U.S. patent application number 10/422234 was filed with the patent office on 2004-03-25 for nanophase precipitation strengthened al alloys processed through the amorphous state.
This patent application is currently assigned to QuesTek Innovations LLC. Invention is credited to Jou, Herng-Jeng, Olson, Gregory B., Qiu, Caian, Tang, Weijia.
Application Number | 20040055671 10/422234 |
Document ID | / |
Family ID | 29739724 |
Filed Date | 2004-03-25 |
United States Patent
Application |
20040055671 |
Kind Code |
A1 |
Olson, Gregory B. ; et
al. |
March 25, 2004 |
Nanophase precipitation strengthened Al alloys processed through
the amorphous state
Abstract
Aluminum alloys having improved strength characteristics at
elevated temperatures (300.degree. C.) are manufactured by
combining selected transition metals (Ni, Co, Ti, Fe, Y, Sc) and
selected rare earth materials (Er, Tm, Tb, Lu) in amounts of about
2 to 12% and 2 to 15% atomic percent respectively in an amorphous,
glassy state and subsequently devitrifying the amorphous material
to form a crystalline mix of fcc and L1.sub.2 phase material.
Devitrification from the amorphous state may be effected by various
means including thermal and thermo mechanical processes.
Inventors: |
Olson, Gregory B.;
(Riverwoods, IL) ; Tang, Weijia; (Wilmette,
IL) ; Qiu, Caian; (Wilmette, IL) ; Jou,
Herng-Jeng; (Wilmette, IL) |
Correspondence
Address: |
BANNER & WITCOFF, LTD.
TEN SOUTH WACKER DRIVE
SUITE 3000
CHICAGO
IL
60606
US
|
Assignee: |
QuesTek Innovations LLC
1820 Ridge Avenue
Evanston
IL
60201
|
Family ID: |
29739724 |
Appl. No.: |
10/422234 |
Filed: |
April 24, 2003 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
60375940 |
Apr 24, 2002 |
|
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60450114 |
Feb 25, 2003 |
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Current U.S.
Class: |
148/561 ;
148/403 |
Current CPC
Class: |
C22F 1/04 20130101; C22C
45/08 20130101; C22C 21/00 20130101 |
Class at
Publication: |
148/561 ;
148/403 |
International
Class: |
C22C 045/08 |
Goverment Interests
[0001] Activities relating to the development of the subject matter
of this invention were funded at least in part by United States
Government, U.S. Army Aviation & Missile Command Contract No.
DAAH01-02-C-R125, and thus may be subject to license rights and
other rights in the United States.
Claims
What is claimed is:
1. An aluminum alloy characterized by high strength including high
strength in the temperature range greater than about 250.degree. C.
comprising, in combination: an alloy mixture in primarily
crystalline form having at least about 30% by volume fcc phase and
at least about 10% by volume L1.sub.2 ductile precipitate phase,
said alloy consisting essentially of at least one transition metal
selected from the group consisting of about 2 to 12 atomic percent
Ni, Co, Ti, Fe, Y, Sc, Cu, Zn, V, Cr, Mn and Li, and at least one
rare earth material selected from the group consisting of about 2
to 15 atomic percent Er, Tm, Yb, Lu, and the balance Al.
2. The alloy of claim 1 further including Mg in an amount up to
about 5% atomic fraction.
3. The alloy of claim 1 having a tensile strength of at least about
275 Mpa at 250.degree. C.
4. A method for manufacture of an aluminum alloy characterized by
high strength including high strength in a temperature range
greater than about 250.degree. C. comprising the steps of: (a)
forming a melt mixture of aluminum, about 2 to about 12 atomic
percent of at least one metal selected from the group consisting of
Ni, Co, Ti, Fe, Y, Sc, Cu, Zn, V, Cr, Mn and Li, and at least one
rare earth selected from the group consisting of about 2 to 15
atomic percent Er, Tm, Yb, and Lu; (b) converting said melt to an
amorphous state wherein the mixture comprises at least about 70% by
volume amorphous material; and (c) devitrifying said amorphous
material to at least in part a crystalline matrix of at least about
30% by volume fcc and at least about 10% by volume L1.sub.2
phase.
5. The method of claim 4 wherein the step of converting the melt
comprises cooling in the range of at least about 10.sup.3.degree.
C. per second.
6. The method of claim 4 wherein the step of devitrification is
selected from a group consisting of thermoprocessing,
thermomechanical processing and combinations thereof. A high
temperature and high strength aluminum-based alloy processed
through a primarily greater than about 70% in volume amorphous
state and then devitrified into greater than about 70% in volume
crystalline microstructure with fee matrix and strengthening by
about a ductile L1.sub.2 precipitate phase greater than about 10%
L1.sub.2, and fcc greater than about 30% in volume.
7. The method of claim 4 wherein converting the melt comprises a
step selected from the group consisting of gas powder atomization,
water powder atomization, melt spinning, spray casting and
combinations thereof.
8. The method of claim 4 wherein the step of devitrifying amorphous
material comprises a step selected from the group consisting of hot
isostatic pressing, thermal aging, extrusion and combinations
thereof.
9. An aluminum alloy characterized by high strength including high
strength at a temperature greater than about 250.degree. C. made by
a process comprising the steps of: (a) formulating a melt comprised
of Al; at least one transition metal (TM) selected from the group
consisting of Ni, Co, Ti, Fe, Y, Sc, Cu, Zn, V, Cr, Mn, Li and Mg;
and at least one rare earth selected from the group consisting of
Er, Tm, Yb, Lu; (b) converting the melt to at least about 70% by
volume amorphous material; and (c) devitrifying at least in part
the amorphous material to a mixture of fee and L1.sub.2 crystalline
precipitate phase material.
10. The alloy product by the process of claim 9 wherein the
transition metal is provided in an amount of about 2 to 10 atomic
percent.
11. The alloy product by the process of claim 9 wherein the rare
earth material is provided in an amount of about 2 to 10 atomic
percent.
12. The alloy product by the process of claim 9 further including
an additive of Mg to the melt.
13. The alloy product by the process of claim 9 wherein
devitrifying the amorphous material comprises forming at least
about 30% by volume fcc material.
14. The alloy product by the process of claim 9 wherein
devitrifying the amorphous material comprises forming at least 10%
by volume L1.sub.2 material.
15. The alloy product by the process of claim 9 wherein converting
the melt to amorphous material comprises at least one step selected
from the group consisting of gas powder atomization, water powder
atomization and melt spinning.
16. The alloy product by the process of claim 9 wherein
devitrification comprises at least one step selected from the group
consisting of hot isostatic pressing, thermal aging, and
extrusion.
17. The method of claim 4 wherein converting the melt comprises the
step of rapid solidification processing.
18. The product by the process of claim 9 wherein converting the
melt comprises rapid solidification processing.
19. The alloy of claim 1 or claim 9 wherein the precipitate phase
has a grain size diameter in the range of less than about 80 nm.
Description
BACKGROUND OF THE INVENTION
[0002] In a principal aspect, the present invention relates to
aluminum-based alloys processed through an amorphous state or
stage, preferably by means of a rapid solidification process from
molten alloy, and then devitrified to a primarily crystalline
structure by thermal or thermomechanical processing. In order to
achieve a desired microstructure during processing, the aluminum
alloy comprises a combination of rare earth and transition metal
components with the aluminum. The final crystalline microstructure
has stable high temperature strength (at or above about 300.degree.
C.) characterized by nanoscale intermetallic precipitates,
preferably L1.sub.2 phase precipitates.
[0003] Currently available commercial aluminum alloys, either
manufactured with ingot or powder processing, are not capable of
simultaneously achieving high strength and high temperature
stability near 300.degree. C.; such characteristics being
particularly important in applications such as the fan components
in turbine engines where short-term tensile strengths of the order
of 500 MPa (40-60 ksi) are desired. Precipitation hardening
introduced by aging is a known method to strengthen aluminum
alloys. Thus, conventional high strength aluminum alloys in
commercial applications utilize GP (Guinier-Preston) zones and
subsequent precipitation when seeking to achieve high material
strength. With this type of strengthening, precipitates are usually
formed during aging at or below 250.degree. C. in order to produce
appropriate strengthening precipitate dispersions.
[0004] Examples of aluminum alloys processed with relatively high
aging temperatures in commercial practice include 2618 (200.degree.
C. for 20 hours), 4032 (170-175.degree. C. for 8-12 hours), and
2218 (240.degree. C. for 6 hours) [Metals Handbook--Properties and
Selection: Nonferrous Alloys and Special-Purpose Materials, Volume
2, 10.sup.th Edition, ASM International]. At the noted aging
temperatures, these alloys have improved microstructure stability
relative to other commercial aluminum alloys. These aluminum
alloys, when age-hardened, usually possess a room temperature yield
strength of about 600 Mpa (85 ksi). However, at temperatures
approaching 300.degree. C., the precipitation strengthening
efficiency in these alloys quickly and significantly degrades as a
result of precipitate coarsening and/or dissolution. Due to the
unstable strengthening precipitate size distribution at such high
temperatures, the yield strength of currently available aluminum
alloys at 300.degree. C. is often only 10% of the yield strength at
room temperature, and thus renders such alloys unsuitable for high
temperature applications above 150.degree. C.
[0005] In order to achieve a combination of high strength and
usable high temperature properties in aluminum alloys, researchers
have investigated a variety of intermetallic precipitation
dispersions. Since selected intermetallic precipitates may
determine the strength and ductility of the final aluminum alloy,
it is desirable to select cubic precipitate phases with high
crystallographic symmetry to promote ductility. The aluminum-based
L1.sub.2 phase (fcc ordered) is one of the best-known precipitates
to achieve a good combination of high strength and high
toughness.
[0006] However, since crystalline aluminum has very limited
solubility for most intermetallics, including L1.sub.2 forming
components, it is difficult to produce fine intermetallic
dispersions through crystalline solid-state heat treatments that
utilize a step to dissolve the strengthening phase particles prior
to reprecipitation during aging. However, rapid solidification
processing (RSP) from the liquid state has been attempted or
utilized. With an RSP process, it is possible to either (1)
directly produce fine crystalline structure, or (2) produce
partially or completely amorphous aluminum alloys. Nonetheless,
crystalline RSP aluminum alloys have not been able to meet the high
temperature strength requirements due to the difficulty of
producing small, stable particle sizes at adequate volume fraction.
Alternatively, the focus on amorphous aluminum alloys has been
primarily on the glass formability. These glassy alloys in general
have low ductility, and, when devitrified, they typically form
crystalline microstructures that are far from optimal for fracture
toughness and other requirements.
[0007] For example, RSP crystalline aluminum alloys, not designed
to form glass with RSP, utilize a metastable and coherent L1.sub.2
dispersion (about 1%) which converts into other phase structures at
high temperature. [Parameswaran, Weertman and Fine, Scripta
Metall., vol.23, pp.147, 1989] U.S. Pat. Nos. 4,950,452, 5,074,935,
5,334,266, 5,458,700, 5,431,751, 5,053,085 and its derived
divisional patents (U.S. Pat. Nos. 5,240,517, 5,320,688, and
5,368,658) utilize RSP to achieve either amorphous, a mixture of
amorphous and crystalline, or crystalline microstructures, without
the use of an L1.sub.2 strengthening dispersion. The aluminum
alloys of U.S. Pat. No. 4,889,582 achieve the high temperature
strength with RSP process and age hardening to produce Al.sub.3Fe
(which has non-cubic DO.sub.3, not L1.sub.2 structure)
strengthening precipitates, using different alloying components.
U.S. Pat. No. 4,464,199 describes high strength aluminum alloys
with 4 to 15% iron processed with RSP to achieve high temperature
stability. Use of iron promotes Al.sub.3Fe precipitation.
[0008] The aluminum alloys of U.S. Pat. No. 4,787,943 are composed
of transition metals and rare earth metal components. However, with
titanium and iron as the preferred transition metals, and
gadolinium and cerium as rare earth elements, the alloy does not
comprise a stable L1.sub.2 strengthening dispersion.
[0009] U.S. Pat. No. 6,248,453B1 describes the use of an L1.sub.2
strengthening dispersion for high strength aluminum alloys
utilizing certain alloying components, but does not utilize a
transition metal such as nicker and cobalt and an intermediate
amorphous state as a means to refine the L1.sub.2
microstructure.
[0010] Finally, the high strength and high temperature aluminum
alloys referenced have not found widespread application due the
lack of robust processing to produce those alloys and the
deficiency of other alloy properties, such as toughness and
ductility.
SUMMARY OF THE INVENTION
[0011] The present invention is directed to a new class of aluminum
alloys characterized by formation from an intermediate amorphous
state to a final face centered cubic (fcc) aluminum matrix
strengthened by a precipitate dispersion in order to achieve high
temperature strength with usable ductility. The disclosed aluminum
alloys thus are comprised of a fine crystalline structure of
primarily aluminum fcc matrix strengthened by a precipitate phase
with high cubic symmetry, preferably the L1.sub.2 phase. An
appropriate melt of aluminum with crystalline formers is first
processed to achieve an intermediate amorphous state to preserve
the crystalline formers. The preferred method to achieve a
primarily (above 70% in volume) amorphous state is a rapid
solidification process (RSP) from the molten alloy by process
techniques such as powder atomization, melt spinning, and spray
casting. The RSP process preferably should have a cooling rate of
at least about 10.sup.3.degree. C./sec, preferably at least
10.sup.4.degree. C./sec. Other methods to achieve amorphous
microstructure through a solid-state process, such as mechanical
milling, may also be used. The intermediate amorphous alloys are
then devitrified to a final primarily (above 70% in volume)
fcc/L1.sub.2 crystalline microstructure with at least about 30% fcc
phase in volume and at least about 10% L1.sub.2 phase in volume.
Thermal treatment, thermomechanical treatment, or a combination of
both, are the preferred methods to produce the final crystalline
microstructure.
[0012] The alloys of the invention demonstrate high strength
capability after both short and long exposure to a high temperature
at or above 300.degree. C. by maintaining a fine and stable
microstructure. For example, a short-term strength requirement of
40-60 ksi yield strength is expected. For the long-term strength,
as characterized by creep strength, it has been demonstrated that
optimal creep resistance at 300.degree. C. for aluminum alloys is
obtained from a dispersion of 25 nm diameter L1.sub.2 particles [E.
A. Marquis, Microstructural Evolution and Strengthening Mechanisms
in Al--Sc and Al--Mg--Sc Alloys, Ph.D. thesis, Northwestern
University, 2002.]. Also, it has been demonstrated [Ohtera, Inoue
and Masumoto, in First International Conference on Processing
Materials for Properties, Ed. Henein and Oki, The Minerals, Metals
& Materials Society, p.713, 1993] that aluminum-based alloys
devitrified by extrusion from amorphous atomized powder can achieve
aluminum fcc grain size of 80 to 170 nm, with intermetallic phase
dispersions as fine as 10 to 80 nm in diameter. Such grain or
particle sizes may therefore encompass the desired optimal particle
diameter for creep strengthening by an L1.sub.2 phase intermetallic
dispersion. Accordingly, a process and an aluminum alloy product
combining a sufficient amount of stable L1.sub.2 refined or
converted by devitrification from a glass phase material by
appropriate thermal or thermomechanical processing has been
discovered. The alloy meets or exceeds high temperature strength
requirements with optimized creep resistance as a result of
intermetallic, small particle phase dispersions.
[0013] The aluminum alloys of this invention are produced by
combining aluminum, selected metals, particularly transition
metals, rare earth elements, and other optional elements into a
melt which is converted to an amorphous or glassy state and then
subsequently to a crystalline state by a devitrification process.
In the broadest sense about 2 to 15 atomic percent in the aggregate
of at least one rare earth selected from the group consisting of
erbium (Er), thulium (Tm), ytterbium (Yb) and lutetium (Lu) is
provided for the melt. Further about 2 to 7 atomic percent in the
aggregate of at least one transition metal selected from the group
consisting of nickel (Ni), cobalt (Co), iron (Fe) and copper (Cu)
is provided to enhance or facilitate the process step of forming an
amorphous or glassy material from the melt. Other materials
included, optionally, in the melt comprise additional transition
metals, lithium (Li) and magnesium (Mg). For example, of less than
about 5% in the aggregate by atomic fraction scandium (Sc), yttrium
(Y), titanium (Ti), zirconium (Zr), vanadium (V), chromium (Cr),
manganese (Mn) and lithium (Li) incorporated into the melt or alloy
as a modifier. Less than about 7% in the aggregate by atomic
fraction magnesium (Mg), zinc (Zn), silver (Ag) and niobium (Nb)
may be incorporated into the melt or alloy to accommodate fcc
misfit in the crystalline state.
[0014] The selection of alloying components is based on the
function or criteria of (1) good L1.sub.2 formers, (2) adequate
glass formability with RSP--this leads to the selection of
components with strong short range ordering effects and slow
long-range diffusing kinetics in molten aluminum, and (3) good
fcc/L1.sub.2 coherency with low misfit which lead to a fine
microstructure and good coarsening resistance at high
temperature.
[0015] Lanthanide rare earth L1.sub.2 formers such as erbium (Er),
thulium (Tm), ytterbium (Yb), and lutetium (Lu) are slow diffusers
and hence enhance glass formability while at the same time allowing
L1.sub.2 formation during devitrification. Scandium (Sc) is a
strong L1.sub.2 former which can be used in conjunction with rare
earth components to promote L1.sub.2 formation and reduce the
misfit between fcc and L1.sub.2 phases. Transition metals such as
nickel (Ni), cobalt (Co), iron (Fe), and copper (Cu) can enhance
short range ordering in the molten aluminum, leading to the
increase of glass transition temperature and promoting glass
formability. Magnesium (Mg), as a component with relatively high
solubility in crystalline aluminum, can be used in fcc aluminum to
further reduce the fcc/L1.sub.2 misfit.
[0016] With efficient precipitation strengthening by the L1.sub.2
dispersion, the alloys of the subject invention achieve a hardness
of 490 VHN at significantly higher aging temperature compared to
conventional aluminum alloys indicating a usable strength at or
beyond 300.degree. C.
[0017] Thus, it is an object of the invention to provide a new
class of high-strength L1.sub.2 phase strengthened aluminum alloys
processed into the amorphous state, preferably with a rapid
solidification process, and then subsequently devitrified into a
crystalline structure with a thermomechanical or thermal
processes.
[0018] A further object of the invention is to provide aluminum
alloys with usable strength at or above 300.degree. C. by selecting
(1) slow diffusing L1.sub.2 formers to prevent fast coarsening
kinetics at high temperature, (2) appropriate matrix fcc and
L1.sub.2 components to reduce the lattice misfit between them to
further slow down the coarsening kinetics and to promote the
initial finer L1.sub.2 morphology.
[0019] Another object of the invention is to utilize the cubic
crystal structure of L1.sub.2 phase to promote the interfacial
coherency with an aluminum fcc matrix and the inherent ductile
behavior of the L1.sub.2 crystalline phase itself, due to high
number of slip systems in cubic L1.sub.2 structure, to maintain
reasonable overall alloy ductility.
[0020] A further object of the invention is to combine alloying
components so that the L1.sub.2 formation is enhanced while the
likelihood of forming other undesirable intermetallics is reduced,
in order to further promote alloy ductility of an aluminum
alloy.
[0021] Another object of the invention is to provide adequate glass
formability during the rapid solidification process such as powder
atomization or melt spinning to form an amorphous aluminum alloy
structure, preserving the L1.sub.2 formers in the solution before a
subsequent devitrification process.
[0022] A further object of the invention is to utilize the
intermediate, highly supersaturated amorphous state, slow L1.sub.2
coarsening kinetics, and low L1.sub.2/fcc misfit to produce a fine
microstructure in a strong aluminum alloy.
[0023] A further object of the invention is to provide aluminum
alloys having high strength at high temperatures with appropriate
toughness when processed with powder metallurgy techniques through
an amorphous state.
[0024] Another object of the invention is to combine transition
metals and rare earth components to provide sufficient glass
formability during a rapid solidification process and a nanoscale
L1.sub.2 dispersion strengthened aluminum alloy fcc crystalline
structure after devitrification.
[0025] These and other objects, advantages and features will be set
forth in the detailed description which follows.
BRIEF DESCRIPTION OF THE DRAWING
[0026] In the detailed description which follows, reference will be
made to the drawing comprised of the following figures:
[0027] FIG. 1 is a chart setting forth criteria for the design of
aluminum alloys of the invention;
[0028] FIG. 2 is a diagram illustrating the viscosity of liquid
aluminum nickel alloys calculated at 1700.degree. C. in comparison
with experimental data wherein the comparison is based upon
calculations made using different thermodynamic descriptions of the
aluminum nickel liquid (substitutional solution model without
associate or associate model with AlNi associate);
[0029] FIG. 3 is an assessed Al--Yb phase diagram in comparison
with experimental data;
[0030] FIG. 4a is a graph of the enthalpy of formation of cobalt
and a rare earth, i.e. Er and Yb predicted by quantum mechanical
calculation;
[0031] FIG. 4b is a graph of the mixing enthalpy of the materials
set forth in FIG. 4a;
[0032] FIG. 4c is a graph of the molar volume of the materials set
forth in FIG. 4a;
[0033] FIG. 5 is an X-ray diffractogram of the alloy of Example 1A
as melt-spun with positions of fcc pure aluminum reflections
indicating a fully amorphous state;
[0034] FIG. 6 is an isochronic DSC thermogram of the alloy of
Example 1A indicating devitrification steps upon heating;
[0035] FIG. 7 is an X-ray diffractogram of the alloy of Example 1A
after devitrification at 425.degree. C. for 22 hours, with
positions of reflections of pure fcc Al, Al.sub.3Yb, and Al.sub.3Er
phases, indicating the desired phases: FCC+L1.sub.2;
[0036] FIG. 8 is a SIMS elemental map of devitrified alloy of
Example 1A indicating Er and Yb partition together to L1.sub.2
phase, and NI and Co which do not partition to L1.sub.2;
[0037] FIG. 9 is a scanning electron microscope (SEM) secondary
electron image of devitrified alloy of Example 1A indicating
submicron morphology;
[0038] FIG. 10 is a 3DAP microscopy reconstruction of
Al.sub.3(Yb,Er)L1.sub.2 phase of the alloy of Example 1A after
devitrification at 424.degree. C. for 19 hours;
[0039] FIG. 11 is an X-ray diffractogram of the alloy of Example 2A
as melt-spun, with positions of fcc pure Al and L1.sub.2
reflections, indicating fully amorphous status;
[0040] FIG. 12 is an X-ray diffractogram of the alloy of Example 2A
after devitrification at 425.degree. C. for 19 hours, with
positions of reflections of pure fcc Al, Al.sub.3Yb, and Al.sub.3Er
phases, indicating the desired phases: FCC+L1.sub.2;
[0041] FIG. 13 is an X-ray diffractogram of the alloy of Example 2B
as melt-spun, with positions of fcc pure Al and L1.sub.2
reflections, indicating fully amorphous status;
[0042] FIG. 14 is an X-ray diffractogram of the alloy of Example 2B
after devitrification at 425.degree. C. for 19 hours, with
positions of reflections of pure fcc Al, Al.sub.3Yb, and Al.sub.3Er
phases, indicating the desired phases: FCC+L1.sub.2 wherein the
lattice parameter of L1.sub.2 is reduced by Sc;
[0043] FIG. 15 is an X-ray diffractogram of the alloy of Example 2C
as melt-spun, with positions of fcc pure Al and L1.sub.2
reflections, indicating almost fully amorphous status;
[0044] FIG. 16 is an X-ray diffractogram of the alloy of Example 2C
after devitrification at 425.degree. C. for 19 hours, with
positions of reflections of pure fcc Al, Al.sub.3Yb, and Al.sub.3Er
phases, indicating the desired phases: FCC+L1.sub.2 wherein the
lattice parameter of L1.sub.2 is expanded by Mg;
[0045] FIG. 17 is an X-ray diffractogram of the alloy of Example 2D
as melt-spun, with positions of fcc pure Al and L1.sub.2
reflections, indicating fully amorphous status;
[0046] FIG. 18 is an X-ray diffractogram of the alloy of Example 2D
after devitrification at 425.degree. C. for 19 hours, with
positions of reflections of pure fcc Al, Al.sub.3Yb, and Al.sub.3Er
phases, indicating the desired phases: FCC+L1.sub.2 and wherein the
lattice parameter of fcc is expanded by Mg and that of L1.sub.2 is
reduced by Sc;
[0047] FIG. 19 is an X-ray diffractogram of the alloy of Example 1C
as melt-spun, with positions of fcc pure Al and L1.sub.2
reflections, indicating non-amorphous status due to the absence of
transition metal elements;
[0048] FIG. 20 is an X-ray diffractogram of the alloy of Example 1C
after devitrification at 425.degree. C. for 19 hours, with
positions of pure fcc Al, Al.sub.3Yb, and Al.sub.3Er phases,
indicating elements ER, Yb, and Sc can be completely intersoluble
to form L1.sub.2 and wherein the lattice parameter of L1.sub.2 is
reduced by Sc; and
[0049] FIG. 21 is a scanning electron microscope (SEM) secondary
electron image of a devitrified alloy of Example 1C showing coarse
microstructure in comparison with the alloy of Example 1A (see FIG.
9) and indicating that passing through the glassy state can reduce
the particle size of precipitation.
DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0050] General Summary
[0051] In general, the subject matter of the invention comprises a
product (alloy) as well as a product by process and a process all
directed to an aluminum alloy in crystalline form having higher or
greater strength particularly at elevated temperatures, i.e.
greater than at least 250.degree. C. and preferably 300.degree. C.
The alloy of aluminum is made by compounding a mixture of aluminum
with selected rare earths (RE) and transition metals (TM) in an
amorphous or glassy state followed by devitrification of the
material from the amorphous state to a mixed crystalline state
comprised of face centered cubic (fcc) and L1.sub.2. (an ordered
fcc) state wherein the ratios of the crystalline states is within
certain preferred ranges. Preferably the resultant alloy has at
least about 30% by volume fcc phase and at least about 10% by
volume L1.sub.2 phase material with limited residual amorphous or
other phase material such as quasi-crystalline or ordered amorphous
material.
[0052] The choice of starting materials may be varied as may the
compounding processes, the glass phase formation processes and the
devitrification processes. Importantly, when in the amorphous
state, there may be some crystalline material contained therein,
but preferably no more than about 30% by volume in the fcc phase.
Thus forming the mixture in the amorphous or glassy intermediate
state and subsequent devitrification constitute very important
aspects of the invention.
[0053] The alloy materials, in addition to aluminum (Al), include
one or more lanthanide rare earths (RE) selected or taken from the
group of erbium (Er), thulium (Tm), ytterbium (Yb), and lutetium
(Lu), and one or more transition metals (TM) taken or selected from
the group of copper (Cu), nickel (Ni), cobalt (Co), titanium (Ti),
iron (Fe), yttrium (Y), and scandium (Sc). Further additives in
minor amounts (less than about five percent (5%) atomic percent)
may be included such as magnesium (Mg). The range of various
constituents is also important to maintain processability and
preserve transition characteristics. Thus RE materials are utilized
in the range of about 2 to 15 atomic percent (%), and the TM and
metals are utilized in the aggregate range of about 2 to 12 atomic
percent (%). Transition metals (TM) in the range of about 2 to 7
atomic percent are utilized.
[0054] The processes for mixing or forming the starting materials
in the amorphous state are not necessarily limiting. Thus, it is
contemplated that solid state processing, liquid or melt processing
as well as gas phase processing may be utilized, though liquid
phase processing is preferred based upon examples as reported
hereinafter. The completeness of the amorphous state is preferably
at least about 70% by volume and preferably greater. The amorphous
state is a solid noncrystalline state or phase prior to
devitrification.
[0055] Development Technique
[0056] Devitrification is accomplished via heat treating (thermal)
processing, thermo-mechanical processing as well as other
processing techniques.
[0057] Precipitation strengthened aluminum (Al) alloys are
difficult to develop for high strength due to limited solubility of
alloying elements in fcc Al. The low solubility of strengthening
dispersion elements prevents the solid-state solution treatment of
the alloy to bring strengthening components into solution. It is
hence not necessarily possible to allow subsequent aging to produce
a desirable strengthening microstructure. Al alloys with high
fractions of precipitate that cannot be completely solution treated
therefore have very coarse particle sizes that tend to limit
strength, corrosion resistance and toughness. In contrast, the Al
alloys exhibit composition and microstructure characteristics that
enable the Al alloys of invention to exhibit high strength, good
ductility, and high temperature stability at or above 300.degree.
C.
[0058] By selecting an appropriate composition, processing
techniques achieve a fully amorphous state after rapid cooling of
the Al alloy composition. Furthermore, with careful selection of
composition this glass can then be thermal or thermo-mechanically
processed such that the glass devitrifies into a fully crystalline
fcc matrix with intermetallic precipitates having a very fine size
scale. By passing through the glass state, the equilibrium
solidification that would produce coarse precipitates is avoided.
Furthermore, the selection of the precipitate phase as L1.sub.2
allows the precipitate to be coherent or semi-coherent with the fcc
matrix improving the ductility of the alloy and the resistance of
the precipitate to coarsen during subsequent thermal treatment.
[0059] FIG. 1 is a system flow-block diagram that illustrates the
processing/structure/properties relationships for alloys of the
invention. The two-stage nature of the chart emphasizes the
generally equal importance of processability and final performance
in the design of alloys of the invention. The processability is
governed by the glass forming ability; strength, ductility, and
thermal stability are the required end properties.
[0060] Certain transition metals (TM) (such as Fe, Co, Ni, Cu)
promote short-range ordering in liquid Al, which leads to low
partial molar volume, low thermal expansion, and high viscosity
that are beneficial to glass forming ability. Rare-earth (RE)
elements (such as Ce, Gd, Yb, Er) with large atomic size exhibit
low diffusivity in Al and thus retard crystal nucleation.
Therefore, Al--TM--RE comprise a class of glass forming system for
Al alloys.
[0061] Viscosity of liquid metals also has a strong effect on the
glass forming ability. Liquid phase metals having high viscosity
usually becomes glass through fast cooling. A binary model to
describe liquid's viscosity is set forth in Equation (1): 1 = - x 1
x 2 R T ( x 1 1 o + x 2 2 o ) 2 G m x 2 2 ( 1 )
[0062] The variables x.sub.i and G.sub.m represent the mole
fraction and molar Gibbs energy of the liquid, respectively. The
viscosity of pure components, .eta..sub.i.sup.o, is available from
literature, see Table 1. From the above equation it is seen that
the composition dependence of the viscosity is fully determined by
the thermodynamic properties of the liquid.
1TABLE 1 Viscosity and related data for pure liquid metals [L.
Battezzati and A. L. Greer, "The Viscosity of Liquid Metals and
Alloys," Acta metall. 37, 1791 (1989).] V.sub.m .eta.(T.sub.m)
C.sub.A T.sub.m (10.sup.-6 (10.sup.-3 [10.sup.-7 (J/K E B
.eta..sub.o Element (K) m.sup.3/mol) Pa s) mol.sup.1/3).sup.1/2]
(kJ/mol) (E/RT.sub.m) (10.sup.-3 Pa s) Al 933 11.3 1.30 1.30 16.5
2.13 0.149 Co 1765 7.59 4.18 1.58 44.4 3.03 0.255 Cr 2133 8.28 5.7
.about.2.2 .about.185 .about.10.43 1.7 .times. 10.sup.-4 Cu 1356
7.94 4.0 1.72 30.5 2.71 0.301 Fe 1809 7.96 5.5 2.18 41.4 2.75 0.370
Mn 1517 9.58 5 2.5 20-46.5 1.6-3.7 0.12-1.02 Ni 1728 7.43 4.90 1.85
50.2 3.49 0.166 V 2175 8.9 2.4 0.98 .about.73 .about.4.04
.about.0.042 Ti 1958 11.6 2.2 1.18 .about.68 .about.4.18
.about.0.034 Zr 2125 .about.15.7 3.5 1.58 .about.88 .about.4.98
.about.0.024 T.sub.m melting temperature V.sub.m molar volume at
T.sub.m .eta.(T.sub.m) viscosity at T.sub.m C.sub.A constant E
activation energy for viscous flow B B = E/RT.sub.m .eta..sub.o
pre-exponential viscosity
[0063] Equation (1) has been validated for several binary systems
where both experimental viscosity data and thermodynamic
descriptions are available. An example is the Al--Ni system where
an intermediate compound AlNi melts congruently at about
1650.degree. C. FIG. 2 shows the calculated viscosity in Al--Ni
liquid at 1700.degree. C. as a function of Ni content, in
comparison with experimental data. It is seen that the calculated
viscosity shows a similar trend as the experimental data, but the
latter indicates a stronger increase of the viscosity around 50 at.
% Ni. The strong increase in viscosity is believed to be caused by
molecular associates of Al--Ni in the liquid. In order to
illustrate the effect of the AlNi associates on the viscosity, the
Al--Ni system was thermodynamically reassessed with an associate
model for the liquid free energy. The viscosity of Al--Ni liquid
calculated using the AlNi associate model is also plotted in FIG. 2
for comparison. The calculation shows a significant increase in the
viscosity due to the formation of liquid AlNi associates.
[0064] This viscosity model has further been extended to
multi-component systems, which provides guidance of glass forming
ability for the design of alloys of the invention.
[0065] A model to predict glass transition temperature, T.sub.g,
has been developed for the Al--TM--RE system. The model is based on
the parabolic relationship between T.sub.g and room temperature
glass shear modulus, G, and it assumes that G is linear in atomic
fraction, X, of each element according to equation (2): 2 T g = (
14100 K i X i ) 0.49 i = Al , Mg , Ni , Co , Er , Yb , Sc , Y , etc
. ( 2 )
[0066] The coefficient for each element, K.sub.i, has been
calibrated to experimental DSC data of melt-spun amorphous ribbons.
These coefficients quantitatively represent the relative
effectiveness of the aforementioned molecular associates in
promoting glass forming ability by varying T.sub.g. This model also
provides guidance of glass forming ability for the design of alloys
of the invention. The prediction error of equation (2) compared to
available experimental data is about 9 K with a standard deviation
of 10 K, which is within the experimental precision.
[0067] The desired performance for applications (e.g. aircraft
structural components, high temperature turbine components, etc.)
determines a set of alloy properties required, including strength,
ductility and thermal stability resistance up to 300.degree. C. The
desired properties can be achieved through a fine, stable,
multicomponent L1.sub.2 phase dispersion in the Al matrix after
devitrification of the glassy Al, as long as a sufficient volume
fraction of L1.sub.2 and low misfit with the Al matrix and thus low
coarsening rate can be obtained.
[0068] The elements Er, Lu, Tm and Yb are known as the only stable
rare earth (RE) L1.sub.2 formers. The melting point and cost of
these rare earth elements are presented in Table 2. Some properties
of corresponding trialuminides are also given in Table 2, including
lattice parameter, density, and L1.sub.2 decomposition or congruent
melting temperature. Among the four RE elements, according to Table
2, Yb has the smallest lattice parameter and relatively lower cost.
Er has the lowest cost. Both Yb and Er have relatively lower
melting point. Usually, the lower the melting temperature, the
higher the glass forming ability. As a consequence, Alloys of the
invention utilize Yb and Er as preferred L1.sub.2 formers rather
than Tm and Lu.
2TABLE 2 Stable L1.sub.2 formers and their properties. Melting
Lattice L1.sub.2 Max. Element Point Cost.sup.a L1.sub.2 Parameter
Density Temp. Er 1529.degree. C. $725/kg Al.sub.3Er 4.226 .ANG.
5.507 g/cm.sup.3 1070.degree. C. Lu 1663.degree. C. $7500/kg
Al.sub.3Lu 4.186 .ANG. 5.783 g/cm.sup.3 1200.degree. C. Tm
1545.degree. C. $6500/kg Al.sub.3Tm 4.203 .ANG. 5.591 g/cm.sup.3 Yb
819.degree. C. $1600/kg Al.sub.3Yb 4.202 .ANG. 5.682 g/cm.sup.3
980.degree. C. Sc 1541.degree. C. $18,000/kg Al.sub.3Sc 4.101 .ANG.
3.03 g/cm.sup.3 1320.degree. C. .sup.aFor 1-5 kg cast metal ingots
from 99.9%-grade oxides
[0069] Sc is the only transition metal (TM) element that can form a
stable L1.sub.2 with Al. In comparison with RE L1.sub.2 formers, Sc
can form L1.sub.2 with a smaller lattice parameter of 4.101 .ANG.,
thus leading to smaller misfit between L1.sub.2 and the Al matrix.
However, Sc is by far the most expensive of the elements
considered, as seen in Table 2. The alloy of the invention
therefore seeks to limit the use of Sc as much as possible. Efforts
have been made to search for other transition metals to substitute
for Sc. A preliminary requirement for such substitution is
solubility. Experiments have shown that Ti has a substantial
solubility in Al.sub.3Sc, up to 12.5 at. %, i.e.
Al.sub.0.75Sc.sub.0.125Ti.sub.0.125- . In addition, Ti has the
lowest diffusion coefficient in solid Al among the transition
metals. For instance, the diffusivity in Al at 300.degree. C. is
9.0.times.10.sup.-20 m.sup.2/s for Sc, 6.3.times.10.sup.-24 m2/S
for Zr, and 2.7.times.10.sup.-25 m.sup.2/S for Ti. Adding Ti to
Al.sub.3Sc will thus reduce the coarsening rate of L1.sub.2
precipitates. Moreover, addition of Ti decreases the lattice
parameter of Al.sub.3(Sc,Ti) and hence minimizes the lattice misfit
with Al. The lattice parameters for seven hypothetical L1.sub.2
compounds with transition metals are calculated at zero degree
Kelvin (0 K) using the first-principle quantum mechanical approach.
Results are shown in Table 3. All seven hypothetic L1.sub.2
compounds show significant smaller lattice parameters than those
stable L1.sub.2 phases. Thus, alloys of the invention incorporate
Yb, Er and Sc as base L1.sub.2 formers but are not limited to these
elements. TMs such as Ti, V, Zr, etc, which will result in low
misfit and thus lower interfacial energy to yield a low driving
force for particle coarsening are considered useful.
3TABLE 3 Lattice Parameters of hypothetical L1.sub.2 TM Compounds
through First Principle Calculations at 0 K. TMs Hypothetic
L1.sub.2 Lattice Parameter, .ANG. Al.sub.3Ti 3.967 Al.sub.3Zr 4.085
Al.sub.3V 4.045 Al.sub.3Fe 3.8005 Al.sub.3Ni 3.8469 Al.sub.3Co
3.7946 Al.sub.3Cu 3.9363
[0070] It should be pointed out that the compound Al.sub.3Ti is
metastable as the L1.sub.2 phase but stable as D0.sub.22. It can be
transformed to the L1.sub.2 structure if Al is substituted
partially by Cr, Mn, Fe, Co or Ni.
[0071] For a multicomponent L1.sub.2 phase, the lattice parameter a
can be calculated based on a linear model similar to Vegard's law.
As an example, we have the following equation for L1.sub.2 with the
formula Al.sub.3(Er.sub.xYb.sub.ySc.sub.z) is utilized:
a[Al.sub.3(Er.sub.xYb.sub.ySc.sub.z]=x*a(Al.sub.3Er)+y*a(Al.sub.3Yb)+z*a(A-
l.sub.3Sc) (3)
[0072] Assuming L1.sub.2 has a linear thermal expansion, its
lattice parameter at elevated temperature can be calculated as
follows (T in .degree. C.):
a(L1.sub.2,T)=a(L1.sub.2,
RT).times.[1+(T-25).times..alpha.(L1.sub.2)] (4)
[0073] where .alpha.(L1.sub.2) represents the thermal expansion
estimated according to the following reciprocal relation: 3 ( M ) (
Al 3 M ) = T m ( A l 3 M ) T m ( M ) ( 5 )
[0074] where M represents the corresponding pure metal in Al.sub.3M
and T.sub.m is the melting point (.degree. K.). Table 4 summarizes
melting point, thermal expansion of relevant elements in alloys of
the invention and the estimated thermal expansion as well as
lattice parameter at 25.degree. C. and 300.degree. C. for stable
L1.sub.2 or meta-stable L1.sub.2.
4TABLE 4 Estimation of Thermal Expansion Coefficients and Lattice
Parameters. Thermal Thermal Melting Expansion, Expansion, Melting
Thermal Lattice Lattice Point, 1/.degree. K. 1/.degree. K. L1.sub.2
Point, Expansion, Parameter Parameter, Element .degree. K.
20.about.100.degree. C. 20.about.300.degree. C. Alloys .degree. K.
20.about.300.degree. C. .ANG., 25.degree. C. .ANG., 300.degree. C.
Ni 1728 13.3 .times. 10.sup.-6 14.4 .times. 10.sup.-6 Al.sub.3Ni
913 27.254 .times. 10.sup.-6 3.8639 3.8929 Co 1768 12.5 .times.
10.sup.-6 13.6 .times. 10.sup.-6 Al.sub.3Co 1408 17.077 .times.
10.sup.-6 3.8114 3.8293 Yb 1092 25 .times. 10.sup.-6 25 .times.
10.sup.-6 Al.sub.3Yb 1253 21.788 .times. 10.sup.-6 4.202 4.2272 Er
1802 9.2 .times. 10.sup.-6 9.2 .times. 10.sup.-6 Al.sub.3Er 1343
12.344 .times. 10.sup.-6 4.226 4.2403 Sc 1814 12 .times. 10.sup.-6
12 .times. 10.sup.-6 Al.sub.3Sc 1593 13.665 4.101 4.1164 Ti 1943
8.9 .times. 10.sup.-6 9.2 .times. 10.sup.-6 Al.sub.3Ti 1623 11.014
3.967 3.979 Y 1795 10.8 .times. 10.sup.-6 10.8 .times. 10.sup.-6
Al.sub.3Y 1253 15.472 4.299 4.3173 Al 25.3 .times. 10.sup.-6 4.0497
4.0779
[0075] In the practice of the invention, decrease of the misfit by
expanding the lattice parameter of the Al matrix is a useful
technique. There are few elements that have significant solubility
in Al. Mg is the only one that can substantially increase the
lattice spacing of Al. According to literature information, in
dilute AlMg alloys, every one atom percent of Mg will expand the Al
lattice by 0.0045 .ANG.. Further, every one atom percent of Mg will
increase the linear coefficient of thermal expansion of Al by
1.179.times.10.sup.-7/.degree. C.
[0076] Thermodynamic modeling and assessments have been performed
for Al--TM--RE systems with relevant elements incorporated in
alloys of the invention. The assessed Al--Yb binary phase diagram
is shown in FIG. 3. The L1.sub.2 Al.sub.3Yb is the strengthening
phase that is used in alloys of the invention. According to FIG. 3,
Al.sub.3Yb is not solution treatable. The only way to avoid coarse
precipitation from Al liquid is using fast cooling. It is possible
to get fine crystalline precipitates directly by rapid
solidification processing (RSP) from the melt. However, much finer
precipitates can be generated through devitrification of amorphous
state due to the larger driving force. In addition, the
precipitation process can be controlled through
devitrification.
[0077] Quantum mechanical calculations have been performed in
Al.sub.3Er--Al.sub.3Co and Al.sub.3Yb--Al.sub.3Co systems. The
predicted enthalpy of formation, mixing enthalpy and molar volume
in the corresponding systems are plotted in FIG. 4. According to
FIG. 4, there is a strong repulsive Co--RE interaction indicating
negligible Co solubility in L1.sub.2 Al.sub.3RE. This sets a
challenge in producing good glass forming alloys through
conventional technology in the same composition intended to
precipitate a significant amount of L1.sub.2 precipitation through
devitrification.
[0078] Experimental Results
[0079] The present invention alloys, through computational design
of multi-component Al--TM--RE systems incorporate, desired
processing properties--glass forming ability and the desired
microstructure--a fine dispersion of L1.sub.2 after devitrification
in the Al matrix. Alloys of the invention are summarized in Table 5
for their composition, status in melt-spun and devitrification
conditions, as well as the lattice parameter of the Al fcc and the
L1.sub.2 phases and the misfit between L1.sub.2 and Al matrix.
5TABLE 5 Summary of Composition and Microstructure of Al Alloys
Alloy Composition Main phases Lattice Parameter Misfit No. at. %
As-spun status after devitrfication L1.sub.2, a(.ANG.) fcc,
a(.ANG.) % 1A Al-2.5Ni-2.5Co- fully glass fcc, L1.sub.2 + 4.215
4.063 3.735 3.5Yb-3.5Er minor x phase 1B Al-2.75Ni-2.75Co-2.17Yb-
fully glass fcc, L1.sub.2 + Al.sub.3Y + 4.217 4.050 4.123
2.17Er-2.17Y minor x phase 1C Al-6Yb-6Er fcc, L1.sub.2 + fcc,
L1.sub.2 4.208 4.057 3.726 some glass 1D Al-4Yb-4Er-4Y fcc,
L1.sub.2 + fcc, L1.sub.2 + Al.sub.3Y 4.229 4.059 4.192 some glass
1E Al-4.5Yb-4.5Er-3Sc fcc, L1.sub.2 + fcc, L1.sub.2 4.183 4.060
3.032 some glass 1F Al-12Er-3Mg fcc, L1.sub.2 + fcc, L1.sub.2 4.216
4.059 3.865 some glass 1G Al-12Er-6Mg fcc, L1.sub.2 + fcc, L1.sub.2
4.217 4.065 3.739 some glass 1H Al-12Yb-3Mg fcc + fcc, L1.sub.2
4.216 4.076 3.435 some glass 1I Al-12Yb-6Mg fcc, L1.sub.2 + fcc,
L1.sub.2 4.191 4.061 3.201 some glass 2A Al-2Ni-2Co-4Yb-4Er almost
fully fcc, L1.sub.2 + 4.201 4.049 3.754 glass + minor fcc minor x
phase 2B Al-2.5Ni-2.5Co-2Yb- almost fully fcc, L1.sub.2 + 4.150
4.048 2.520 2Er-3Sc glass + minor fcc minor x phase 2C
Al-3Mg-2.5Ni-2.5Co- mostly glass + fcc, L1.sub.2 + 4.197 4.060
3.374 4.5Yb-2.5Er minor fcc minor x phase 2D Al-3Mg-2.5Ni-2.5Co-
fully glass fcc, L1.sub.2 + 4.155 4.057 2.416 2.5Yb-1.5Er-3Sc minor
x phase
[0080] According to Table 5, alloys with TM elements (e.g. Ni and
Co) are fully glass or almost fully glass in the as melt-spun
conditions. These alloys include alloy 1A, 1B, 2A, 2B, 2C and 2D.
Alloys without Ni and Co primarily show FCC Al and L1.sub.2
crystalline phases although they are partially amorphous in the as
melt-spun conditions. These alloys include alloy 1C, 1D, 1E, 1F,
1G, 1H and 1I. The results confirm the prediction that TM elements
such as Ni and Co are critical elements to promote the glass
forming ability in the alloys of the invention. Through
devitrification, all alloys precipitate mainly L1.sub.2 in the Al
matrix. It should be emphasized that the particle size of alloys
passing through a fully or almost fully glassy state is much finer
than that of alloys without passing through the glassy state or
only passing through a partially glassy state with L1.sub.2 already
present in the as-spun condition.
[0081] Following is a detailed demonstration of characteristics of
alloys in Table 5.
EXAMPLE 1
[0082] Alloy 1A in Table 5 was fabricated into both powders using
gas atomization and ribbons using melt spinning technology (wheel
speed 55 m/s). In the as melt-spun condition, Alloy 1A was
initially characterized by X-ray diffraction (XRD) and differential
scanning calorimetry (DSC). The diffractometer trace (Cu--K.alpha.
radiation ) of as-spun Alloy 1A in FIG. 5 shows a broad peak
characteristic of glass, and the ribbons show very good ductility,
also characteristic of the glassy state. A DSC trace run at
40.degree. C./min. shown in FIG. 6 indicates three stages of glass
devitrification at temperatures near 200.degree. C., 315.degree. C.
and 370.degree. C.
[0083] DSC thermograms indicated that for heating at a constant
rate of 10-40.degree. C./min, the dominant crystallization
reactions occur below about 400.degree. C. for all alloys under
investigation. Based on this information, a temperature of
425.degree. C. was selected for the devitrification anneal.
Specifically, alloy samples were encapsulated in quartz tubes under
a protective atmosphere of approximately 500 mbar Argon with
99.999% purity. The encapsulated specimens were kept at 425.degree.
C. for periods of 19 or 22 hrs, respectively, to ensure completion
of the crystallization reaction and grain growth to a size that
facilitates experimental characterization.
[0084] FIG. 7 shows an X-ray diffractogram of Alloy 1A after
devitrification at 425.degree. C. for 22 hours. At this point
transformations I-III are fully completed and the ribbons are fully
devitrified in the sense that no significant amount of amorphous
material can be detected by XRD and TEM. The peak positions of
(111), (200), and (220) reflections calculated for FCC pure Al
(lattice parameter 4.0496 .ANG.) are indicated in red. Within the
accuracy of this experiment (that is 0.01.degree. at
2.theta.=38.505.degree. (Al (111)), the measured peak positions are
identical to pure Al, indicating the equilibrium solubility of Ni,
Co, Yb, and Er in Al is very small. Also marked in red are the
positions of the `forbidden` reflections (110), (210), (211), and
(221) for FCC Al, which are superlattice reflections of the
L1.sub.2 structure. The peak positions for L1.sub.2 Al.sub.3Yb
(lattice parameter 4.202 .ANG. marked in green, and that for
L1.sub.2 Al.sub.3Er (lattice parameter 4.215 .ANG.) are marked in
brown. It can be seen that the measured peak positions are
perfectly matched by L1.sub.2 structure reflections. Therefore,
Alloy 1A mainly consists of FCC Al and L1.sub.2 phases as
devitrified microstructure. It is also noticed that there is a
minor third phase x, which crystalline structure is to be
identified.
[0085] The FCC Al+L1.sub.2 microstructure of Alloy 1A was
investigated by the secondary ion mass spectrometry (SIMS)
technology. FIG. 8 indicates that Yb and Er partition together into
L1.sub.2 precipitates, while Ni and Co have limited partitions.
[0086] The morphology of FCC+L1.sub.2 microstructure of Alloy 1A is
shown in FIG. 9. This SEM secondary electron image indicates that
the majority of L1.sub.2 precipitates are less than 1 micron after
devitrification at 425.degree. C., 22 hours. Based upon previously
discussed experience in devitrified Al alloys, this indicated the
potential to obtain nanoscale L1.sub.2 precipitates with
thermomechanical processing.
[0087] The three dimension atom probe microscopy (3DAPM) analysis
confirms that L1.sub.2 exists in Alloy 1A with formula
Al.sub.0.72(Yb,Er).sub.0.28- . Yb and Er have about the same site
fraction. Ni and Co are negligible in L1.sub.2 phase. See FIG.
10.
[0088] Alloy 1A was also studied through microhardness testing in
ribbon conditions after thermal cycles required for powder
consolidation and extrusion. Results are presented in Table 6.
Through devitrification at a temperature near 300.degree. C., Alloy
IA has demonstrated a hardness of 490 VHN, which is equivalent to
the room temperature tensile strength exceeding 1000 MPa according
to Inoue's work in similar alloy systems. (Akihisa Inoue,
"Amorphous, nanoquasicrystalline and nanocrystalline alloys in
Al-based systems", Progress in Materials Science 43 (1998)
365-520). This result indicates high efficient strengthening of
L1.sub.2.
6TABLE 6 Microhardness testing of Alloy 1A Condition Hardness (VHN)
As-spun 369 177.degree. C. 1 hr 276 177.degree. C. 10 hrs 310
177.degree. C. 20 hrs 335 177.degree. C. 20 hrs then 288.degree. C.
2 hrs. 490 177.degree. C. 20 hrs then 316.degree. C. 2 hrs. 340
177.degree. C. 20 hrs then 400.degree. C. 2 hrs. 275
EXAMPLE 2
[0089] Alloy 2A in Table 5 was fabricated into ribbons using melt
spinning technology (wheel speed 55 m/s). In the as melt-spun
condition, Alloy 2A was characterized by X-ray diffraction. The
diffractometer trace (Cu--K.alpha. radiation ) of as-spun Alloy 2A
in FIG. 11 shows a broad peak characteristic of glass, and the
ribbons show reasonable good ductility, also characteristic of the
glassy state.
[0090] FIG. 12 shows an X-ray diffractogram of Alloy 2A after
devitrification at 425.degree. C. for 19 hours. At this point the
ribbon is fully devitrified in the sense that no significant amount
of amorphous material can be detected by XRD. The peak positions of
(111), (200), and (220) reflections calculated for FCC pure Al
(lattice parameter 0.40496 .ANG.) are indicated in red. The
measured peak positions are identical to pure Al, indicating the
equilibrium solubility of Ni, Co, Yb, and Er in Al is very small.
Also marked in red are the positions of the `forbidden` reflections
(110), (210), (211), and (221) for FCC Al, which are superlattice
reflections of the L1.sub.2 structure. The peak positions for
L1.sub.2 Al.sub.3Yb (lattice parameter 4.202 .ANG.) are marked in
green, and that for L1.sub.2 Al.sub.3Er (lattice parameter 4.215
.ANG.) are marked in brown. It can be seen that the measured peak
positions are perfectly matched by L1.sub.2 structure reflections.
Therefore, Alloy 2A mainly consists of FCC Al and L1.sub.2 phases
as devitrified microstructure. It is noticed that the third phase
x, in Example 2A, is less than in Example 1A, reflecting the desire
to reduce the amount of x phase by reducing Ni and Co in Alloy 2A
while retaining reasonable glass forming ability.
EXAMPLE 3
[0091] Alloy 2B in Table 5 was fabricated into ribbons using melt
spinning technology (wheel speed 55 m/s). In the as melt-spun
condition, Alloy 2B was characterized by X-ray diffraction. The
diffractometer trace (Cu--K.alpha.radiation) of as-spun Alloy 2B in
FIG. 13 shows a broad peak characteristic of glass, and the ribbons
show reasonable good ductility, also characteristic of the glassy
state.
[0092] FIG. 14 shows an X-ray diffractogram of Alloy 2B after
devitrification at 425.degree. C. for 19 hours. At this point the
ribbon is fully devitrified in the sense that no significant amount
of amorphous material can be detected by X-ray diffraction (XRD).
The peak positions of (111), (200), and (220) reflections
calculated for FCC pure Al (lattice parameter 4.0496 .ANG.) are
indicated in red. The measured peak positions are identical to pure
Al, indicating the equilibrium solubility of Ni, Co, Yb, Er and Sc
in Al is very small. Also marked in red are the positions of the
`forbidden` reflections (110), (210), (211), and (221) for FCC Al,
which are superlattice reflections of the L1.sub.2 structure. It
can be seen that the measured peak positions are matched by
L1.sub.2 structure reflections with lattice parameter 4.150 .ANG..
It demonstrates that Yb, Er and Sc can be completely intersoluble
to form L1.sub.2, and the lattice parameter has been significantly
reduced by adding Sc, which is another means of reducing misfit to
keep the L1.sub.2 precipitates coarsening resistant. The misfit
between FCC Al and L1.sub.2 is 2.52% for Alloy 2B.
EXAMPLE 4
[0093] Alloy 2C in Table 5 was fabricated into ribbons using melt
spinning technology (wheel speed 55 m/s). In the as melt-spun
condition, Alloy 2C was characterized by X-ray diffraction. The
diffractometer trace (Cu--K.alpha. radiation ) of as-spun Alloy 2C
in FIG. 15 shows a broad peak characteristic of glass, and the
ribbons show reasonable good ductility, also characteristic of the
glassy state.
[0094] FIG. 16 shows an X-ray diffractogram of Alloy 2C after
devitrification at 425.degree. C. for 19 hours. At this point the
ribbon is fully devitrified in the sense that no significant amount
of amorphous material can be detected by XRD. The peak positions of
(111), (200), and (220) reflections calculated for FCC pure Al
(lattice parameter 4.0496 .ANG.) are indicated in red. The measured
peak positions indicate that the lattice space of Al has been
dilated to 4.060 .ANG., reflecting another feature of the
invention: adding Mg to expand the lattice space of Al and thus
decrease the misfit between Al and L1.sub.2. Also marked in red are
the positions of the `forbidden` reflections (110), (210), (211),
and (221) for FCC Al, which are superlattice reflections of the
L1.sub.2 structure. It can be seen that the measured peak positions
are perfectly matched by L1.sub.2 structure reflections. Therefore,
Alloy 2C mainly consists of FCC Al and L1.sub.2 phases as
devitrified microstructure.
EXAMPLE 5
[0095] Alloy 2D in Table 5 was fabricated into ribbons using melt
spinning technology (wheel speed 55 m/s). In the as melt-spun
condition, Alloy 2D was characterized by X-ray diffraction. The
diffractometer trace (Cu--K.alpha. radiation ) of as-spun Alloy 2D
in FIG. 17 shows a broad peak characteristic of glass, and the
ribbons show very good ductility, also characteristic of the glassy
state.
[0096] FIG. 18 shows an X-ray diffractogram of Alloy 2D after
devitrification at 425.degree. C. for 19 hours. At this point the
ribbon is fully devitrified in the sense that no significant amount
of amorphous material can be detected by XRD. The peak positions of
(111), (200), and (220) reflections calculated for FCC pure Al
(lattice parameter 4.0496 .ANG.) are indicated in red. The measured
peak positions indicate that the lattice space of Al has been
dilated to 4.057 .ANG.. Also marked in red are the positions of the
`forbidden` reflections (110), (210), (211), and (221) for FCC Al,
which are superlattice reflections of the L1.sub.2 structure. It
can be seen that the measured peak positions are matched by
L1.sub.2 structure reflections with lattice parameter 4.155 .ANG..
It demonstrates that Yb, Er and Sc can be completely intersoluble
to form L1.sub.2, and the lattice parameter has been significantly
reduced by adding Sc. Alloy 2D reflects the means of reducing
misfit by adding Mg to dilate the lattice space of Al and adding Sc
to decrease the lattice parameter of L1.sub.2. Thus Alloy 2D
achieved the lowest misfit so far:2.42% without losing the glass
forming ability. Alloy 2D mainly consists of FCC Al and L1.sub.2
phases as devitrified microstructure.
EXAMPLE6
[0097] Alloy 1E in Table 5 was fabricated into ribbons using melt
spinning technology (wheel speed 55 m/s). In the as melt-spun
condition, Alloy 1E was characterized by X-ray diffraction. The
diffractometer trace (Cu--K.alpha. radiation ) of as-spun Alloy 1E
in FIG. 19 shows crystalline reflection peaks, indicating
non-amorphous or partially non-amorphous state. It demonstrates
that TMs are necessary elements to promote glass in Al--TM--RE
system. In addition, the ribbons are very brittle, also
characteristic of the non-glassy state.
[0098] FIG. 20 shows an X-ray diffractogram of Alloy 1E after
devitrification at 425.degree. C. for 19 hours. This diffractogram
shows that Alloy 1E only has two phases: fcc Al+L1.sub.2. It also
demonstrates that Yb, Er and Sc can be completely intersoluble to
form L1.sub.2, and the lattice parameter has been significantly
reduced by adding Sc. The misfit between FCC Al and L1.sub.2 is
3.03%.
[0099] The morphology of fcc+L1.sub.2 microstructure of Alloy 1E is
shown in FIG. 21. This SEM secondary electron image indicates that
the particle size of L1.sub.2 precipitates is about 3 times larger
than that of Alloy 1A, see FIG. 9, although the misfit of Alloy 1A
is a little bit larger:3.735%. It indicates that devitrification of
glassy state can significantly reduce the particle size of
precipitation in comparison with that by crystallization directly
from fast cooling at the same rate.
[0100] In review, the range of composition processing parameters,
microstructure characterization and physical properties is
summarized in the following Table 7:
7 TABLE 7 Composition Transition Metals and Mg, Li Processing
Parameters Late FCC thermo- Rare TM for elem. thermal mechanical
Earth glass L12 For aging temp. hot Elem. forming modifier misfit
glass for thermal extrusion (1) (2) (3) (4) at. % RSP devitrifying
parameters Maximum Er, Lu, Ni, Co, Sc, Y Mg, Zn, 2%<sum(group
cooling >300.degree. C. Ranges Yb, Tm Fe, Cu Ti, Zr, Ag
(1))<15% rate U V, Cr, 2%<sum(group >10.sup.4.degree. C./
Mn, Li, (2))<7% sec Nb sum(group (3))<5% Preferred Er, Lu,
Ni, Co, Sc Mg 5%<sum(group cooling >325.degree. C.
>350.degree. C. Ranges Yb, Tm (1))<12% rate <475.degree.
C. <450.degree. C. 2%<sum(group >10.sup.4.degree. C./ 5,
9, 11:1 (2))<7% sec ratio Sc<4% Mg<6% Optimum Er, Yb Ni,
Co, Sc 7%<sum(Er, Yb, >10.sup.5.degree. C./ >370.degree.
C. 400.degree. C. Ranges Sc)<8% sec, <450.degree. C. 11:1
ratio 3%<sum(Ni, Co) atomized <6% powder with -325 mesh
Examples Er, Yb Ni, Co Sc, Y Mg 7%<sum(Er, Yb) about 425.degree.
C. 400.degree. C. <8% sum(Ni, Co) 10.sup.6.degree. C./ 11:1
ratio <5.5% sum(Sc, Y) sec <5%, Mg<3% Preferred Methods:
Atomization and consolidation with extrusion Properties high
temperature microstructure characterization (vol %) room UTS of
amorphous devitrified state temperature 275-410 fcc glass others
fcc L12 glass others UTS Mpa Maximum <15% bal. <10% bal.
>10% <15% <20% >500Mpa >250.degree. C. Ranges
Preferred <8% bal. <3% bal. >20%- <8% <15%
>800Mpa >275.degree. C. Ranges 25% Optimum <5% bal. <1%
bal. >30% <3% <5% >1000Mpa >300.degree. C. Ranges
Examples 100% or bal. >28% <3% <15- 490VHN <20% <32%
20% (1A) which is >1000MPa
[0101] Variations of the described aluminum alloy as well as the
process for manufacture thereof and the product created by the
process are available to provide the expected functionality of high
short-term and long-term strength at temperatures above about
300.degree. C. Thus the invention is to be limited only by the
following claims and equivalents thereof.
* * * * *