U.S. patent application number 12/589104 was filed with the patent office on 2011-04-21 for method for fabrication of tubes using rolling and extrusion.
This patent application is currently assigned to United Technologies Corporation. Invention is credited to Awadh B. Pandey.
Application Number | 20110091345 12/589104 |
Document ID | / |
Family ID | 43558355 |
Filed Date | 2011-04-21 |
United States Patent
Application |
20110091345 |
Kind Code |
A1 |
Pandey; Awadh B. |
April 21, 2011 |
Method for fabrication of tubes using rolling and extrusion
Abstract
A method for producing a high strength aluminum alloy tubing
containing L1.sub.2 dispersoids from an aluminum alloy powder
containing the L1.sub.2 dispersoids. The powder is consolidated
into a billet having a density of about 100 percent. The tube is
formed by at least one of direct extrusion, Mannesmann process,
pilgering, and rolling.
Inventors: |
Pandey; Awadh B.; (Jupiter,
FL) |
Assignee: |
United Technologies
Corporation
Hartford
CT
|
Family ID: |
43558355 |
Appl. No.: |
12/589104 |
Filed: |
October 16, 2009 |
Current U.S.
Class: |
419/12 ; 419/13;
419/15; 419/17; 419/19; 419/28; 75/228; 75/235; 75/236; 75/244;
75/249 |
Current CPC
Class: |
B22F 2999/00 20130101;
B22F 2998/10 20130101; C22C 21/00 20130101; B22F 2999/00 20130101;
B22F 5/106 20130101; B22F 3/14 20130101; B22F 3/14 20130101; B22F
2998/10 20130101; B22F 2998/10 20130101; B22F 2998/10 20130101;
C22C 32/00 20130101; B22F 3/17 20130101; C22C 1/0416 20130101; B22F
3/14 20130101; B22F 3/18 20130101; B22F 3/14 20130101; B22F 2201/20
20130101; B22F 3/20 20130101 |
Class at
Publication: |
419/12 ; 419/28;
419/19; 419/13; 419/17; 419/15; 75/249; 75/228; 75/235; 75/236;
75/244 |
International
Class: |
B22F 5/12 20060101
B22F005/12; B22F 3/24 20060101 B22F003/24; B22F 3/12 20060101
B22F003/12; B22F 3/20 20060101 B22F003/20; B22F 3/17 20060101
B22F003/17; B22F 3/18 20060101 B22F003/18; B22F 9/00 20060101
B22F009/00; C22C 21/00 20060101 C22C021/00 |
Claims
1. A method for forming a high strength aluminum alloy component
containing L1.sub.2 dispersoids, comprising the steps of: placing
in a container a quantity of an aluminum alloy powder containing an
L1.sub.2 dispersoid L1.sub.2 comprising 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;
the alloy powder having a mesh size of less than 350 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 solid
billet; removing the container from the formed billet; and forming
the billet into a tube.
2. The method of claim 1, wherein the billet is rolled into a sheet
and the sheet is rolled into a tube.
3. The method of claim 1, wherein the billet is formed into a tube
by the Mannesmann process.
4. The method of claim 1, wherein the billet is formed into a tube
by the pilgering process.
5. The method of claim 1, wherein the billet is extruded directly
into a tube.
6. The method of claim 1, wherein the aluminum alloy powder
contains at least one metal selected from the group comprising:
about 1.0 to about 8.0 weight percent magnesium, (4-25) weight
percent silicon, (0.1-3) weight percent manganese, about 0.5 to
about 3.0 weight percent lithium, about 0.2 to about 6.0 weight
percent copper, about 3.0 to about 12.0 weight percent zinc, and
about 1.0 to about 12.0 weight percent nickel.
7. The method of claim 1, wherein the aluminum alloy powder
contains at least one ceramic selected from the group comprising:
about 5 to about 40 volume percent aluminum oxide, about 5 to about
40 volume percent silicon carbide, about 5 to about 40 volume
percent aluminum nitride, about 5 to about 40 volume percent
titanium diboride, about 5 to about 40 volume percent titanium
boride, about 5 to about 40 volume percent boron carbide and about
5 to about 40 volume percent titanium carbide.
8. The method of claim 5, wherein the tube extrusion is carried out
at a temperature of from about 300.degree. F. (149.degree. C.) to
about 850.degree. F. (454.4.degree. C.).
9. The method of claim 5, wherein the tube extrusion rate is about
0.1 min.sup.-1 to 25 min.sup.-1.
10. The method of claim 1, wherein the billet temperature ranges
from about 300.degree. F. (149.degree. C.) to about 850.degree. F.
(454.4.degree. C.) and the billet has a soak time ranging from
about 1 hour to about 10 hours.
11. A high strength aluminum alloy tube, comprising: a vacuum hot
pressed aluminum alloy billet containing an L1.sub.2 dispersoid
comprising 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; and the billet
being formed into a tube.
12. The alloy tube of claim 11, wherein the alloy powder contains
at least one metal selected from the group comprising: about 1.0 to
about 8.0 weight percent magnesium, (4-25) weight percent silicon,
(0.1-3) weight percent manganese, about 0.5 to about 3.0 weight
percent lithium, about 0.02 to about 6.0 weight percent copper,
about 3.0 to about 12.0 weight percent zinc and about 1.0 to about
12.0 weigh percent nickel.
13. The alloy tube of claim 11, wherein the aluminum alloy tube
powder contains at least one ceramic selected from the group
comprising: about 5 to about 40 volume percent aluminum oxide,
about 5 to about 40 volume percent silicon carbide, about 5 to
about 40 volume percent aluminum nitride, about 5 to about 40
volume percent titanium diboride, about 5 to about 40 volume
percent titanium boride, about 5 to about 40 volume percent boron
carbide and about 5 to about 40 volume percent titanium
carbide.
14. The alloy tube of claim 11, wherein the tube is formed by
rolling a rolled sheet into a tube.
15. The alloy tube of claim 11, wherein the tube is formed by the
Mannesmann process.
16. The alloy tube of claim 11, wherein the tube is formed by the
pilgering process.
17. The alloy tube of claim 11, wherein the tube is formed by
direct extrusion.
18. The alloy tube of claim 17, wherein the extrusion is carried
out at a temperature from about 300.degree. F. (149.degree. C.) to
about 850.degree. F. (454.4.degree. C.).
19. The alloy tube of claim 17, wherein the extrusion rate is about
0.1 min.sup.-1 to about 25 min.sup.-1.
20. The alloy tube of claim 11, wherein the tube forming
temperature is about 300.degree. F. (149.degree. C.) to about
850.degree. F. (454.4.degree. C.).
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 into
aluminum tubing.
[0004] The combination of high strength, ductility, and fracture
toughness, as well as low density, make aluminum alloys natural
candidates for a variety of applications. Because of its low weight
and corrosion resistance, aluminum alloys are of interest in the
manufacture and use of tubing for many applications.
[0005] Powder processing plays an important role in the mechanical
properties of L1.sub.2 strengthened aluminum alloys due to the
finer grain size and finer precipitate size than can be produced by
other processing techniques. Powder metallurgy also allows greater
volume fractions of alloying elements, which further increases the
strength of the alloy.
[0006] 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.
[0007] 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.
[0008] U.S. Pat. No. 6,248,453 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 discloses a high strength aluminum
alloy that contains scandium and other elements that is
strengthened by L1.sub.2 dispersoids.
[0009] L1.sub.2 strengthened aluminum alloys have high strength and
improved fatigue properties compared to commercial 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 high ductility.
SUMMARY
[0010] The present invention is a method for consolidating aluminum
alloy powders into useful components such as tubing with 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.
[0011] The aluminum alloy tubes are formed from consolidated
billets by rolling of extruded bars to produce sheets and then
rolling the sheets into tubes, by the Mannesmann process, by
pilgering, by direct extrusion and by other processes known to
those in the art. Deformation processing of these alloys produces
considerable improvement in mechanical properties, especially
ductility compared to the consolidated billet.
BRIEF DESCRIPTION OF THE DRAWINGS
[0012] FIG. 1 is an aluminum scandium phase diagram.
[0013] FIG. 2 is an aluminum erbium phase diagram.
[0014] FIG. 3 is an aluminum thulium phase diagram.
[0015] FIG. 4 is an aluminum ytterbium phase diagram.
[0016] FIG. 5 is an aluminum lutetium phase diagram.
[0017] FIG. 6A is a schematic diagram of a vertical gas
atomizer.
[0018] FIG. 6B is a close up view of nozzle 108 in FIG. 6A.
[0019] FIGS. 7A and 7B are SEM photos of the inventive aluminum
alloy powder.
[0020] FIGS. 8A and 8B are optical micrographs showing the
microstructure of gas atomized L1.sub.2 aluminum alloy powder.
[0021] FIG. 9 is a diagram showing the steps of the gas atomization
process.
[0022] FIG. 10 is a diagram showing the processing steps to
consolidate the L1.sub.2 aluminum alloy powder into tubing.
[0023] FIG. 11 is a schematic diagram of blind die compaction.
[0024] FIG. 12 is a schematic diagram of tube rolling.
[0025] FIG. 13 is a schematic diagram of the Mannesmann
process.
[0026] FIG. 14 is a schematic diagram of the pilgering process.
[0027] FIG. 15 A-D are schematic diagrams showing reduction in wall
thickness during pilgering.
[0028] FIG. 16 is a schematic diagram showing direct extrusion of
tubes.
DETAILED DESCRIPTION
1. L1.sub.2 Aluminum Alloys
[0029] 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, manganese, 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.
[0030] The alloy may also include at least one ceramic
reinforcement. Aluminum oxide, silicon carbide, aluminum nitride,
titanium diboride, boron carbide and titanium carbide are suitable
ceramic reinforcements.
[0031] 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.
[0032] The binary aluminum silicon system is a simple eutectic at
12.6 weight percent silicon and 1070.6.degree. F. (577.degree. C.).
There is complete solubility of silicon and aluminum in the rapidly
solidified inventive alloys discussed herein.
[0033] The binary aluminum manganese system is a simple eutectic at
about 2 weight percent manganese and 1216.4.degree. F. (658.degree.
C.). There is complete solubility of manganese and aluminum in the
rapidly solidified inventive alloys discussed herein.
[0034] 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.
[0035] 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.
[0036] 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 supersaturated 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.
[0037] 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.
[0038] 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.
[0039] 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, manganese 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.
[0040] 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, manganese 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.
[0041] 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, manganese 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.
[0042] 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, manganese 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.
[0043] 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, manganese 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.
[0044] 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.
[0045] 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.
[0046] 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.
[0047] 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.
[0048] 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.
[0049] 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.
[0050] The aluminum oxide, silicon carbide, aluminum nitride,
titanium di-boride, titanium boride, boron carbide and titanium
carbide locate at the grain boundary and within the grain boundary
to restrict dislocations from going around particles of the ceramic
particles when the alloy is under stress. When dislocations form,
they become attached with the ceramic particles on the departure
side. Thus, more energy is required to detach the dislocation and
the alloy has increased strength. To accomplish this, the particles
of ceramic have to have a fine size, a moderate volume fraction in
the alloy, and form a good interface between the matrix and the
reinforcement. A working range of particle sizes is from about 0.5
to about 50 microns, more preferably about 1 to about 20 microns,
and even more preferably about 1 to about 10 microns. The ceramic
particles can break during blending and the average particle size
will decrease as a result.
[0051] 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.
[0052] 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.
[0053] Exemplary aluminum alloys for this invention include, but
are not limited to (in weight percent unless otherwise
specified):
[0054] about Al-M-(0.1-4)Sc-(0.1-20)Gd;
[0055] about Al-M-(0.1-20)Er-(0.1-20)Gd;
[0056] about Al-M-(0.1-15)Tm-(0.1-20)Gd;
[0057] about Al-M-(0.1-25)Yb-(0.1-20)Gd;
[0058] about Al-M-(0.1-25)Lu-(0.1-20)Gd;
[0059] about Al-M-(0.1-4)Sc-(0.1-20)Y;
[0060] about Al-M-(0.1-20)Er-(0.1-20)Y;
[0061] about Al-M-(0.1-15)Tm-(0.1-20)Y;
[0062] about Al-M-(0.1-25)Yb-(0.1-20)Y;
[0063] about Al-M-(0.1-25)Lu-(0.1-20)Y;
[0064] about Al-M-(0.1-4)Sc-(0.05-4)Zr;
[0065] about Al-M-(0.1-20)Er-(0.05-4)Zr;
[0066] about Al-M-(0.1-15)Tm-(0.05-4)Zr;
[0067] about Al-M-(0.1-25)Yb-(0.05-4)Zr;
[0068] about Al-M-(0.1-25)Lu-(0.05-4)Zr;
[0069] about Al-M-(0.1-4)Sc-(0.05-10)Ti;
[0070] about Al-M-(0.1-20)Er-(0.05-10)Ti;
[0071] about Al-M-(0.1-15)Tm-(0.05-10)Ti;
[0072] about Al-M-(0.1-25)Yb-(0.05-10)Ti;
[0073] about Al-M-(0.1-25)Lu-(0.05-10)Ti;
[0074] about Al-M-(0.1-4)Sc-(0.05-10)Hf;
[0075] about Al-M-(0.1-20)Er-(0.05-10)Hf;
[0076] about Al-M-(0.1-15)Tm-(0.05-10)Hf;
[0077] about Al-M-(0.1-25)Yb-(0.05-10)Hf;
[0078] about Al-M-(0.1-25)Lu-(0.05-10)Hf;
[0079] about Al-M-(0.1-4)Sc-(0.05-5)Nb;
[0080] about Al-M-(0.1-20)Er-(0.05-5)Nb;
[0081] about Al-M-(0.1-15)Tm-(0.05-5)Nb;
[0082] about Al-M-(0.1-25)Yb-(0.05-5)Nb; and
[0083] about Al-M-(0.1-25)Lu-(0.05-5)Nb.
[0084] M is at least one of about (1-8) weight percent magnesium,
(4-25) weight percent silicon, (0.1-3) weight percent manganese,
(0.5-3) weight percent lithium, (0.2-6) weight percent copper,
(3-12) weight percent zinc, and (1-12) weight percent nickel.
[0085] 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.
[0086] 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 5 to about 20 weight percent, and even more
preferably from about 6 to about 14 weight percent.
[0087] The amount of manganese present in the fine grain matrix, if
any, may vary from about 0.1 to about 3 weight percent, more
preferably from about 0.2 to about 2 weight percent, and even more
preferably from about 0.3 to about 1 weight percent.
[0088] 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.
[0089] 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.
[0090] 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.
[0091] 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.
[0092] 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 103.degree. C./second.
[0093] 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 103.degree. C./second.
[0094] 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 103.degree. C./second.
[0095] 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 103.degree. C./second.
[0096] 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 103.degree. C./second.
[0097] 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.
[0098] 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.
[0099] 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.
[0100] 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.
[0101] 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.
[0102] 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.
[0103] 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 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. L1.sub.2 Alloy Powder Formation and Consolidation
[0104] The highest cooling rates observed in commercially viable
processes are achieved by gas atomization of molten metals to
produce powder. Gas atomization is a two fluid process wherein a
stream of molten metal is disintegrated by a high velocity gas
stream. The end result is that the particles of molten metal
eventually become spherical due to surface tension and finely
solidify in powder form. Heat from the liquid droplets is
transferred to the atomization gas by convection. The
solidification rates, depending on the gas and the surrounding
environment, can be very high and can exceed 10.sup.6.degree.
C./second. Cooling rates greater than 10.sup.3.degree. C./second
are typically specified to ensure supersaturation of alloying
elements in gas atomized L12 aluminum alloy powder in the inventive
process described herein.
[0105] A schematic of typical vertical gas atomizer 100 is shown in
FIG. 6A. FIG. 6A is taken from R. Germain, Powder Metallurgy
Science Second Edition MPIF (1994) (chapter 3, p. 101) and is
included herein for reference. Vacuum or inert gas induction melter
102 is positioned at the top of free flight chamber 104. Vacuum
induction melter 102 contains melt 106 which flows by gravity or
gas overpressure through nozzle 108. A close up view of nozzle 108
is shown in FIG. 6B. Melt 106 enters nozzle 108 and flows downward
till it meets the high pressure gas stream from gas source 110
where it is transformed into a spray of droplets. The droplets
eventually become spherical due to surface tension and rapidly
solidify into spherical powder 112 which collects in collection
chamber 114. The gas recirculates through cyclone collector 116
which collects fine powder 118 before returning to the input gas
stream. As can be seen from FIG. 6A, the surroundings to which the
melt and eventual powder are exposed are completely controlled.
[0106] There are many effective nozzle designs known in the art to
produce spherical metal powder. Designs with short gas-to-melt
separation distances produce finer powders. Confined nozzle designs
where gas meets the molten stream at a short distance just after it
leaves the atomization nozzle are preferred for the production of
the inventive L1.sub.2 aluminum alloy powders disclosed herein.
Higher superheat temperatures cause lower melt viscosity and longer
cooling times. Both result in smaller spherical particles.
[0107] A large number of processing parameters are associated with
gas atomization that affect the final product. Examples include
melt superheat, gas pressure, metal flow rate, gas type, and gas
purity. In gas atomization, the particle size is related to the
energy input to the metal. Higher gas pressures, higher superheat
temperatures and lower metal flow rates result in smaller particle
sizes. Higher gas pressures provide higher gas velocities for a
given atomization nozzle design.
[0108] To maintain purity, inert gases are used, such as helium,
argon, and nitrogen. Helium is preferred for rapid solidification
because the high heat transfer coefficient of the gas leads to high
quenching rates and high supersaturation of alloying elements.
[0109] Lower metal flow rates and higher gas flow ratios favor
production of finer powders. The particle size of gas atomized
melts typically has a log normal distribution. In the turbulent
conditions existing at the gas/metal interface during atomization,
ultra fine particles can form that may reenter the gas expansion
zone. These solidified fine particles can be carried into the
flight path of molten larger droplets resulting in agglomeration of
small satellite particles on the surfaces of larger particles. An
example of small satellite particles attached to inventive
spherical L1.sub.2 aluminum alloy powder is shown in the scanning
electron microscopy (SEM) micrographs of FIGS. 7A and 7B at two
magnifications. The spherical shape of gas atomized aluminum powder
is evident. The spherical shape of the powder is suggestive of
clean powder without excessive oxidation. Higher oxygen in the
powder results in irregular powder shape. Spherical powder helps in
improving the flowability of powder which results in higher
apparent density and tap density of the powder. The satellite
particles can be minimized by adjusting processing parameters to
reduce or even eliminate turbulence in the gas atomization process.
The microstructure of gas atomized aluminum alloy powder is
predominantly cellular as shown in the optical micrographs of
cross-sections of the inventive alloy in FIGS. 8A and 8B at two
magnifications. The rapid cooling rate suppresses dendritic
solidification common at slower cooling rates resulting in a finer
microstructure with minimum alloy segregation.
[0110] Oxygen and hydrogen in the powder can degrade the mechanical
properties of the final part. It is preferred to limit the oxygen
in the L1.sub.2 alloy powder to about 1 ppm to 2000 ppm. Oxygen is
intentionally introduced as a component of the helium gas during
atomization. An oxide coating on the L1.sub.2 aluminum powder is
beneficial for two reasons. First, the coating prevents
agglomeration by contact sintering and secondly, the coating
inhibits the chance of explosion of the powder. A controlled amount
of oxygen is important in order to provide good ductility and
fracture toughness in the final consolidated material. Hydrogen
content in the powder is controlled by ensuring the dew point of
the helium gas is low. A dew point of about minus 50.degree. F.
(minus 45.5.degree. C.) to minus 100.degree. F. (minus 73.3.degree.
C.) is preferred.
[0111] In preparation for final processing, the powder is
classified according to size by sieving. To prepare the powder for
sieving, if the powder has zero percent oxygen content, the powder
may be exposed to nitrogen gas which passivates the powder surface
and prevents agglomeration. Finer powder sizes result in improved
mechanical properties of the end product. While minus 325 mesh
(about 45 microns) powder can be used, minus 450 mesh (about 30
microns) powder is a preferred size in order to provide good
mechanical properties in the end product. During the atomization
process, powder is collected in collection chambers in order to
prevent oxidation of the powder. Collection chambers are used at
the bottom of atomization chamber 104 as well as at the bottom of
cyclone collector 116. The powder is transported and stored in the
collection chambers also. Collection chambers are maintained under
positive pressure with nitrogen gas which prevents oxidation of the
powder.
[0112] A schematic of the L1.sub.2 aluminum powder manufacturing
process is shown in FIG. 9. In the process aluminum 200 and L12
forming (and other alloying) elements 210 are melted in furnace 220
to a predetermined superheat temperature under vacuum or inert
atmosphere. Preferred charge for furnace 220 is prealloyed aluminum
200 and L1.sub.2 and other alloying elements before charging
furnace 220. Melt 230 is then passed through nozzle 240 where it is
impacted by pressurized gas stream 250. Gas stream 250 is an inert
gas such as nitrogen, argon or helium, preferably helium. Melt 230
can flow through nozzle 240 under gravity or under pressure.
Gravity flow is preferred for the inventive process disclosed
herein. Preferred pressures for pressurized gas stream 250 are
about 50 psi (10.35 MPa) to about 750 psi (5.17 MPa) depending on
the alloy.
[0113] The atomization process creates molten droplets 260 which
rapidly solidify as they travel through agglomeration chamber 270
forming spherical powder particles 280. The molten droplets
transfer heat to the atomizing gas by convention. The role of the
atomizing gas is two fold: one is to disintegrate the molten metal
stream into fine droplets by transferring kinetic energy from the
gas to the melt stream and the other is to extract heat from the
molten droplets to rapidly solidify them into spherical powder. The
solidification time and cooling rate vary with droplet size. Larger
droplets take longer to solidify and their resulting cooling rate
is lower. On the other hand, the atomizing gas will extract heat
efficiently from smaller droplets resulting in a higher cooling
rate. Finer powder size is therefore preferred as higher cooling
rates provide finer microstructures and higher mechanical
properties in the end product. Higher cooling rates lead to finer
cellular microstructures which are preferred for higher mechanical
properties. Finer cellular microstructures result in finer grain
sizes in consolidated product. Finer grain size provides higher
yield strength of the material through the Hall-Petch strengthening
model.
[0114] Key process variables for gas atomization include superheat
temperature, nozzle diameter, helium content and dew point of the
gas, and metal flow rate. Superheat temperatures of from about
150.degree. F. (66.degree. C.) to 200.degree. F. (93.degree. C.)
are preferred. Nozzle diameters of about 0.07 in. (1.8 mm) to 0.12
in. (3.0 mm) are preferred depending on the alloy. The gas stream
used herein was a helium nitrogen mixture containing 74 to 87 vol.
% helium. The metal flow rate ranged from about 0.8 lb/min (0.36
kg/min) to 4.0 lb/min (1.81 kg/min). The oxygen content of the
L1.sub.2 aluminum alloy powders was observed to consistently
decrease as a run progressed. This is suggested to be the result of
the oxygen gettering capability of the aluminum powder in a closed
system. The dew point of the gas was controlled to minimize
hydrogen content of the powder. Dew points in the gases used in the
examples ranged from -10.degree. F. (-23.degree. C.) to
-110.degree. F. (-79.degree. C.).
[0115] The powder is then classified by sieving process 290 to
create classified powder 300. Sieving of powder is performed under
an inert environment to minimize oxygen and hydrogen pickup from
the environment. While the yield of minus 450 mesh powder is
extremely high (95%), there are always larger particle sizes,
flakes and ligaments that are removed by the sieving. Sieving also
ensures a narrow size distribution and provides a more uniform
powder size. Sieving also ensures that flaw sizes cannot be greater
than minus 450 mesh which will be required for nondestructive
inspection of the final product.
[0116] Processing parameters of exemplary gas atomization runs are
listed in Table 1.
TABLE-US-00001 TABLE 1 Gas atomization parameters used for
producing powder Nozzle He Gas Dew Charge Average Metal Oxygen
Oxygen Diameter Content Pressure Point Temperature Flow Rate
Content Content Run (in) (vol %) (psi) (.degree. F.) (.degree. F.)
(lbs/min) (ppm) Start (ppm) End 1 0.10 79 190 <-58 2200 2.8 340
35 2 0.10 83 192 -35 1635 0.8 772 27 3 0.09 78 190 -10 2230 1.4 297
<0.01 4 0.09 85 160 -38 1845 2.2 22 4.1 5 0.10 86 207 -88 1885
3.3 286 208 6 0.09 86 207 -92 1915 2.6 145 88
[0117] The role of powder quality is extremely important to produce
material with higher strength and ductility. Powder quality is
determined by powder size, shape, size distribution, oxygen
content, hydrogen content, and alloy chemistry. Over fifty gas
atomization runs were performed to produce the inventive powder
with finer powder size, finer size distribution, spherical shape,
and lower oxygen and hydrogen contents. Processing parameters of
some exemplary gas atomization runs are listed in Table 1. It is
suggested that the observed decrease in oxygen content is
attributed to oxygen gettering by the powder as the runs
progressed.
[0118] Inventive L1.sub.2 aluminum alloy powder was produced with
over 95% yield of minus 450 mesh (30 microns) which includes powder
from about 1 micron to about 30 microns. The average powder size
was about 10 microns to about 15 microns. As noted above, finer
powder size is preferred for higher mechanical properties. Finer
powders have finer cellular microstructures. As a result, finer
cell sizes lead to finer grain size by fragmentation and
coalescence of cells during powder consolidation. Finer grain sizes
produce higher yield strength through the Hall-Petch strengthening
model where yield strength varies inversely as the square root of
the grain size. It is preferred to use powder with an average
particle size of 10-15 microns. Powders with a powder size less
than 10-15 microns can be more challenging to handle due to the
larger surface area of the powder. Powders with sizes larger than
10-15 microns will result in larger cell sizes in the consolidated
product which, in turn, will lead to larger grain sizes and lower
yield strengths.
[0119] Powders with narrow size distributions are preferred.
Narrower powder size distributings produce product microstructures
with more uniform grain size. Spherical powder was produced to
provide higher apparent and tap densities which help in achieving
100% density in the consolidated product. Spherical shape is also
an indication of cleaner and lower oxygen content powder. Lower
oxygen and lower hydrogen contents are important in producing
material with high ductility and fracture toughness. Although it is
beneficial to maintain low oxygen and hydrogen content in powder to
achieve good mechanical properties, lower oxygen may interfere with
sieving due to self sintering. An oxygen content of about 25 ppm to
about 500 ppm is preferred to provide good ductility and fracture
toughness without any sieving issue. Lower hydrogen is also
preferred for improving ductility and fracture toughness. It is
preferred to have about 25-200 ppm of hydrogen in atomized powder
by controlling the dew point in the atomization chamber. Hydrogen
in the powder is further reduced by heating the powder in vacuum.
Lower hydrogen in final product is preferred to achieve good
ductility and fracture toughness.
[0120] A schematic of the L1.sub.2 aluminum powder consolidation
process is shown in FIG. 10. The starting material is sieved and
classified L1.sub.2 aluminum alloy powders (step 310). Blending
(step 320) 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 320) 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.
[0121] Following blending (step 320), the powders are transferred
to a can (step 330) where the powder is vacuum degassed (step 340)
at elevated temperatures. The can (step 330) is an aluminum
container having a cylindrical, rectangular or other configuration
with a central axis. Cylindrical configurations are preferred with
hydraulic extrusion presses. Vacuum degassing times can range from
about 0.5 hours to about 8 days. A temperature range of about
300.degree. F. (149.degree. C.) to about 900.degree. F.
(482.degree. C.) is 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.
[0122] Following vacuum degassing (step 340), the vacuum line is
crimped and welded shut (step 350). The powder is then fully
densified by blind die compaction or closed die forging as the
process is sometimes called (step 360). At this point the can may
be removed by machining (step 380) to form a useful billet (step
390).
[0123] A schematic showing blind die compaction (process 400) is
shown in FIGS. 11A and 11B. The equipment comprises base 410, die
420, ram 430, and means to apply pressure to ram 430 indicated by
arrow 450. Prior to compaction, billet 440 does not fill die cavity
460. After compaction, billet 445 completely fills the die cavity
and has taken the shape of die cavity 460. The die cavities can
have any shape provided they have a central symmetrical axis
parallel to arrow 450. Cylindrical shapes adopt well for extrusion
billets. Rectangular shapes are suitable shapes for rolling
preforms. Canned L1.sub.2 aluminum alloy powder preforms are easily
densified due to the large capacity of modern hydraulic
presses.
3. Tube Fabrication
[0124] Consolidated L1.sub.2 aluminum alloy powders can be
fabricated into tube form by a number of processes. The processes
described here are only descriptive of general practice and are not
to be taken as exclusive.
[0125] FIG. 12 is a schematic illustration of tube rolling process
500. The starting workpiece is rolled sheet 510. Sheet 510 moves in
the direction of arrow A through a series of rolls comprising
progressively rounder cross sectional shapes that gradually
transform workpiece 510 into tube 590. Roll pair 515/520 initiates
the process by forming rolled sheet 510 into gradually curved cross
section 525. Roll pairs 530/535 and 545/550 produce profiles 540
and 555 wherein a tube cross section is starting to form. Vertical
roll pair 560/565 in combination with horizontal finishing roll
pair 570/575 produce tube 590 having cross section 580. The seam of
tube 590 is then sealed by welding or other methods known in the
art to produce a finished product. Tube rolling can be performed at
ambient or elevated temperatures depending on the dimensions and
mechanical properties of starting sheet 510.
[0126] FIGS. 13A and 13B are schematic diagrams showing the
Mannesmann process for forming tubing from a solid round bar. FIG.
13A illustrates that when a round bar is radially compressed by
pressure P, a biaxial stress state is generated wherein a tensile
stress forms in the center of the bar in a direction perpendicular
to the compressive stress axes. If the bar is rotating under these
conditions of stress, a cavity will form at the center of the bar
to relieve the tensile stress. The Mannesmann process simply takes
advantage of the cavity initiation and uses a mandrel to pierce the
bar and fabricate a tube between working rolls. A perspective
sketch of Mannesmann process 600 is shown in FIG. 13B. Shaped rolls
620 and 625 are turning in the direction of arrows B and C on an
axis canted to the centerline of solid bar 610. Under these
conditions, the rolls impart forces to bar 610 that reduce its
cross section as well as propel it forward from right to left in
the FIG. towards mandrel 630. Mandrel 630 is held in place by a
long rod or by other techniques (not shown) in which the mandrel
remains in place without the rod. For aluminum alloys, the
Mannesmann process is typically used to form thicker wall
tubes.
[0127] FIG. 14 is a perspective sketch of pilgering process 700
used to make seamless tube or pipe. The process rapidly and
efficiently reduces the diameter of a tube and generates a product
with homogeneous characteristics because of the relatively large
amount of plastic deformation in each forming step. Pilgering
process 700 starts with tubular workpiece 710 positioned over
tapered mandrel 720. Ring dies 730 and 740 move back and forth as
indicated by double headed arrow D while workpiece 710 rotates as
indicated by arrow E. As the dies move back and forth, they reduce
the thickness of workpiece 710 in a manner similar to how a rolling
pin rolls out dough. In the middle of each reciprocal stroke of the
ring dies, workpiece 710 becomes unloaded and is advanced forward
by a predetermined increment.
[0128] FIGS. 15A, 15B and 15C illustrate deformation of workpiece
710 during a forward pass of ring dies 730 (not shown) and 740
rotating in the direction of arrow F. FIG. 15A is a longitudinal
cross section of workpiece 710, mandrel 720, and ring die 740. FIG.
15B is a radial cross section of ring dies 740, workpiece 710, and
mandrel 720 at the point of contact between ring die 740, workpiece
710, and mandrel 720 in FIG. 15A. At the point of contact workpiece
710 conforms to dies 730 and 740 and mandrel 720. FIG. 15A shows
the longitudinal cross section of workpiece 710 in the process of
being "ironed" as noted by the conformal fit between workpiece 710
and ring die 740. The radial cross section in FIG. 15C illustrates
the separation of workpiece 710 and mandrel 720 before being
contacted by die 730 and 740 during the completion of the forward
pass.
[0129] FIGS. 15D, 15E, and 15F illustrate a similar comparison of
cross sections during the reverse pass of the ring dies. Die 740 is
rolling in the direction indicated by arrow G. Cross section 15E at
the point of contact between workpiece 710 and ring die 740 shows
workpiece 710 conforming to die 740 and mandrel 720. Cross section
15F ahead of the point of contact shows workpiece 710 has not been
contacted by the dies at this point during the reverse pass of the
ring dies. Since workpiece 710 is continually rotating during
pilgering and the deformation is substantial, L1.sub.2 alloy tubes
with homogeneous mechanical properties are produced by this
process.
[0130] FIG. 16 is a schematic cross section of direct tube
extrusion process 800 to produce L1.sub.2 aluminum alloy tube. The
process comprises billet 810 and billet container 815. Billet
container 815 is preferably heated or situated in a furnace to
preferably maintain billet 810 at an elevated extrusion
temperature. Extrusion ram 820 contains extrusion mandrel 830 such
that extrusion mandrel 830 is slideable within extrusion ram 820.
During tube extrusion, extrusion ram 820 under pressure P moves in
direction of arrows P such that billet 810 is forced through a
cavity in die 840 defined by exterior cavity wall 845. Die backer
850 serves to maintain die 840 in position during extrusion.
L1.sub.2 aluminum alloy tube 860, formed during extrusion has inner
diameter defined by diameter of mandrel 830. Pressure P is
preferably hydraulic pressure due to the high capacity of hydraulic
presses available in the art.
[0131] 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 with
departing from the spirit and scope of the invention.
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