U.S. patent application number 12/316046 was filed with the patent office on 2010-06-10 for method for forming high strength aluminum alloys containing l12 intermetallic dispersoids.
This patent application is currently assigned to United Technologies Corporation. Invention is credited to Awadh B. Pandey.
Application Number | 20100143177 12/316046 |
Document ID | / |
Family ID | 42231295 |
Filed Date | 2010-06-10 |
United States Patent
Application |
20100143177 |
Kind Code |
A1 |
Pandey; Awadh B. |
June 10, 2010 |
Method for forming high strength aluminum alloys containing L12
intermetallic dispersoids
Abstract
A method and apparatus produces high strength aluminum alloys
from a powder containing L1.sub.2 intermetallic dispersoids. The
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: |
42231295 |
Appl. No.: |
12/316046 |
Filed: |
December 9, 2008 |
Current U.S.
Class: |
419/30 ;
425/78 |
Current CPC
Class: |
C22C 1/0416 20130101;
B22F 2998/00 20130101; C22C 1/0491 20130101; B22F 2999/00 20130101;
B22F 2999/00 20130101; B22F 2998/00 20130101; B22F 2998/10
20130101; B22F 2998/10 20130101; B22F 3/1208 20130101; B22F 3/14
20130101; B22F 3/17 20130101; B22F 3/14 20130101; B22F 3/18
20130101; B22F 2201/20 20130101; B22F 3/20 20130101 |
Class at
Publication: |
419/30 ;
425/78 |
International
Class: |
B22F 1/00 20060101
B22F001/00; B22F 3/12 20060101 B22F003/12 |
Claims
1. A method for producing a high strength aluminum alloy billet
containing L1.sub.2 dispersoids, comprising the steps of: placing a
quantity of an aluminum alloy powder containing an L1.sub.2
dispersoid therein having a mesh size of less than 450 mesh in a
container; vacuum degassing the powder at a temperature of about
300.degree. F. (149.degree. C.) to about 900.degree. F.
(482.degree. C.) for about 0.5 hours to about 8 days; sealing the
degassed powder in the container under vacuum; heating the sealed
container at about 300.degree. F. (149.degree. C.) to about
900.degree. F. (482.degree. C.) for about 15 minutes 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 configuration with 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 minute to about 8 hours at a
temperature of 300.degree. F. (149.degree. C.) to about 900.degree.
F. (482.degree. C.) under uni-axial pressure of about 5 ksi to
about 100 ksi.
4. The method of claim 1, wherein the vacuum hot pressing produces
a billet of the aluminum alloy powder having a 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 load of about 100 tons to about 10,000 tons.
7. The method of claim 6, wherein the extrusion temperature is
about 300.degree. F. (149.degree. C.) to about 900.degree. F.
(482.degree. C.), the billet soak time is about 15 minutes to about
8 hours at a rate of about 0.2 inch per minute to about 20 inches
per minute and an extrusion ratio of about 2:1 to about 500:1.
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 6 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. Apparatus for producing a high strength aluminum alloy billet
containing L1.sub.2 dispersoids, comprising: a container for
holding a quantity of an aluminum alloy powder containing an
L1.sub.2 dispersoid therein having a mesh size of less than 450
mesh; a vacuum and heat source for degassing the powder at a
temperature of about 300.degree. F. (149.degree. C.) to about
900.degree. F. (482.degree. C.) for about 0.5 hours to about 8
days; sealing means for sealing the degassed powder in the
container under vacuum; a heater for heating the sealed container
at about 300.degree. F. (149.degree. C.) to about 900.degree. F.
(482.degree. C.) for about 15 minutes to eight hours; a vacuum hot
press for forming the heated container into a billet; and means for
removing the container from the thus formed billet.
13. The apparatus of claim 12, wherein the container is aluminum
having a central axis, and vacuum hot pressing is done along the
axis while restraining transverse movement of the container.
14. The apparatus of claim 12, wherein the vacuum hot pressing
includes blind die compaction for about 1 minute to about 8 hours
at a temperature of 300.degree. F. (149.degree. C.) to about
900.degree. F. (482.degree. C.) under uni-axial pressure of about 5
ksi to about 100 ksi.
15. The apparatus of claim 12, wherein the vacuum hot pressing
produces a billet of the aluminum alloy powder having a density of
about 100 percent.
16. The apparatus of claim 12, wherein the degassing includes
rotating the aluminum alloy powder to heat and exposing to vacuum
all the powder.
17. The apparatus of claim 12, wherein the formed billet is
extruded at a load of about 100 tons to about 10,000 tons.
18. The apparatus of claim 17, wherein the extrusion temperature is
about 300.degree. F. (149.degree. C.) to about 900.degree. F.
(482.degree. C.), the billet soak time is about 15 minutes to about
8 hours at a rate of about 0.2 inch per minute to about 20 inches
per minute, and an extrusion ratio of about 2:1 to about 500:1.
19. The apparatus 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 apparatus of claim 19, wherein the aluminum alloy powder
contains at least one third element selected from the group
consisting of silicon, magnesium, lithium, copper, zinc, and
nickel.
21. The apparatus of claim 20, 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 6 weight percent copper, about
3 to about 12 weight percent zinc, about 1 to about 12 weight
percent nickel.
22. The apparatus of claim 19, 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.
Description
CROSS-REFERENCE TO RELATED APPLICATION(S)
[0001] This application is related to the following co-pending
applications that are filed on even date herewith and are assigned
to the same assignee: VERSION PROCESS FOR HEAT TREATABLE L1.sub.2
ALUMINUM ALLOYS, Ser. No. 12/316,020, Attorney Docket No.
PA0006942U-U73.12-324KL; and A METHOD FOR PRODUCING HIGH STRENGTH
ALUMINUM ALLOY POWDER CONTAINING L1.sub.2 INTERMETALLIC
DISPERSOIDS, Ser. No. 12/316,047, Attorney Docket No.
PA0007539U-U73.12-338KL.
[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 Al 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 nanometer range
of about 30 to 100 nm. These alloys also have lower ductility.
SUMMARY
[0009] The present invention is a method for consolidating 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. The powders are then vacuum degassed in a container
that is then sealed. The sealed container (i.e. can) is vacuum hot
pressed to densify the powder charge and then compacted further by
blind die compaction or other suitable method. 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] FIGS. 6A and 6B are SEM photos of gas atomized L1.sub.2
aluminum alloy powder.
[0017] FIGS. 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. 8 is a diagram showing the processing steps to
consolidate L1.sub.2 aluminum alloy powder.
[0019] FIG. 9 is a photo of a 3-inch diameter copper jacketed
L1.sub.2 aluminum alloy billet.
[0020] FIG. 10 is a photo of extrusion dies for 3-inch diameter
billet.
[0021] FIG. 11 is a photo of extruded L1.sub.2 aluminum alloy rods
from 3-inch diameter billets.
[0022] FIG. 12 is a photo of machined L1.sub.2 aluminum alloy
billets.
[0023] FIG. 13 is a photo of a machined three-piece L1.sub.2
aluminum alloy billet assembly for 6-inch copper jacketed extrusion
billet.
[0024] FIG. 14 is a photo of extruded L1.sub.2 aluminum alloy rods
from 6-inch diameter billets.
DETAILED DESCRIPTION
1. L1.sub.2 Aluminum Alloys
[0025] Alloy powders of 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.
[0026] 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
[0027] 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
[0028] 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 is complete
solubility of lithium in the rapid solidified inventive alloys
discussed herein.
[0029] The binary aluminum copper system is a simple eutectic at 32
weight percent copper and 1018.degree. F. (548.degree. C.). There
is complete solubility of copper in the rapidly solidified
inventive alloys discussed herein.
[0030] 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.
[0031] 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.
[0032] 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.
[0033] 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.
[0034] 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.
[0035] 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.
[0036] 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.
[0037] 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.
[0038] 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.
[0039] 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.
[0040] 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.
[0041] Titanium forms Al.sub.3Ti dispersoids in the aluminum matrix
that have an L1.sub.2 structure in the metastable condition and
D0.sub.22 structure in the equilibrium condition. The metastable
Al.sub.3Ti despersoids 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.
[0042] 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.
[0043] 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.
[0044] 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.
[0045] 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.
[0046] Exemplary aluminum alloys for this invention include, but
are not limited to (in weight percent unless otherwise
specified):
[0047] about Al-M-(0.1-4)Sc-(0.1-20)Gd;
[0048] about Al-M-(0.1-20)Er-(0.1-20)Gd;
[0049] about Al-M-(0.1-15)Tm-(0.1-20)Gd;
[0050] about Al-M-(0.1-25)Yb-(0.1-20)Gd;
[0051] about Al-M-(0.1-25)Lu-(0.1-20)Gd;
[0052] about Al-M-(0.1-4)Sc-(0.1-20)Y;
[0053] about Al-M-(0.1-20)Er-(0.1-20)Y;
[0054] about Al-M-(0.1-15)Tm-(0.1-20)Y;
[0055] about Al-M-(0.1-25)Yb-(0.1-20)Y;
[0056] about Al-M-(0.1-25)Lu-(0.1-20)Y;
[0057] about Al-M-(0.1-4)Sc-(0.05-4)Zr;
[0058] about Al-M-(0.1-20)Er-(0.05-4)Zr;
[0059] about Al-M-(0.1-15)Tm-(0.05-4)Zr;
[0060] about Al-M-(0.1-25)Yb-(0.05-4)Zr;
[0061] about Al-M-(0.1-25)Lu-(0.05-4)Zr;
[0062] about Al-M-(0.1-4)Sc-(0.05-10)Ti;
[0063] about Al-M-(0.1-20)Er-(0.05-10)Ti;
[0064] about Al-M-(0.1-15)Tm-(0.05-10)Ti;
[0065] about Al-M-(0.1-25)Yb-(0.05-10)Ti;
[0066] about Al-M-(0.1-25)Lu-(0.05-10)Ti;
[0067] about Al-M-(0.1-4)Sc-(0.05-10)Hf;
[0068] about Al-M-(0.1-20)Er-(0.05-10)Hf;
[0069] about Al-M-(0.1-15)Tm-(0.05-10)Hf;
[0070] about Al-M-(0.1-25)Yb-(0.05-10)Hf;
[0071] about Al-M-(0.1-25)Lu-(0.05-10)Hf;
[0072] about Al-M-(0.1-4)Sc-(0.05-5)Nb;
[0073] about Al-M-(0.1-20)Er-(0.05-5)Nb;
[0074] about Al-M-(0.1-15)Tm-(0.05-5)Nb;
[0075] about Al-M-(0.1-25)Yb-(0.05-5)Nb; and
[0076] about Al-M-(0.1-25)Lu-(0.05-5)Nb.
[0077] M is at least one of about (4-25) weight percent silicon,
(1-8) weight percent magnesium, (0.5-3) weight percent lithium,
(0.2-6) weight percent copper, (3-12) weight percent zinc, and
(1-12) weight percent nickel.
[0078] 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.
[0079] 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.
[0080] 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.
[0081] The amount of copper present in the fine grain matrix, if
any, may vary from about 0.2 to about 6 weight percent, more
preferably from about 0.5 to about 5 weight percent, and even more
preferably from about 2 to about 4.5 weight percent.
[0082] 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.
[0083] 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.
[0084] 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.
[0085] 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.
[0086] 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.
[0087] 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
[0088] 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.
[0089] 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.
[0090] 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.
[0091] 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.
[0092] 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.
[0093] 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.
[0094] 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.
[0095] 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.
[0096] 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
[0097] 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 the
inventive 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. The process of consolidating the alloy powders into
useful forms is schematically illustrated in FIG. 8.
[0098] 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.
[0099] Sieving (step 20) is a critical 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. Sufficient 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 minimize reaction of the
powder with atmosphere. The powder is stored in a nitrogen
atmosphere to prevent oxidation. However, if the powder is
completely free from oxides, it sticks together reducing the
efficiency of sieveing. If oxygen in the powder is too high, it has
a deleterious effect on mechanical properties. There is an optimal
oxygen level which is desired so that it does not create problems
with sieving and yields good mechanical properties. The oxygen
content of the powder is between about 1 ppm and 2000 ppm,
preferred between about 10 ppm to 1000 ppm and most preferred
between about 25 ppm to about 500 ppm. Ultrasonic sieving is
preferred for its efficiency.
[0100] 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 (step 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.
[0101] Following sieving (step 20) and blending (step 30), the
powders are transferred to 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 about 0.5 hours to about 8 days, more preferably it can range
from about 4 hours to 7 days, even more preferably it can range
from about 8 hours to about 6 days. A temperature range of about
300.degree. F. (149.degree. C.) to about 900.degree. F.
(482.degree. C.) is preferred and about 600.degree. F. (316.degree.
C.) to about 850.degree. F. (454.degree. C.) is more preferred and
650.degree. F. (343.degree. C.) to about 850.degree. F.
(454.degree. C.) is most preferred. Dynamic degassing of large
amounts of powder is 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
oxygen and hydrogen from the powder.
[0102] The role of dynamic degassing is to remove oxygen and
hydrogen more efficiently than static degassing. Dynamic degassing
is very important for large lbillets to reduce time and temperature
required for degassing. Static degassing works well for small sizes
of billets and small quantity of powder as it does not take long
time to degas effectively. For large billets, it can take several
days to degas at high temperatures which can coarsen the material
microstructure and reduce the strength. In addition, the process
efficiency goes down with longer time for degassing.
[0103] Following vacuum degassing (step 60), the vacuum line is
crimped and welded shut. The powder is then consolidated further by
unaxially 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. The billet can be
compressed by blind die compaction (step 90) to further densify the
structure if it is not 100% dense. At this point the can may be
removed by machining.
[0104] Following blind die compaction, the billet is machined into
an extrusion billet, copper jacketed and extruded (step 100).
Alternatively, the billet can be extruded directly after blind die
compaction without machining and without a copper jacket. A copper
jacket is preferred to provide improved lubrication. However, it is
not essential for extrusion of billets. The extrusion process
preferably improves the hardness and improves the tensile
ductility. Extrusion imparts directional mechanical properties to
the material. Forging and/or rolling (step 110) can improve the
transverse mechanical properties leading to isotropic
properties.
[0105] FIG. 9 shows a 3-inch diameter copper jacketed L1.sub.2
aluminum alloy billet ready for extrusion. FIG. 10 is a photo of
three 3-inch diameter extrusion dies. Representative extrusions
using the 3-inch diameter dies are shown in FIG. 11. A 12-inch
ruler is included in the photo for size comparison. Larger 6-inch
diameter billets were also extruded. Machined 6-inch diameter
L1.sub.2 aluminum alloy extrusion billets are shown in FIG. 12.
FIG. 13 is a photo of a machined three-piece copper jacketed 6-inch
diameter billet assembly. A 12-inch 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.
[0106] Extruded L1.sub.2 aluminum alloy rods from 6-inch diameter
billets are shown in FIG. 14. The top rod is 46 inches long.
[0107] Representative processing parameters for 3-inch diameter
L1.sub.2 aluminum alloy billets are listed in Table 1.
TABLE-US-00001 TABLE 1 Extrusion Billet Powder Vacuum Hot
Temperature/Die Billet Degassing Pressing and Container
Breakthrough Extrusion Extrusion # Temp/Time Temp/Time Temperature
Load Speed Ratio 1 500.degree. F./19 h 500.degree. F./1 h
700.degree. F./650.degree. F. 640 Tons 0.5''/min 10:01 2
550.degree. F./19 h 550.degree. F./1 h 600.degree. F./600.degree.
F. 601 Tons 0.75''/min 6:01 3 600.degree. F./19 h 600.degree. F./1
h 650.degree. F./650.degree. F. 648 Tons 0.5''/min 10:01 4
650.degree. F./19 h 700.degree. F./1 h 700.degree. F./650.degree.
F. 655 Tons 0.5''/min 10:01 5 700.degree. F./19 h 700.degree. F./1
h 700.degree. F./650.degree. F. 634 Tons 0.5''/min 10:01 6
700.degree. F./19 h 700.degree. F./1 h 600.degree. F./600.degree.
F. 621 Tons 0.5''/min 10:01 7 750.degree. F./19 h 700.degree. F./1
h 700.degree. F./650.degree. F. 550 Tons 0.5''/min 10:01 8
600.degree. F./19 h 700.degree. F./1 h 700.degree. F./650.degree.
F. 612 Tons 0.5''/min 10:01
[0108] Table 1 shows powder processing data that includes degassing
temperature, time, consolidation temperature and time, extrusion
temperature, ratio and load experienced during extrusion, extrusion
die and billet temperatures. These processing parameters were used
to degas, consolidate and extrude L1.sub.2 aluminum alloy powders.
The degassing temperature range of 500.degree. F. to 750.degree. F.
(260.degree. C. to 399.degree. C.) for a constant degas time of 19
hours was evaluated. Vacuum hot pressing at a temperature range of
500.degree. F. to 700.degree. F. (260.degree. C. to 371.degree. C.)
for a constant time of 1 hour was evaluated. Since the billet does
not usually have good ductility to provide sufficient integrity for
testing, billets were extruded for providing deformation to impart
ductility in the billet. Extrusion billet temperature, die
temperature and container temperature varied from 650.degree. F. to
700.degree. F. (343.degree. C. to 371.degree. C.). Extrusion speed
varied from 0.5 inch per minute to 0.75 inch per minute and
extrusion ratio varied from 6:1 to 10:1. Extrusion load varied from
550 tons to 655 tons depending on process parameters used for the
powder. Breakthrough load depends on degassing temperature and
vacuum hot pressing temperature in addition to extrusion
temperature, extrusion speed and extrusion ratio. Breakthrough load
decreased with an increase in degassing temperature and vacuum hot
pressing temperature. The load has decreased from 640 tons to 550
tons as we increased the degassing temperature from 500.degree. F.
to 750.degree. F. (260.degree. C. to 399.degree. C.). Breakaway
load is important in order to make successful extrusions. If the
load requirement is higher than the capacity of the extrusion
press, then the press will stall and material will not be extruded.
It is very important to select the degassing and vacuum hot
pressing temperature in such a way that successful extrusions are
produced with good mechanical properties.
[0109] Metallographic examination of microstructures of extruded
bars in longitudinal and transverse directions showed deformation
bands from the extrusion process in longitudinal sections. The
transverse microstroctures showed more uniform microstructures
without any deformation bands. The grain size could not be resolved
by optical microscopy as it is very fine. Very fine dispersoids
were observed throughout the material.
[0110] X-ray diffraction patterns were taken of the powders and
forgings to examine phase content. Diffraction patterns of the
powders showed only two phases: aluminum and aluminum nickel. Since
the lattice parameters of aluminum and Al.sub.3Sc dispersoids are
very similar, the peaks for aluminum and Al.sub.3Sc dispersoids
cannot be resolved. Diffraction patterns of extrusions showed
additional phases based on gadolinium nickel and nickel zirconium.
These phases were produced during powder processing.
[0111] The hydrogen content in extrusions produced from powders
which were degassed at different temperatures from 500.degree. F.
to 750.degree. F. (260.degree. C. to 399.degree. C.) showed the
hydrogen content in the extrusions from powder degassed at
500.degree. F. to 600.degree. F. (260.degree. C. to 343.degree. C.)
were higher than those degassed at 650.degree. F.-750.degree. F.
(343.degree. C. to 399.degree. C.). 650.degree. F. (343.degree. C.)
is a critical temperature above which degassing is more effective
in L12 powder. There is no appreciable benefit in degassing at
higher temperature than 650.degree. F. (343.degree. C.) in terms of
hydrogen content for a constant time of 19 hours. If time is varied
for degassing, results will change based on the diffusion kinetics.
For a given temperature, longer time will give better degassing
based on the diffusion kinetics of L12 powder. It is desired to
have low hydrogen in the material as hydrogen has deleterious
effects on ductility of the material.
[0112] Representative mechanical properties of extruded aluminum
alloy billets are listed in Table 2.
TABLE-US-00002 TABLE 2 Tensile properties of extruded L1.sub.2
aluminum alloy Ultimate Tensile Yield Reduction in Billet #
Strength, ksi Strength, ksi Elongation, % Area, % 1 115.4 102.4 4
4.3 2 113.2 100.8 4 13 3 115.4 104.2 3.7 8.5 4 113.5 101.9 5.4 7.9
5 110 101.3 4.7 15 6 101.2 93.2 10.3 18 7 107.8 96.1 6.6 17.5 8
115.9 101.7 5 8.5
[0113] Table 2 shows tensile properties of extrusions made from
powders degassed at different temperatures. The yield strengths and
ultimate tensile strengths of these L1.sub.2 based alloys are
excellent. These strength values are much higher than the strengths
of commercial aluminum alloys including 6061 Al, 2124 Al and 7075
Al suggesting that the processing parameters used for the
production of this inventive material has worked well. A tensile
strength over 100 ksi for L1.sub.2 aluminum alloy is remarkable as
it can provide significant weight savings by replacing high
strength aluminum alloy, titanium nickel and steel alloys. In
addition, the elongation and reduction in area values for this
L1.sub.2 alloy are also very good. The yield strength remains
fairly constant over 100 ksi for degassing and vacuum hot pressing
temperature ranges of 500.degree. F.-650.degree. F. (260.degree. C.
to 343.degree. C.). The yield strength decreased slightly for
degassing and vacuum hot pressing temperature ranges of 700.degree.
F. to 750.degree. F. (371.degree. C. to 399.degree. C.). The
ductility measured by elongation and reduction in area, however,
increased significantly with an increase in degassing temperature.
Reduction in area has increased almost two times for material
degassed in the temperature range of 700.degree. F.-750.degree. F.
(371.degree. C. to 399.degree. C.) compared to material that was
degassed in the temperature range of 500.degree. F.-650.degree. F.
(260.degree. C. to 343.degree. C.). These results are expected
based on strengthening models including an Orowan strengthening
model and a Hall-Petch strengthening model. Vacuum degassing is
more effective when the powder is degassed at higher temperatures
as indicated by lower hydrogen content. Lower hydrogen results in
higher ductility of material as measured by elongation and
reduction in area. However, strength is expected to decrease with
an increase in degassing temperature because strengthening
precipitates will coarsen with an increase in degassing
temperature, which is consistent with observed results from
material degassed at 700.degree. F.-750.degree. F. (371.degree. C.
to 399.degree. C.). These results indicate that the properties of
L1.sub.2 alloy can be varied by controlling the degassing and
vacuum hot pressing temperature. In order to have a balanced
combination of strength and ductility in these L1.sub.2 alloys, the
alloys need to be degassed and vacuum hot pressed at the
recommended temperatures and times. The results obtained here
demonstrate that the present invention worked very well.
[0114] The fracture surfaces of tested samples show presence of
dimples indicating ductile fracture where voids nucleate, grow and
finally coalesce to failure. The fracture surface morphology
provided an evidence of ductile failure mode which is consistent
with good elongation and reduction in area values.
[0115] 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.
* * * * *