9 weight percent.
[0085] The amount of nickel present in the fine grain matrix, if
any, may vary from about 1 to about 12 weight percent, more
preferably from about 2 to about 10 weight percent, and even more
preferably from about 4 to about 10 weight percent.
[0086] The amount of scandium present in the fine grain matrix, if
any, may vary from 0.1 to about 4 weight percent, more preferably
from about 0.1 to about 3 weight percent, and even more preferably
from about 0.2 to about 2.5 weight percent. The Al--Sc phase
diagram shown in FIG. 1 indicates a eutectic reaction at about 0.5
weight percent scandium at about 1219.degree. F. (659.degree. C.)
resulting in a solid solution of scandium and aluminum and
Al.sub.3Sc dispersoids. Aluminum alloys with less than 0.5 weight
percent scandium can be quenched from the melt to retain scandium
in solid solution that may precipitate as dispersed L1.sub.2
intermetallic Al.sub.3Sc following an aging treatment. Alloys with
scandium in excess of the eutectic composition (hypereutectic
alloys) can only retain scandium in solid solution by rapid
solidification processing (RSP) where cooling rates are in excess
of about 10.sup.3.degree. C./second.
[0087] The amount of erbium present in the fine grain matrix, if
any, may vary from about 0.1 to about 20 weight percent, more
preferably from about 0.3 to about 15 weight percent, and even more
preferably from about 0.5 to about 10 weight percent. The Al--Er
phase diagram shown in FIG. 2 indicates a eutectic reaction at
about 6 weight percent erbium at about 1211.degree. F. (655.degree.
C.). Aluminum alloys with less than about 6 weight percent erbium
can be quenched from the melt to retain erbium in solid solutions
that may precipitate as dispersed L1.sub.2 intermetallic Al.sub.3Er
following an aging treatment. Alloys with erbium in excess of the
eutectic composition can only retain erbium in solid solution by
rapid solidification processing (RSP) where cooling rates are in
excess of about 10.sup.3.degree. C./second.
[0088] The amount of thulium present in the alloys, if any, may
vary from about 0.1 to about 15 weight percent, more preferably
from about 0.2 to about 10 weight percent, and even more preferably
from about 0.4 to about 6 weight percent. The Al--Tm phase diagram
shown in FIG. 3 indicates a eutectic reaction at about 10 weight
percent thulium at about 1193.degree. F. (645.degree. C.). Thulium
forms metastable Al.sub.3Tm dispersoids in the aluminum matrix that
have an L1.sub.2 structure in the equilibrium condition. The
Al.sub.3Tm dispersoids have a low diffusion coefficient, which
makes them thermally stable and highly resistant to coarsening.
Aluminum alloys with less than 10 weight percent thulium can be
quenched from the melt to retain thulium in solid solution that may
precipitate as dispersed metastable L1.sub.2 intermetallic
Al.sub.3Tm following an aging treatment. Alloys with thulium in
excess of the eutectic composition can only retain Tm in solid
solution by rapid solidification processing (RSP) where cooling
rates are in excess of about 10.sup.3.degree. C./second.
[0089] The amount of ytterbium present in the alloys, if any, may
vary from about 0.1 to about 25 weight percent, more preferably
from about 0.3 to about 20 weight percent, and even more preferably
from about 0.4 to about 10 weight percent. The Al--Yb phase diagram
shown in FIG. 4 indicates a eutectic reaction at about 21 weight
percent ytterbium at about 1157.degree. F. (625.degree. C.).
Aluminum alloys with less than about 21 weight percent ytterbium
can be quenched from the melt to retain ytterbium in solid solution
that may precipitate as dispersed L1.sub.2 intermetallic Al.sub.3Yb
following an aging treatment. Alloys with ytterbium in excess of
the eutectic composition can only retain ytterbium in solid
solution by rapid solidification processing (RSP) where cooling
rates are in excess of about 10.sup.3.degree. C./second.
[0090] The amount of lutetium present in the alloys, if any, may
vary from about 0.1 to about 25 weight percent, more preferably
from about 0.3 to about 20 weight percent, and even more preferably
from about 0.4 to about 10 weight percent. The Al--Lu phase diagram
shown in FIG. 5 indicates a eutectic reaction at about 11.7 weight
percent Lu at about 1202.degree. F. (650.degree. C.). Aluminum
alloys with less than about 11.7 weight percent lutetium can be
quenched from the melt to retain Lu in solid solution that may
precipitate as dispersed L1.sub.2 intermetallic Al.sub.3Lu
following an aging treatment. Alloys with Lu in excess of the
eutectic composition can only retain Lu in solid solution by rapid
solidification processing (RSP) where cooling rates are in excess
of about 10.sup.3.degree. C./second.
[0091] The amount of gadolinium present in the alloys, if any, may
vary from about 0.1 to about 20 weight percent, more preferably
from about 0.3 to about 15 weight percent, and even more preferably
from about 0.5 to about 10 weight percent.
[0092] The amount of yttrium present in the alloys, if any, may
vary from about 0.1 to about 20 weight percent, more preferably
from about 0.3 to about 15 weight percent, and even more preferably
from about 0.5 to about 10 weight percent.
[0093] The amount of zirconium present in the alloys, if any, may
vary from about 0.05 to about 4 weight percent, more preferably
from about 0.1 to about 3 weight percent, and even more preferably
from about 0.3 to about 2 weight percent.
[0094] The amount of titanium present in the alloys, if any, may
vary from about 0.05 to about 10 weight percent, more preferably
from about 0.2 to about 8 weight percent, and even more preferably
from about 0.4 to about 4 weight percent.
[0095] The amount of hafnium present in the alloys, if any, may
vary from about 0.05 to about 10 weight percent, more preferably
from about 0.2 to about 8 weight percent, and even more preferably
from about 0.4 to about 5 weight percent.
[0096] The amount of niobium present in the alloys, if any, may
vary from about 0.05 to about 5 weight percent, more preferably
from about 0.1 to about 3 weight percent, and even more preferably
from about 0.2 to about 2 weight percent.
[0097] In order to have the best properties for the fine grain
matrix, it is desirable to limit the amount of other elements.
Specific elements that should be reduced or eliminated include no
more than about 0.1 weight percent iron, 0.1 weight percent
chromium, 0.1 weight percent manganese, 0.1 weight percent
vanadium, and 0.1 weight percent cobalt. The total quantity of
additional elements should not exceed about 1% by weight, including
the above listed impurities and other elements.
2. Consolidation of Aluminum L1.sub.2 Alloy Powder
[0098] Gas atomized high temperature L1.sub.2 aluminum alloy powder
needs to be consolidated into solid-state forms suitable for
engineering applications. Scanning electron micrographs of gas
atomized L1.sub.2 aluminum alloy powder are shown in FIGS. 6A and
6B. The powder is spherical and capable of high packing density. As
a result of the high solidification rate, e.g. greater than
10.sup.3.degree. C./second, the microstructure is a finely divided
cellular structure instead of a dendritic structure common to
conventionally cooled alloys. SEM photos illustrating the fine
cellular structure of the L1.sub.2 aluminum powder are shown in
FIGS. 7A and 7B. The fine structure allows for a uniform
distribution of alloying elements and resulting even dispersion of
L1.sub.2 strengthening dispersoids in the final consolidated alloy
structure. FIGS. 8A and 8B show photomicrographs of cryomilled
powders at different magnifications indicating that the shape of
the spherical powder has changed from spherical to irregular due to
the milling operation. The microstructure of cryomilled powder
shown in FIGS. 9A and 9B indicates that the cellular structure is
refined and there is a more uniform distribution of fine particles
present in the powder. The process of consolidating the alloy
powders into useful forms is schematically illustrated in FIG.
10.
[0099] L1.sub.2 aluminum alloy powders 10 are first classified
according to size by sieving (step 20). Fine particle sizes are
required for optimum mechanical properties in the final part. The
starting stock should be at least -325 mesh powder (step 30). Other
benefits of blending will be discussed later. Powders are then
cryomilled to decrease grain size and improve strength (step 40).
Cryomilling is carried out in a high-energy ball mill under liquid
nitrogen, and offers several benefits that will be discussed
later.
[0100] The sieved, blended and cryomilled powders are then put in a
can (step 50) and vacuum degassed (step 60). Following vacuum
degassing the can is sealed (step 70) under vacuum and hot pressed
(step 80) to densify the powder compact. The compact is then hot
worked (step 90) to refine the microstructure. Finally, the
densified compact is machined into an extrusion billet and extruded
(step 100) to produce a product with improved mechanical properties
useful for subsequent service as a high temperature L1.sub.2
strengthened aluminum alloy. Extrusion results in optimum
mechanical properties in the extrusion direction. If more uniform
directional properties are required, forging and/or rolling (step
110) is necessary.
[0101] Sieving (step 20) is a preferred step in consolidation
because the final mechanical properties relate directly to the
particle size. Finer particle size results in finer L1.sub.2
particle dispersion. Optimum mechanical properties have been
observed with -450 mesh (30 micron) powder. Sieving (step 20) also
limits the defect size in the powder. Before sieving, the powder is
passivated with nitrogen gas in order to improve the efficiency of
sieving. If the as-atomized powder is oxygen deficient, the powder
can have a tendancy to stick together which will lower the
efficiency for sieving. Ultrasonic sieving is preferred for its
efficiency.
[0102] Blending (step 30) is a preferred step in the consolidation
process because it results in improved uniformity of particle size
distribution. Gas atomized L1.sub.2 aluminum alloy powder generally
exhibits a bimodal particle size distribution and cross blending of
separate powder batches tends to homogenize the particle size
distribution. Blending 30 is also preferred when separate metal
and/or ceramic powders are added to the L1.sub.2 base powder to
form bimodal or trimodal consolidated alloy microstructures.
[0103] Cryomilling (step 40) is a preferred step and is used to
refine the grain size of gas atomized L1.sub.2 aluminum alloy
powder as well as the final consolidated alloy microstructure.
Cryomilling is described in U.S. Pat. No. 6,902,699, Fritzemeier et
al. and in U.S. Pat. No. 7,344,675, Van Daam et al., both owned by
the assignee of the present invention, and are incorporated herein
in their entirety by reference. Cryomilling involves high-energy
ball milling under liquid nitrogen. The liquid nitrogen facilitates
efficient breaking up of powder particles. The liquid nitrogen
environment prevents oxidation and prevents frictional heating of
the powder and the resulting grain coarsening. During the process,
the powder particles are repeatedly sheared, fractured and cold
welded which results in a severely deformed structure containing a
high dislocation density that, with continued deformation, evolves
into a cellular structure consisting of extremely small dislocation
free grains separated by high angle grain boundaries with high
dislocation density. The grain size of the cellular microstructure
is typically less than 100 nm (0.04 microinch) and the structure is
considered a nanostructure.
[0104] In addition, the nitrogen environment results in the
formation of nitride particles that reside at the grain boundaries
and inside grains themselves and resist coarsening at higher
temperatures. Stearic acid is preferably added to the powder charge
to prevent excessive agglomeration and to promote fracturing and
rewelding of the L1.sub.2 aluminum alloy particles during
milling.
[0105] Following sieving (step 20), blending (step 30) and
cryomilling (step 40), the powders are transferred to a can (step
50) where the powder is vacuum degassed (step 60) at elevated
temperatures. The can (step 50) is an aluminum container having a
cylindrical, rectangular or other configuration with a central
axis. Vacuum degassing times can range from 12 hours to over 8
days. A temperature range of about 500.degree. F. (260.degree. C.)
to about 900.degree. F. (482.degree. C.) is preferred and about
750.degree. F. (399.degree. C.) is more preferred. Dynamic
degassing of large amounts of powder are preferred to static
degassing. In dynamic degassing, the can is preferably rotated
during degassing to expose all of the powder to a uniform
temperature. Degassing removes the stearic acid lubricant as well
as oxygen and hydrogen from the powder.
[0106] Following vacuum degassing (step 60), the vacuum line is
crimped and welded shut. The powder is then consolidated further by
uniaxially hot pressing the evacuated can (along its central axis
while radial movement is restrained) in a die or by hot isostatic
pressing (HIP) the can in an isostatic press. At this point the
powder charge is nearly 100 percent dense. The billet can be
compressed by blind die compaction (step 90) to further densify the
structure. Blind die compaction is preferred to further densify the
billet as prior consolidation processes may not provide 100%
density due to insufficient load available with the press. If
uniaxial hot pressing or hot isostatic pressing are used 100%
density can be achieved and blind die compaction may not be
required. Following densification, the can may be removed by
machining.
[0107] Following blind die compaction, the billet is machined into
an extrusion billet, copper jacketed and extruded (step 100). The
extrusion process preferably increases the hardness and improves
the tensile ductility. Extrusion imparts directional mechanical
properties to the material. Forging and/or rolling (step 110) can
improve the uniformity of the short transverse mechanical
properties.
[0108] FIG. 11 shows a 3-inch (7.62 cm) diameter copper jacketed
L1.sub.2 aluminum alloy billet ready for extrusion. FIG. 12 is a
photo of three 3-inch (7.62 cm) diameter extrusion dies.
Representative extrusions using the 3-inch (7.62 cm) diameter dies
are shown in FIG. 13. A 12-inch (30.48 cm) ruler is included in the
photo for size comparison. Larger 6-inch (15.24 cm) diameter
billets were also extruded. Machined 6-inch (15.24 cm) diameter
L1.sub.2 aluminum alloy extrusion billets are shown in FIG. 14.
FIG. 15 is a photo of a machined three-piece copper jacketed 6-inch
(15.24 cm) diameter billet assembly. A 12-inch (30.48 cm) ruler is
included in the photo for size comparison. The upright cylinder
behind the three-piece assembly is another machined, copper
jacketed L1.sub.2 aluminum alloy extrusion billet.
[0109] Extruded L1.sub.2 aluminum alloy rods from 6-inch diameter
billets are shown in FIG. 16. The top rod is 46 inches (116.8 cm)
long.
[0110] Representative processing parameters for 3-inch diameter
L1.sub.2 aluminum alloy billets are listed in Table 1.
TABLE-US-00001 TABLE 1 Processing Details of L12 Strengthened
Alloys Billet Degassing/VHP Extrusion Billet Extrusion Number
Temperature Temperature Temperature/Time Liner Temperature Speed 1
600.degree. F. (316.degree. C.) 450.degree. F. (232.degree. C.)
450.degree. F. (232.degree. C.) 450.degree. F. (232.degree. C.)
0.5''/min (1.2 cm/min) 2 700.degree. F. (371.degree. C.)
400.degree. F. (204.degree. C.) 400.degree. F. (204.degree. C.)
400.degree. F. (204.degree. C.) 0.5''/min (1.2 cm/min) 3
600.degree. F. (316.degree. C.) 650.degree. F. (343.degree. C.)
650.degree. F. (343.degree. C.) 650.degree. F. (343.degree. C.)
0.5''/min (1.2 cm/min) 4 550.degree. F. (288.degree. C.)
600.degree. F. (316.degree. C.) 600.degree. F. (316.degree. C.)
600.degree. F. (316.degree. C.) 0.5''/min (1.2 cm/min) 5
700.degree. F. (371.degree. C.) 450.degree. F. (232.degree. C.)
450.degree. F. (232.degree. C.) 450.degree. F. (232.degree. C.)
0.5''/min (1.2 cm/min) 6 700.degree. F. (371.degree. C.)
350.degree. F. (177.degree. C.) 350.degree. F. (177.degree. C.)
350.degree. F. (177.degree. C.) 0.5''/min (1.2 cm/min)
TABLE-US-00002 TABLE 2 Tensile Properties of L12 Strengthened
Aluminum Alloys Billet Tensile Strength, Yield Strength,
Elongation, Reduction Number ksi (MPa) Ksi (MPa) % in Area, % 1 113
(779) 103 (710) 3.8 6.7 2 119 (820) 112 (772) 3.2 7.0 3 101 (696)
97 (669) 3.9 9.7 4 111 (765) 105 (724) 2.7 8.3 5 118 (813) 102
(703) 2.1 3.5 6 120 (827) 108 (745) 0.5 2.0
Representative mechanical properties of extruded aluminum alloy
billets are listed in Table 2. Table 1 shows processing details of
degassing, vacuum hot pressing (VHP), and extrusion parameters used
for fabrication of this material. Degassing and consolidation of
billets were performed in the range of 550.degree. F. (288.degree.
C.) to 700.degree. F. (371.degree. C.). Extrusion was performed in
the 350.degree. F. (177.degree. C.) to 650.degree. F. (343.degree.
C.) temperature range where billet, die and liner temperatures were
maintained equal. Extrusion speed was maintained at 0.5 inches per
minute (1.27 cm per minute). Lower speed is desired for higher
strength due to less adiabatic heat generation during extrusion.
Table 2 includes tensile properties of extrusions that resulted
from the above processing parameters. The measured tensile strength
ranges from 101 ksi (696 MPa) to 120 ksi MPa) and yield strength
ranges from 97 ksi (667 MPa) to 108 ksi (745 MPa). These strength
values are significantly higher than commercially available
existing aluminum alloys. It should be noted that higher strength
was obtained for lower extrusion temperature conditions. For
example, billet number 3 showed a yield strength of 97 ksi (669
MPa) and tensile strength of 101 ksi (696 MPa) for an extrusion
temperature of 650.degree. F. (343.degree. C.), whereas billet
number 6 showed a yield strength of 108 ksi (745 MPa) and a tensile
strength of 120 ksi (827 MPa) for an extrusion temperature of
350.degree. F. (177.degree. C.). Billet number 2 showed a yield
strength of 112 ksi (772 MPa) and a tensile strength of 118 (813
MPa) for an extrusion temperature of 400.degree. F. (204.degree.
C.) which is very close to the strength obtained for an extrusion
temperature of 350.degree. F. (177.degree. C.). Ductility which is
measured by elongation and reduction in area also showed variations
with extrusion temperature. Higher extrusion temperatures resulted
in higher ductility. Higher strength values obtained for these
extrusions made from cryomilled billets suggest that cryomilling
has worked very effectively for L1.sub.2 strengthened aluminum
alloys.
[0111] The above mentioned processing steps (and others) for
consolidating L1.sub.2 aluminum alloy powder are well known to
those versed in the art and need not be described in more
detail.
[0112] Although the present invention has been described with
reference to preferred embodiments, workers skilled in the art will
recognize that changes may be made in form and detail without
departing from the spirit and scope of the invention.
* * * * *