High strength, ductile, low density aluminum alloys and process for making same

The present invention provides a process for making high strength, high ductility, low density rapidly solidified aluminum-base alloys, consisting essentially of the formula Al.sub.bal Zr.sub.a Li.sub.b X.sub.c, wherein X is at least one element selected from the group consisting of Cu, Mg, Si, Sc, Ti, U, Hf, Be, Cr, V, Mn, Fe, Co and Ni, "a" ranges from about 0.2-0.6 wt %, "b" ranges from about 2.5-5 wt %, "c" ranges from about 0-5 wt % and the balance is aluminum. The alloy is given multiple aging treatments after being solutionized. The microstructure of the alloy is characterized by the precipitation of a composite phase in the aluminum matrix thereof.

DESCRIPTION 
1. Field of the Invention 
The invention relates to a process for making high strength, high 
ductility, low density rapidly solidified aluminum-based alloys and, in 
particular, to alloys that are characterized by a homogeneous distribution 
of composite precipitates in the aluminum matrix hereof. The 
microstructure is developed by a heat treatment method consisting of 
initial solutionizing treatment followed by multiple aging treatments. 
2. Background of the Invention 
There is a growing need for structural alloys with improved specific 
strength to achieve substantial weight savings in aerospace applications. 
Aluminum-lithium alloys offer the potential of meeting the weight savings 
due to the pronounced effects of lithium on the mechanical and physical 
properties of aluminum alloys. The addition of one weight percent lithium 
(3.5 atom percent) decreases the density by 3% and increases the elastic 
modulus by 6%, hence giving substantial increase in the specific modulus 
(E/P). Moreover, heat treatment of alloys results in the precipitation of 
a coherent, metastable phase, .delta. (Al.sub.3 Li) which offers 
considerable strengthening. Nevertheless, development and widespread 
application of the Al-Li alloy system have been impeded mainly due to its 
inherent brittleness. It has been shown that the poor toughness of alloys 
in the Al-Li system is due to brittle fracture along the grain or subgrain 
boundaries. The two dominant microstructural features responsible for 
their brittleness appear to be the precipitation of intermetallic phases 
along the grain and/or subgrain boundaries and the marked planar slip in 
the alloys, which create stress concentrations at the grain boundaries. 
The intergranular precipitates tend to embrittle the boundary, and 
simultaneously extract Li from the boundary region to form precipitate 
free zones which act as sites of strain localization. The planar slip is 
largely due to the shearable nature of .delta.' precipitates which result 
in decreased resistance to dislocation slip on planes containing the 
sheared .delta.' precipitates. 
Several metallurgical approaches have been undertaken to circumvent these 
problems. It has been found that the PFZ (precipitate free zone) and 
precipitate induced intergranular fracture can be reduced by controlling 
processing to avoid the intergranular precipitation of stable Al-Li, 
Al-Cu-Li, Al-Mg-Li phases. The problem of planar slip can be partly 
alleviated by promoting slip dispersion through the addition of dispersoid 
forming elements and the controlled co-precipitation of Al-Cu-Li, Al-Cu-Mg 
and/or Al-Li-Mg intermetallics. The dispersoid forming elements include 
Mn, Fe, Co, etc. The co-precipitation of Cu and/or Mg containing 
intermetallics appears to be relatively effective in dispersing the 
dislocation movement. However, the sluggish formation of these 
intermetallics requires the thermomechanical treatments involving (P. J. 
Gregson and M. M. Flower, Acta Metallurgica, vol. 33, pp. 527-537, 1985), 
or a high Cu content which adversely affects the density of alloys (B. van 
der Brandt, P. J. von den Brink, H. F. de Jong, L. Katgerman, and H. 
Kleinjan, in "Aluminum-Lithium Alloy II", Metallurgical Society of AIME, 
pp. 433-446, 1984). Moreover, the properties of alloys thus processed were 
less than satisfactory. 
Recently, a new approach has been suggested to modify the deformation 
behavior of Al-Li alloy system through the development of Zr modified 
.delta.' precipitate. This approach is based on the observation that the 
metastable Al.sub.3 Zr phase in the Al-Zr alloy system is highly resistant 
to dislocation shear and is of the same crystal structure (Ll.sub.2) as 
.delta.'. In this regard, attempts have been made to produce a ternary 
ordered composite Al.sub.3 (Li, Zr) phase in the aluminum matrix with an 
alloy of Al-2.34 Li-1.07Zr (F. W. Gayle and J. B. van der Sande, Scripta 
Metallurgica, vol. 18, pp. 473-478, 1984). However, the process for 
developing a homogeneous distribution of such phase has required the 
strict control of processing parameters during the thermomechanical 
processing, as well as prolonged solutionizing and/or aging treatments. 
From the practical point of view, this process is quite undesirable and 
may also result in undesirable microstructural features such as 
recrystallization and wide precipitate free zones. Moreover, the process 
cannot be effectively applied to low Zr (e.g., 0.2 wt % Zr) containing 
alloys which produce a small volume fraction of heterogeneously 
distributed coarse composite precipitates (P. L. Makin and B. Ralph, 
Journal of Materials Science, vol. 19, pp. 3835-3843, 1984; P.J. Gregson 
and H. M. Flower, Journal of Materials Science Letters, vol. 3, pp. 
829-834, 1984; P. L. Makin, D. J. Lloyd and W. M. Stobbs, Philosophical 
Magazine A, vol. 51, pp. L41-L47, 1985). 
Alternatively, whilst the process can be applied to high Zr (e.g. 1.0 wt % 
Zr) containing alloys which produce a large volume fraction of shear 
resistant composite precipitates (F. W. Gayle et al., U.S. Pat. No. 
4,747,884), the high Zr content also increases the density of the alloy. 
There remains a need in the art for an alloy and process wherein the 
characteristics of strength, toughness and ductility are combined with a 
lower density than has heretofore been achieved with extant zirconium 
content. 
Despite considerable efforts to develop low density aluminum alloys, 
conventional techniques, such as those discussed above, have been unable 
to provide low density aluminum alloys having the sought for combination 
of high strength, high ductility and low density. As a result, 
conventional aluminum-lithium alloy systems have not been entirely 
satisfactory for applications such as aircraft structural components, 
wherein high strength, high ductility and low density are required. 
SUMMARY OF THE INVENTION 
The present invention provides a process for making rapidly solidified 
aluminum-lithium alloys containing a high density of substantially 
uniformly distributed shear resistant dispersoids which markedly improve 
the strength and ductility thereof. The low density rapidly solidified 
aluminum-base alloys of the invention consist essentially of the formula 
Al.sub.bal Zr.sub.a Li.sub.b X.sub.c, wherein X is at least one element 
selected from the group consisting of Cu, Mg, Si, Sc, Ti, U, Hf, Cr, V, 
Mn, Fe, Co and Ni, "a" ranges from about 0.2-0.6 wt %, "b" ranges from 
about 2.5-5 wt %, "c" ranges from about 0-5 wt % and the balance is 
aluminum. The microstructure of these alloys is characterized by the 
precipitation of composite Al.sub.3 (Li, Zr) phase in the aluminum matrix 
thereof. This microstructure is developed in accordance with the process 
of the present invention by subjecting a rapidly solidified alloy having 
the formula delineated above to solutionizing treatment followed by 
multiple aging treatments. An improved process for making high strength, 
high ductility, low density aluminum-based alloy is thereby provides 
wherein the aluminum-based alloy produced has an improved combination of 
strength and ductility (at the same density). 
The high strength, high ductility, low density rapidly solidified 
aluminum-based alloy produced in accordance with the present invention has 
a controlled composite Al.sub.3 (Li, Zr) precipitate which, 
advantageously, offers a wide range of strength and ductility combinations 
.

DESCRIPTION OF THE PREFERRED EMBODIMENTS 
In general, the present invention relates to the process of making high 
strength, high ductility, and low density rapidly solidified Al-Li-Zr-X 
alloys. Alloys of the invention are produced by rapidly quenching and 
solidifying a melt of a desired composition at a rate of at least about 
10.sup.5 .degree. Cs.sup.-1 onto a moving, chilled casting surface. The 
casting surface may be, for example, the peripheral surface of a chill 
roll or the chill surface of an endless casting belt. Preferably, the 
casting surface moves at a speed of at least about 4000 ft.min.sup.-1 
(1220 m.min-.sup.-1) to provide a cast alloy strip approximately 30 to 75 
mm in thickness, which has been uniformly quenched at the desired quench 
rate. Such strip can be 4" or more in width, depending upon the casting 
method and apparatus employed. Suitable casting techniques include, for 
example, jet casting and planar flow casting through a slot-type orifice. 
The strip is cast in an inert atmosphere, such as argon atmosphere, and 
means are employed to deflect or otherwise disrupt the high speed boundary 
layer moving along with the high speed casting surface. The disruption of 
the boundary layer ensures that the cast strip maintains contact with the 
casting surface and is cooled at the required quench rate. Suitable 
disruption means include vacuum devices around the casting surface and 
mechanical devices that impede the boundary layer motion. Other rapid 
solidification techniques, such as melt atomization and quenching 
processes, can also be employed to produce the alloys of the invention in 
non-strip form, provided the technique produces a uniform quench rate of 
at least about 10.sup.5 .degree. Cs.sup.-1. 
Rapidly solidified alloys having the Al.sub.bal Zr.sub.a Li.sub.b X.sub.c 
composition described above have been processed into ribbons and then 
formed into particles by conventional comminution devices such as 
pulverizers, knife mills, rotating hammer mills and the like. Preferably, 
the comminuted powder particles have a size ranging from about -40 to 200 
mesh, US standard sieve size. 
The particles are placed in a vacuum of less than 10.sup.-4 torr 
(1.33.times.10.sup.-3 pa.) preferably less than 10.sup.-5 torr 
(1.33.times.10.sup.-3 Pa.), and then compacted by conventional powder 
metallurgy techniques. In addition, the particles are heated at a 
temperature ranging from about 300.degree. C. to 550.degree. C., 
preferably ranging from about 325.degree. C. to 450.degree. C., minimizing 
the growth or coarsening of the intermetallic phases therein. The heating 
of the powder particles preferably occurs during the compacting step. 
Suitable powder metallurgy techniques include direct powder extrusion by 
putting the powder in a can which has been evacuated and sealed under 
vacuum, vacuum hot compaction, blind die compaction in an extrusion or 
forging press, direct and indirect extrusion, conventional and impact 
forging, impact extrusion and combinations of the above. 
The strengthening process involves the use of multiple aging steps during 
heat treatment of the alloy following rapid solidification thereof. The 
alloy is characterized by a unique microstructure consisting essentially 
of "composite" Al.sub.3 (Li, Zr) precipitate in an aluminum matrix (FIG. 
1) due to the heat treatment as hereinafter described. The alloy may also 
contain other Li, Cu and/or Mg containing precipitates provided such 
precipitates do not significantly deteriorate the mechanical and physical 
properties of the alloy. 
The factors governing the properties of the Al-Li-Zr-X alloys are primarily 
its Li content and micro-structure and secondarily the residual alloying 
elements. The microstructure is determined largely by the composition and 
the final thermomechanical treatments such as extrusion, forging and/or 
heat treatment parameters. Normally, an alloy in the as processed 
condition (cast, extruded or forged) has large intermetallic particles. 
Further processing is required to develop certain microstructural features 
for certain characteristic properties. 
The alloy is given an initial solutionizing treatment, that is, heating at 
a temperature (T.sub.1) for a period of time sufficient to substantially 
dissolve most of the intermetallic particles present during the forging or 
extrusion process, followed by cooling to ambient temperature at a 
sufficiently high rate to retain alloying elements in said solution. 
Generally, the time at temperature T.sub.1, will be dependent on the 
composition of the alloy and the method of fabrication (e.g., ingot cast, 
powder metallurgy processed) and will typically range from about 0.1 to 10 
hours. The alloy is then reheated to an aging temperature, T.sub.2, for a 
period of time sufficient to activate the nucleation of composite Al.sub.3 
(Li, Zr) precipitates, and cooled to ambient temperature, followed by a 
second aging treatment at temperature, for a period of time, T.sub.3, for 
a period of time sufficient for the growth of the composite Al.sub.3 (Li, 
Zr) precipitate and a dissolution of .delta.' precipitate whose nucleation 
is not aided by Zr. The alloy at this point is characterized by a unique 
microstructure which consists essentially of composite Al.sub.3 (Li, Zr) 
precipitate. This composite Al.sub.3 (Li, Zr) precipitate is resistant to 
dislocation shear and quite effective in dispersing dislocation motion 
(FIG. 2). The result is that the alloy containing an optimum amount of 
composite Al.sub.3 (Li, Zr) precipitate deform by a homogeneous mode of 
deformation resulting in improved mechanical properties. FIG. 3(b) clearly 
shows the homogeneous mode of deformation in an alloy subjected to the 
process claimed in this invention, while FIG. 3(a) shows the severe planar 
slip observed in a conventionally processed alloy due to the shearing of 
.delta.' precipitates by dislocations (see FIG. 4). The combination of 
ductility with high strength is best achieved in accordance with the 
invention when the density of the shear resistant dispersoids ranges from 
about 10 to 60 percent by volume, and preferably from about 20-40 percent 
by volume. 
The optimum and preferred amount of composite Al.sub.3 (Li,Zr) precipitate 
thus described is accomplished through the claimed chemistry and 
processing steps which maintain low density. 
The exact temperature, T.sub.1, to which the alloy is heated in the 
solutionizing step is not critical as long as there is a dissolution of 
intermetallic particles at this temperature. The exact temperature, 
T.sub.2, in the first aging step where the nucleation of composite 
Al.sub.3 (Li, Zr) precipitate is promoted, depends upon the alloying 
elements present and upon the final aging step. The optimum temperature 
range for T.sub.2, is from about 100.degree. C. to 180.degree. C. The 
exact temperature, T.sub.3, whose range is from 120.degree. C. to 
200.degree. C., depends on the alloying elements present and mechanical 
properties desired. Generally, the times at temperatures T.sub.2 and 
T.sub.3 are different depending upon the composition of the alloy and the 
thermomechanical processing history, and will typically range from about 
0.1 to 100 hours. 
The following examples are presented to provide a more complete 
understanding of the invention. The specific techniques, conditions, 
materials, proportions and reported data set forth to illustrate the 
principles of the invention are exemplary and should not be construed as 
limiting the scope of the invention. 
EXAMPLE 1 
The ability of composite Al.sub.3 (Li, Zr) precipitates to modify the 
deformation behavior of rapidly solidified Al-Li-Zr alloys is illustrated 
as follows: 
FIG. 2 is a weak beam dark field transmission electron micrograph showing 
microstructure of a deformed alloy (Al-3.7Li-0.5Zr) which has been rapidly 
solidified, solutionized at 540.degree. C. for 4 hrs. and subsequently 
aged at 160.degree. C. for 4 hrs. followed by final aging at 180.degree. 
C. for 16 hrs. Such heat treatment promotes the precipitation of composite 
Al.sub.3 (Li, Zr) which is highly resistant to dislocation shear and is 
quite effective in dispersing the dislocation movement. 
FIG. 3(a) shows a bright field electron micrograph showing microstructure 
of a deformed alloy (Al-3.7Li-0.5Zr) which has not been given the claimed 
process. The alloy following rapid solidification had been aged for 16 
hrs. at 180.degree. C. after solutionizing at 540.degree. C. for 4 hrs. 
This alloy showed the pronounced planar slip which is the common 
deformation characteristic of brittle alloy. 
In contrast, FIG. 3(b) illustrates the beneficial effect of the claimed 
process on the deformation behavior of an alloy having the composition 
Al-3.7Li-0.5Zr. After rapid solidification and then solutionizing at 
540.degree. C. for 4 hrs., the alloy had been subjected to the double 
aging treatment of 160.degree. C. for 4 hrs. and 180.degree. C. for 16 
hrs. the deformation mode of this alloy is quite homogeneous indicating 
high ductility. 
EXAMPLE 2 
A rapidly solidified alloy having a composition of Al-3.1Li-2Cu-1Mg-0.5Zr 
was developed for medium strength applications as shown in Table I. The 
alloy was rapidly solidified and then solutionized at 540.degree. C. for 
2.5 hrs., quenched into water at about 20.degree. C. and given 
conventional single aging and the claimed double aging treatments. 
TABLE I 
______________________________________ 
Ultimate Elongation 
0.2% Yield 
Tensile to Failure 
Strength (MPa) 
Strength (MPa) 
(%) 
______________________________________ 
Aged at 190.degree. C. 
524 592 3.6 
for 16 hrs. 
Aged at 170.degree. C. 
530 606 6.1 
for 4 hrs. and 
190.degree. C. for 16 hrs. 
______________________________________ 
Conventional aging treatment (190.degree. C. for 16 hrs.) showed poor 
ductility (3.6%) due to the shearing of .delta.' precipitate (FIG. 4), 
while composite precipitate developed by double aging (FIG. 1) improve 
both strength and ductility (6.1% elongation). 
EXAMPLE 3 
A high strength Al-Li alloy was made to satisfy the requirements for high 
strength applications for aerospace structure. A rapidly solidified alloy 
having a composition of Al-3.2Li-2Cu-2Mg-0.5Zr was rapidly solidified then 
solutionized at 542.degree. C. for 4 hrs. As shown in Table II, 
conventional aging treatment (190.degree. C. for 16 hrs.) showed lower 
strength (yield strength of MPa) and ductility (3.6%). However, double 
aging of the alloy (160.degree. C. for 4 hrs. followed by 180.degree. C. 
for 16 hrs.) gave significantly higher strength (yield strength of 554 
MPa) and ductility (5.5%), which meets property requirements for high 
strength alloys needed for aerospace structural applications. 
TABLE II 
______________________________________ 
Ultimate Elongation 
0.2% Yield 
Tensile to Failure 
Strength (MPa) 
Strength (MPa) 
(%) 
______________________________________ 
Aged at 190.degree. C. 
521 595 3.6 
for 16 hrs. 
Aged at 170.degree. C. 
554 631 5.5 
for 4 hrs. and 
190.degree. C. for 16 hrs. 
______________________________________ 
EXAMPLE 4 
This example illustrates the beneficial effect of the claimed process on 
the mechanical properties of a simple ternary alloy Al-3.7Li-0.5Zr. The 
rapidly solidified alloy was rapidly solidified, solutionized at 
540.degree. C. for 4 hrs., and subsequently aged as shown in Table III. 
The resulting tensile properties show that the claimed process results in 
improved strength and ductility compared to the conventional process. 
TABLE III 
______________________________________ 
Ultimate 
Aging 0.2% Yield Tensile Elongation 
Treatment Strength (MPa) 
Strength (MPa) 
Failure (%) 
______________________________________ 
140.degree. C., 16 hr. 
424 442 4.2 
120.degree. C., 4 hr. + 
434 460 6.0 
140.degree. C., 16 hr. 
160.degree. C., 16 hr. 
419 431 3.2 
140.degree. C., 4 hr. + 
425 448 4.8 
160.degree. C., 16 hr. 
140.degree. C., 16 hr. + 
426 451 4.6 
160.degree. C., 16 hr. 
______________________________________ 
EXAMPLE 5 
A wide range of mechanical properties can be achieved by subjecting a 
rapidly solidified alloy to multiple aging conditions. For example, a 
triple aging treatment (120.degree. C., 4 hrs. +140.degree. C., 16 hrs. 
+160.degree. C., 4 hrs.) produced yield strength of 446 MPa and ultimate 
tensile strength of 464 MPa with 4.6% elongation. As a result, a variety 
of heat treatments of the rapidly solidified alloys according to the 
claims can be employed to produce alloys having a variety of mechanical 
properties. 
EXAMPLE 6 
This example illustrates the potential of the claimed process for the 
development of composite precipitate in low Zr containing rapidly 
solidified Al-Li alloys. FIG. 5 shows the dark field electron micrograph 
of a typical rapidly solidified alloy Al-3.2Li-3Cu-1.5Mg-0.2Zr which had 
been rapidly solidified, solutionized at 540.degree. C. for 4 hrs., 
reheated to 170.degree. C. for 4 hrs. followed by final aging at 
190.degree. C. for 16 hrs. The large volume fraction of composite Al.sub.3 
(Li, Zr) precipitate observed in such an alloy indicates that the claimed 
process is also quite effective in Al-Li alloys having low Zr content of 
0.2% 
EXAMPLE 7 
This example illustrates the potential of the claimed process for the 
development of composite precipitates in a rapidly solidified alloy as 
specified in Example 4. The specific strength of the alloy (UTS) can be 
compared with the conventional ageing process conducted on an alloy 
outside the scope of the invention with high Zr content. It is evident 
from the specific strength that alloys having Zr content within 0.2 to 06 
wt % range of the present invention produce an improved combination of 
high strength at low density. 
TABLE IV 
______________________________________ 
Specific 
Aging UTS Density Strength 
Alloy Treatment (MPa) p (gm/cm.sup.3) 
(UTS/p) 
______________________________________ 
Al-3.7 wt % 140.degree. C./ 
442 2.32 190.5 
Li-0.5 wt % Zr 
16 hrs 
Al-3.7 wt % 120.degree. C./ 
460 2.32 198.3 
Li-0.5 wt % Zr 
4 hrs + 
140.degree. C./ 
16 hrs 
Al-2.34 wt % 
190.degree. C./ 
479 2.45 195.5 
Li-1.07 wt % Zr 
2 hrs 
______________________________________ 
Having thus described the invention in rather full detail, it will be 
understood that such detail need not be strictly adhered to but that 
further changes and modifications may suggest themselves to one skilled in 
the art, all falling within the scope of the present invention as defined 
by the subjoined claims.