CA1280342C - High strength, ductile, low density, aluminum alloys and process for making the same - Google Patents
High strength, ductile, low density, aluminum alloys and process for making the sameInfo
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- CA1280342C CA1280342C CA000513291A CA513291A CA1280342C CA 1280342 C CA1280342 C CA 1280342C CA 000513291 A CA000513291 A CA 000513291A CA 513291 A CA513291 A CA 513291A CA 1280342 C CA1280342 C CA 1280342C
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- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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- C22C21/00—Alloys based on aluminium
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Abstract
ABSTRACT OF THE DISCLOSURE
The present invention provides a process for making high strength, high ductility, low density aluminum-base alloys, consisting essentially of the formula AlbalZraLibXc, 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 0.15-2 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.
The present invention provides a process for making high strength, high ductility, low density aluminum-base alloys, consisting essentially of the formula AlbalZraLibXc, 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 0.15-2 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
~280342 DESCRIPTION
IGH STRENGTH, DUCTILE, ',OW DENSITY AL~MINUM
ALLOYS AND PROCESS FOR MAKING SAME
5 1. Field of the Invention The invention relates to a process for making high strength, high ductility, low density aluminum-based alloys, and, in particular, to the alloys that are characteriæed by a homogeneous distribution of composite 10 precipitates in the aluminum matrix thereof. The microstructure is developed by heat treatment method consisting of initial solutionizing treatment followed by multiple aging treatments.
IGH STRENGTH, DUCTILE, ',OW DENSITY AL~MINUM
ALLOYS AND PROCESS FOR MAKING SAME
5 1. Field of the Invention The invention relates to a process for making high strength, high ductility, low density aluminum-based alloys, and, in particular, to the alloys that are characteriæed by a homogeneous distribution of composite 10 precipitates in the aluminum matrix thereof. The microstructure is developed by 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 20 mechanical and physical properties of aluminum alloys.
The addition of one weight percent lithium t ~3.5 atom percent) decreases the density by ~3~ and increases the elastic modulus by ~6~ , hence giving a substantial increase in the specific modulus ~/ p ). Moreover, heat treatment of alloys results in the precipitation of a coherent, metastable phase, ~' (A13Li) 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 micro-structural 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 inter-`.~
~2~)342 granular 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 ~' precipitates which result in decreased resistance to dislocation 51ip on plane~
containing the sheared ~' precipitates.
Several metallurgical approaches have been under-taken to circumvent these problems. It has been found lQ 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 15 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 20 f 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 stretching operations and multiple aging treatments ~P.J. Gregson and M.M. Flower, Acta Metallurgica, vol. 33, pp. 527-~37, 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 ~' precipitate.
This approach is based on the observation that the metastable A13Zr phase in the Al-Zr alloy system is highly resistant to dislocation shear and is of the same 3~
crystal structure (L12) as ~'. In this regard, attempts have been made to produce a ternary ordered composite A13(Li, zr) phase in the aluminum matrix with an alloy of Al-2.34 Li-1.07Zr (F.W. Gayle and J .B . vander Sande, 5 Scripta Metallurgica, vol. 18, pp. 473-478, 1984).
However, the process for developing a homogeneou~ dis-tributlon of such phase has required the strict control of processing parameters during the thermomechanical processing, as well as prolonged solutionizing and/or 10 aging treatments. From the practical point of view, this process is quite undesirable and may also result in undesirable microstructural features such as recrystal-lization and wide precipitate free zones. Moreover, the process cannot be effectively applied to low Zr ~e.g., 15 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 20 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).
Despite considerable efforts to develop low density aluminum alloys, conventional techniques, such as those 25 discussed above, have been unable to provide low density aluminium 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 aluminium-lithium alloys containing a high density of substantially uniformly distributed shear resistant dispersoids which markedly improve the strength and ductility thereof. The low density aluminum-base alloys, of the invention consist essentially of the ~28~33~L2 formula AlbalZraLibXc, 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 0.15-2 wt~, ~b" ranges from about 2.5-S wt~, ~c"
The addition of one weight percent lithium t ~3.5 atom percent) decreases the density by ~3~ and increases the elastic modulus by ~6~ , hence giving a substantial increase in the specific modulus ~/ p ). Moreover, heat treatment of alloys results in the precipitation of a coherent, metastable phase, ~' (A13Li) 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 micro-structural 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 inter-`.~
~2~)342 granular 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 ~' precipitates which result in decreased resistance to dislocation 51ip on plane~
containing the sheared ~' precipitates.
Several metallurgical approaches have been under-taken to circumvent these problems. It has been found lQ 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 15 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 20 f 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 stretching operations and multiple aging treatments ~P.J. Gregson and M.M. Flower, Acta Metallurgica, vol. 33, pp. 527-~37, 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 ~' precipitate.
This approach is based on the observation that the metastable A13Zr phase in the Al-Zr alloy system is highly resistant to dislocation shear and is of the same 3~
crystal structure (L12) as ~'. In this regard, attempts have been made to produce a ternary ordered composite A13(Li, zr) phase in the aluminum matrix with an alloy of Al-2.34 Li-1.07Zr (F.W. Gayle and J .B . vander Sande, 5 Scripta Metallurgica, vol. 18, pp. 473-478, 1984).
However, the process for developing a homogeneou~ dis-tributlon of such phase has required the strict control of processing parameters during the thermomechanical processing, as well as prolonged solutionizing and/or 10 aging treatments. From the practical point of view, this process is quite undesirable and may also result in undesirable microstructural features such as recrystal-lization and wide precipitate free zones. Moreover, the process cannot be effectively applied to low Zr ~e.g., 15 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 20 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).
Despite considerable efforts to develop low density aluminum alloys, conventional techniques, such as those 25 discussed above, have been unable to provide low density aluminium 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 aluminium-lithium alloys containing a high density of substantially uniformly distributed shear resistant dispersoids which markedly improve the strength and ductility thereof. The low density aluminum-base alloys, of the invention consist essentially of the ~28~33~L2 formula AlbalZraLibXc, 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 0.15-2 wt~, ~b" ranges from about 2.5-S wt~, ~c"
5 ranges from about 0-s wt% and the balance is aluminum.
The microstructure of these alloys is characterized by the precipitation of composite A13(Li, Zr) phase in the aluminum matrix thereof. This microstructure is developed in accordance with the process of the present 10 invention by subjecting an 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 provided wherein the 15 aluminum-based alloy produced has an improved combination of strength and ductility (at the same density).
The high strength, high ductility, low density aluminum-based alloy produced in accordance with the 20 present invention has a controlled composite A13(Li, Zr) precipitate which, advantageously, offers a wide range of strength and ductility combinations.
BRIEF DESCRIPTION OF THE DRAWINGS
The invention will be more fully understood and further advantages will become apparent when reference is made to the following detailed description and the accompanying drawings, in which:
Fig. 1 is a dark field transmission electron micro-~raph of an alloy having the composition Al-3.1Li-2Cu-lMg-0.5Zr~ the alloy having been subjected to double aging treatments (170C for 4 hrs. followe~ by 190C for 16 hrs.) to develop a composite precipitate in the aluminum matrix thereof;
Fig. 2 is a weak beam dark field micrograph of an alloy having the composition Al-3.7Li-O.SZr, illustrating the resistance of the composite precipitate to dislocation shear during deformation;
Fig. 3(a) shows the planar slip observed in an ~L28~3~2 alloy having the composition Al-3.7Li-0.5Zr, the alloy having been subjected to a conventional aging treatment (180C for 16 hours);
Fig. 3(b) shows the beneficial effect of subjecting 5 the alloy of Fig.3(a) to treatment in accordance with the claimed process tl60C for 4 hr5. followed by 180C
for 16-hrs.), thereby promoting the homogeneous deformation thereof;
Fig. 4 shows the sheared ~' precipitates observed 10 in an alloy having the composition Al-3.1Li-2Cu-lMg-0.5zr, the alloy havin~ been subjected to a conventional aging treatment (190C for 16 hours); and Fig. 5 shows the development of composite precipi-tates in an alloy having the composition Al-3.2Li-3Cu-15 1.5Mg-0.2Zr~ the alloy having been subjected to treatment in accordance with the claimed process tl70C
for 4 hrs. followed by 190C for 16 hrs).
DESCRIPTION OF THE PREFERRED EMBODIMENTS
In general, the present invention relates to the 2~ process of making high strength, high ductility, and low density Al-Li-Zr-X alloys. The process involves the use of multiple aging steps during heat treatment of the alloy. The alloy is characterized by a unique micro-structure consisting essentially of "compositen 25 A13(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~
zr-X alloys are primarily its I.i 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 extusion, forging and/or heat ~reatment parametersO Normally, an alloy in the as-processed condition (~ast, extruded or forged) has large . . .
~Z3~3~:
intermetallic particles. Further processing is required to develop certain microstructural features for certain characteristic properties.
The alloy is given an initial solutionizing treat-5 ment, that is, heating a~ a temperature (Tl) for a period of time sufficient to substantially dis501ve most of the intermetallic particles present during the forging or extrusion process, followed by cooling to ambient temperature at a sufficiently high rate to 10 retain alloying elements in said solution~ Generally, the time at temperature Tl, will be dependent on the composition of the alloy and the method of fabrication ~e.g., ingot cast, powder metallargy processed) and will typically range from about 0.1 to 10 hours. The alloy 15 is then reheated to an aging temperature, T2, for a period of time sufficient to activate the nucleation of composite A13(Li, Zr) precipitates, and cooled to ambient temperature, followed by a second aging treatment at temperature, T3, for a period of time 20 sufficient for the growth of the composite A13(Li, Zr) precipitate and a dissolution of ~' precipitate whose nucleation is not aided by Zr. The alloy at this p~int is characterized by a unique microstructure which con-sists essentially of composite A13(Li, zr) precipi-tate. This composite A13 (Li, Zr) precipitate is re-sistant to dislocation shear and quite effective in dispersing dislocation motion (Fig. 2). The result is that the alloy containing an optimum amount of composite A13(Li, Zr) precipitate deforms by a homogeneous mode of deformation resulting in improved mechanical properties. Fig. 3(b) clearly shows the h~mogeneous 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 ~' 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 exact temperature, Tl, to which the alloy is 5 heated in the solutionizing step is not critical as long as there is a dissolution of intermetallic particles at this temperature. The exact temperature, T2, in the first aging ~tep where the nucleation of composite A13 ~Li, Zr) precipitate is promoted, depends upon the 10 alloying elements present and upon the final aging step The optimum temperature range for T2, is ~rDm about 100C to 180C. The exact temperature, T3, whose range is from 120C to 200C, depends on the alloying elements present and mechanical properties desired.
15 Generally, the times at temperatures T2 and T3 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 ability of composite A13(Li, Zr) precipitates to modify the deformation behavior of 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 solutionized at 540C for 4 hrs. and subseguently aged at 160C for 4 hrs. followed by final aging at 180C for 16 hrs. Such heat treatment promotes the precipitation of composite A13(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-O.SZr) which has not been given the claimed process~
The alloy had been aged, for 16 hrs. at 180QC after solutionizing at 540 for 4 hrs. This alloy showed the pronounced planar slip which is the common deformation characteristic of brittle alloy.
34~2 In contrast, Fig. 3(b) illustrates the beneficial effect of the claimed process on the deformation benavior of an alloy having the composition Al-3.7Li-0.5Zr. After solutionizing at 540C for 4 hrs., the 5 alloy had been subjected to the double aging treatment of 160C for 4 hrs. and 180C for 16 hrs. The deform~tion mode of this alloy is quite homogeneous indicating high ductility.
Example 2 An alloy having a composition of Al-3.1Li-2Cu-lMg-0.5Zr was developed for medium strength applications as shown in Table I. The alloy was solutionized at 540C
for 2.5 hrs./ quenched into water at about 20C and given conventional single aging and the claimed double 15 aging treatments.
~28~3~2 TABLE I
0.2~ Yield Ultimate Tensile Elongation to Strength (MPa) Strength (MPa) Failure (~) Aged at 190C
for 16 hrs. 524 592 3.6 Aged at 170C --for 4 hrs. and 10190C for 16 hrs. 530 606 6.1 Conventional aginy treatment (19ODC for 16 hrs.) showed poor ductility t3.6%) due to the shearing of ~' precipitate (Fig. 4), while composite precipitate developed by double aging (Fig. 1) improve both strength 15 and ductility (6.1% elongation).
Example 3 A high strength Al-Li alloy was made to satisfy the requirements for high strength applications for aero-space structure. An alloy having a composition of Al-3.2Li-2Cu-2Mg-0.5Zr was solutionized at 542C for 4 hrs. As shown in Table II, conventional aging treatment (190C for i6 hrs.) showed lower strength (yield strength of 521 MPa) and ductility ~3.6%). However, double aging of the alloy (160C for 4 hrs. followed by 180C for 16 hrs.) gave significantly higher strength Iyield strength of 554 MPa) and ductility t5.5%), which meets property requirements for high strength alloys needed for aerospace structural applications.
TABLE II
0.2~ Yield Ultimate Tensile Elongation Strength (Mæa) Strength (MRa) to Failure (~) Aged at 190C
for 16 hrs. 521 595 3.6 Aged at 160C for 4 hrs. and 180C
for 16 hrs. 554 631 5.5 - ~ -. .
' .
33~
Example 4 This example illustrates the beneficial effect of the claimed process on the mechanical properties of a simple ternary alloy Al-3.~Li-0.5Zr. The alloy was 5 solutionized at 540C for 4 hrs., and subseg~ently aged as shown in Table III. The resulting tensile properties show that the claimed process results in improved ~trength and ductility compared to the conventional process.
TABLE III
Aging Treatment 0.2% Yield Ultimate Tensile ~longation to Strength (MPa) Strength (Mæa) Fract~e (~) . _ 15 140C, 16 hr. 424 442 4.2 120C, 4 hr. +
140C, 16 hr. 434 460 6.0 160C, 16 hr. 419 431 3.
20 140C, 4 hr. ~
160C, 16 hr. 425 448 4.8 140C, 16 hr. +
160C, 16 hr. 426 451 4.6 Example 5 A wide range of mechanical properties can be achieved by using multiple aging conditions. For example, a triple aging treatment (120C, 4 hrs. ~
140C, 16 hrs. ~ 160C, 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 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 pre-cipitate in low Zr containing Al-Li alloysO Fig. 5 shows the dark field electron micrograph of a typical :...
8633~2 alloy Al-3.2Li-3Cu-1.5Mg-0.2Zr which had been solutionized at 540C for 4 hrs., reheated to 170C for 4 hrs. followed by final aging at 190C for 16 hrs. The large volume fraction of composite A13 (Li, Zr) precipi-5 tate observed in such an alloy indicates that theclaimed process is also quite effective in Al-Li alloys having low ~r content of 0.2~. Having thus described the invention in rather full detail, ~t will ~e understood that such detail need not be strictly adhered 10 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.
The microstructure of these alloys is characterized by the precipitation of composite A13(Li, Zr) phase in the aluminum matrix thereof. This microstructure is developed in accordance with the process of the present 10 invention by subjecting an 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 provided wherein the 15 aluminum-based alloy produced has an improved combination of strength and ductility (at the same density).
The high strength, high ductility, low density aluminum-based alloy produced in accordance with the 20 present invention has a controlled composite A13(Li, Zr) precipitate which, advantageously, offers a wide range of strength and ductility combinations.
BRIEF DESCRIPTION OF THE DRAWINGS
The invention will be more fully understood and further advantages will become apparent when reference is made to the following detailed description and the accompanying drawings, in which:
Fig. 1 is a dark field transmission electron micro-~raph of an alloy having the composition Al-3.1Li-2Cu-lMg-0.5Zr~ the alloy having been subjected to double aging treatments (170C for 4 hrs. followe~ by 190C for 16 hrs.) to develop a composite precipitate in the aluminum matrix thereof;
Fig. 2 is a weak beam dark field micrograph of an alloy having the composition Al-3.7Li-O.SZr, illustrating the resistance of the composite precipitate to dislocation shear during deformation;
Fig. 3(a) shows the planar slip observed in an ~L28~3~2 alloy having the composition Al-3.7Li-0.5Zr, the alloy having been subjected to a conventional aging treatment (180C for 16 hours);
Fig. 3(b) shows the beneficial effect of subjecting 5 the alloy of Fig.3(a) to treatment in accordance with the claimed process tl60C for 4 hr5. followed by 180C
for 16-hrs.), thereby promoting the homogeneous deformation thereof;
Fig. 4 shows the sheared ~' precipitates observed 10 in an alloy having the composition Al-3.1Li-2Cu-lMg-0.5zr, the alloy havin~ been subjected to a conventional aging treatment (190C for 16 hours); and Fig. 5 shows the development of composite precipi-tates in an alloy having the composition Al-3.2Li-3Cu-15 1.5Mg-0.2Zr~ the alloy having been subjected to treatment in accordance with the claimed process tl70C
for 4 hrs. followed by 190C for 16 hrs).
DESCRIPTION OF THE PREFERRED EMBODIMENTS
In general, the present invention relates to the 2~ process of making high strength, high ductility, and low density Al-Li-Zr-X alloys. The process involves the use of multiple aging steps during heat treatment of the alloy. The alloy is characterized by a unique micro-structure consisting essentially of "compositen 25 A13(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~
zr-X alloys are primarily its I.i 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 extusion, forging and/or heat ~reatment parametersO Normally, an alloy in the as-processed condition (~ast, extruded or forged) has large . . .
~Z3~3~:
intermetallic particles. Further processing is required to develop certain microstructural features for certain characteristic properties.
The alloy is given an initial solutionizing treat-5 ment, that is, heating a~ a temperature (Tl) for a period of time sufficient to substantially dis501ve most of the intermetallic particles present during the forging or extrusion process, followed by cooling to ambient temperature at a sufficiently high rate to 10 retain alloying elements in said solution~ Generally, the time at temperature Tl, will be dependent on the composition of the alloy and the method of fabrication ~e.g., ingot cast, powder metallargy processed) and will typically range from about 0.1 to 10 hours. The alloy 15 is then reheated to an aging temperature, T2, for a period of time sufficient to activate the nucleation of composite A13(Li, Zr) precipitates, and cooled to ambient temperature, followed by a second aging treatment at temperature, T3, for a period of time 20 sufficient for the growth of the composite A13(Li, Zr) precipitate and a dissolution of ~' precipitate whose nucleation is not aided by Zr. The alloy at this p~int is characterized by a unique microstructure which con-sists essentially of composite A13(Li, zr) precipi-tate. This composite A13 (Li, Zr) precipitate is re-sistant to dislocation shear and quite effective in dispersing dislocation motion (Fig. 2). The result is that the alloy containing an optimum amount of composite A13(Li, Zr) precipitate deforms by a homogeneous mode of deformation resulting in improved mechanical properties. Fig. 3(b) clearly shows the h~mogeneous 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 ~' 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 exact temperature, Tl, to which the alloy is 5 heated in the solutionizing step is not critical as long as there is a dissolution of intermetallic particles at this temperature. The exact temperature, T2, in the first aging ~tep where the nucleation of composite A13 ~Li, Zr) precipitate is promoted, depends upon the 10 alloying elements present and upon the final aging step The optimum temperature range for T2, is ~rDm about 100C to 180C. The exact temperature, T3, whose range is from 120C to 200C, depends on the alloying elements present and mechanical properties desired.
15 Generally, the times at temperatures T2 and T3 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 ability of composite A13(Li, Zr) precipitates to modify the deformation behavior of 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 solutionized at 540C for 4 hrs. and subseguently aged at 160C for 4 hrs. followed by final aging at 180C for 16 hrs. Such heat treatment promotes the precipitation of composite A13(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-O.SZr) which has not been given the claimed process~
The alloy had been aged, for 16 hrs. at 180QC after solutionizing at 540 for 4 hrs. This alloy showed the pronounced planar slip which is the common deformation characteristic of brittle alloy.
34~2 In contrast, Fig. 3(b) illustrates the beneficial effect of the claimed process on the deformation benavior of an alloy having the composition Al-3.7Li-0.5Zr. After solutionizing at 540C for 4 hrs., the 5 alloy had been subjected to the double aging treatment of 160C for 4 hrs. and 180C for 16 hrs. The deform~tion mode of this alloy is quite homogeneous indicating high ductility.
Example 2 An alloy having a composition of Al-3.1Li-2Cu-lMg-0.5Zr was developed for medium strength applications as shown in Table I. The alloy was solutionized at 540C
for 2.5 hrs./ quenched into water at about 20C and given conventional single aging and the claimed double 15 aging treatments.
~28~3~2 TABLE I
0.2~ Yield Ultimate Tensile Elongation to Strength (MPa) Strength (MPa) Failure (~) Aged at 190C
for 16 hrs. 524 592 3.6 Aged at 170C --for 4 hrs. and 10190C for 16 hrs. 530 606 6.1 Conventional aginy treatment (19ODC for 16 hrs.) showed poor ductility t3.6%) due to the shearing of ~' precipitate (Fig. 4), while composite precipitate developed by double aging (Fig. 1) improve both strength 15 and ductility (6.1% elongation).
Example 3 A high strength Al-Li alloy was made to satisfy the requirements for high strength applications for aero-space structure. An alloy having a composition of Al-3.2Li-2Cu-2Mg-0.5Zr was solutionized at 542C for 4 hrs. As shown in Table II, conventional aging treatment (190C for i6 hrs.) showed lower strength (yield strength of 521 MPa) and ductility ~3.6%). However, double aging of the alloy (160C for 4 hrs. followed by 180C for 16 hrs.) gave significantly higher strength Iyield strength of 554 MPa) and ductility t5.5%), which meets property requirements for high strength alloys needed for aerospace structural applications.
TABLE II
0.2~ Yield Ultimate Tensile Elongation Strength (Mæa) Strength (MRa) to Failure (~) Aged at 190C
for 16 hrs. 521 595 3.6 Aged at 160C for 4 hrs. and 180C
for 16 hrs. 554 631 5.5 - ~ -. .
' .
33~
Example 4 This example illustrates the beneficial effect of the claimed process on the mechanical properties of a simple ternary alloy Al-3.~Li-0.5Zr. The alloy was 5 solutionized at 540C for 4 hrs., and subseg~ently aged as shown in Table III. The resulting tensile properties show that the claimed process results in improved ~trength and ductility compared to the conventional process.
TABLE III
Aging Treatment 0.2% Yield Ultimate Tensile ~longation to Strength (MPa) Strength (Mæa) Fract~e (~) . _ 15 140C, 16 hr. 424 442 4.2 120C, 4 hr. +
140C, 16 hr. 434 460 6.0 160C, 16 hr. 419 431 3.
20 140C, 4 hr. ~
160C, 16 hr. 425 448 4.8 140C, 16 hr. +
160C, 16 hr. 426 451 4.6 Example 5 A wide range of mechanical properties can be achieved by using multiple aging conditions. For example, a triple aging treatment (120C, 4 hrs. ~
140C, 16 hrs. ~ 160C, 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 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 pre-cipitate in low Zr containing Al-Li alloysO Fig. 5 shows the dark field electron micrograph of a typical :...
8633~2 alloy Al-3.2Li-3Cu-1.5Mg-0.2Zr which had been solutionized at 540C for 4 hrs., reheated to 170C for 4 hrs. followed by final aging at 190C for 16 hrs. The large volume fraction of composite A13 (Li, Zr) precipi-5 tate observed in such an alloy indicates that theclaimed process is also quite effective in Al-Li alloys having low ~r content of 0.2~. Having thus described the invention in rather full detail, ~t will ~e understood that such detail need not be strictly adhered 10 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.
Claims (8)
1. A process for increasing the strength and ductility of low density aluminum-base alloys comprising the steps of subjecting an Al-Li alloy, to multiple aging treatments to form therein a microstructure wherein a high density of shear resistant dispersoids are substantially uniformly distributed, said alloy consisting essentially of the formula AlbalZraLibXc, 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 0.15-2 wt%, "b" ranges from about 2.5-5 wt%, "c" ranges from 0 to about 5 wt% and the balance is aluminum.
2. A process according to claim 1 wherein said alloy is characterized by the precipitation of composite A13 (Li, Zr) phase in an aluminum matrix.
3. A process as recited by claim 1, wherein the number of aging treatments ranges from 2 to 10.
4. A process as recited by claim 1, wherein the number of aging treatments ranges from 2 to 5.
5. A process for making high strength, high ductility, low density aluminum-lithium alloy, comprising the steps of:
heating an aluminum alloy, consisting essentially of formula AlbalZraLibxc, wherein X is at least one element selected from the group consisting of Cu, Mg, V, Si, Sc, Ti, U, Hf, Be, Cr, Mn, Fe, Co and Ni, "a" ranges from about 0.15-2 wt%, "b" ranges from about 2.5-5wt%, "c" ranges from 0 to about 5 wt% and balance of aluminum, to a temperature, T1, for a period of time sufficient to substantially dissolve most of the intermetallic particles therein;
cooling said alloy to ambient temperature at rates sufficient to retain its elements in supersaturated solid solution;
heating said alloy to a temperature, T2, for a period of time sufficient to activate nucleation of composite A13 (Li, Zr) precipitates;
cooling said alloy to ambient temperature;
heating said alloy to a temperature, T3, for a period of time sufficient to effect additional growth of composite A13 (Li, Zr) precipitates, and dissolution of .delta.' precipitates whose nucleation is not aided by Zr;
and cooling said alloy to ambient temperature to produce therein a controlled precipitation of composite A13 (Li, Zr) phase in said aluminum matrix,
heating an aluminum alloy, consisting essentially of formula AlbalZraLibxc, wherein X is at least one element selected from the group consisting of Cu, Mg, V, Si, Sc, Ti, U, Hf, Be, Cr, Mn, Fe, Co and Ni, "a" ranges from about 0.15-2 wt%, "b" ranges from about 2.5-5wt%, "c" ranges from 0 to about 5 wt% and balance of aluminum, to a temperature, T1, for a period of time sufficient to substantially dissolve most of the intermetallic particles therein;
cooling said alloy to ambient temperature at rates sufficient to retain its elements in supersaturated solid solution;
heating said alloy to a temperature, T2, for a period of time sufficient to activate nucleation of composite A13 (Li, Zr) precipitates;
cooling said alloy to ambient temperature;
heating said alloy to a temperature, T3, for a period of time sufficient to effect additional growth of composite A13 (Li, Zr) precipitates, and dissolution of .delta.' precipitates whose nucleation is not aided by Zr;
and cooling said alloy to ambient temperature to produce therein a controlled precipitation of composite A13 (Li, Zr) phase in said aluminum matrix,
6. A process according to claim 1, further comprising the step of stretching said solutionized alloy.
7. A process according to claim 5 ! further comprising the step of stretching said alloy.
8. A process according to claim 5 wherein T1 ranges from about 500°C to 555°C, T2 ranges from about 100°C to 180°C and T3 ranges from about 120°C to 200°C.
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| US75243385A | 1985-07-08 | 1985-07-08 | |
| US752,433 | 1985-07-08 |
Publications (1)
| Publication Number | Publication Date |
|---|---|
| CA1280342C true CA1280342C (en) | 1991-02-19 |
Family
ID=25026317
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CA000513291A Expired - Lifetime CA1280342C (en) | 1985-07-08 | 1986-07-08 | High strength, ductile, low density, aluminum alloys and process for making the same |
Country Status (6)
| Country | Link |
|---|---|
| EP (1) | EP0229075B1 (en) |
| JP (1) | JPS62502295A (en) |
| AU (1) | AU578828B2 (en) |
| CA (1) | CA1280342C (en) |
| DE (1) | DE3665884D1 (en) |
| WO (1) | WO1987000206A1 (en) |
Families Citing this family (9)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US5178695A (en) * | 1990-05-02 | 1993-01-12 | Allied-Signal Inc. | Strength enhancement of rapidly solidified aluminum-lithium through double aging |
| AUPQ485399A0 (en) | 1999-12-23 | 2000-02-03 | Commonwealth Scientific And Industrial Research Organisation | Heat treatment of age-hardenable aluminium alloys |
| AUPR360801A0 (en) * | 2001-03-08 | 2001-04-05 | Commonwealth Scientific And Industrial Research Organisation | Heat treatment of age-hardenable aluminium alloys utilising secondary precipitation |
| DE102007056298A1 (en) * | 2007-11-22 | 2009-05-28 | Bayerische Motoren Werke Aktiengesellschaft | Piston for internal combustion engine, suitable for use in motor sports, is hardened by very rapid cooling of specified composition |
| RU2513492C1 (en) * | 2013-02-21 | 2014-04-20 | Открытое акционерное общество "Всероссийский институт легких сплавов" (ОАО "ВИЛС") | Aluminium-based wrought nonhardenable alloy |
| CN104694786B (en) * | 2015-01-29 | 2016-09-07 | 东莞劲胜精密组件股份有限公司 | A kind of aluminium alloy |
| CN106756272A (en) * | 2016-12-14 | 2017-05-31 | 张家港市广大机械锻造有限公司 | A kind of alloy manufacturing methods for airborne vehicle housing |
| WO2019152664A1 (en) * | 2018-01-31 | 2019-08-08 | Arconic Inc. | Corrosion resistant aluminum electrode alloy |
| KR102494830B1 (en) * | 2022-03-22 | 2023-02-06 | 국방과학연구소 | Fabrication Method of Al-Li Alloy Using Multi-Stage Aging Treatment |
Family Cites Families (5)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| DE3366165D1 (en) * | 1982-02-26 | 1986-10-23 | Secr Defence Brit | Improvements in or relating to aluminium alloys |
| CA1198656A (en) * | 1982-08-27 | 1985-12-31 | Roger Grimes | Light metal alloys |
| JPS59118848A (en) * | 1982-12-27 | 1984-07-09 | Sumitomo Light Metal Ind Ltd | Structural aluminum alloy having improved electric resistance |
| DE3411760A1 (en) * | 1983-03-31 | 1984-10-04 | Alcan International Ltd., Montreal, Quebec | METHOD FOR PRODUCING SHEET OR STRIP FROM A ROLLING BAR OF AN ALUMINUM ALLOY |
| US4661172A (en) * | 1984-02-29 | 1987-04-28 | Allied Corporation | Low density aluminum alloys and method |
-
1986
- 1986-04-11 AU AU57749/86A patent/AU578828B2/en not_active Ceased
- 1986-04-11 JP JP50235586A patent/JPS62502295A/en active Granted
- 1986-04-11 DE DE8686902711T patent/DE3665884D1/en not_active Expired
- 1986-04-11 EP EP19860902711 patent/EP0229075B1/en not_active Expired
- 1986-04-11 WO PCT/US1986/000757 patent/WO1987000206A1/en active IP Right Grant
- 1986-07-08 CA CA000513291A patent/CA1280342C/en not_active Expired - Lifetime
Also Published As
| Publication number | Publication date |
|---|---|
| DE3665884D1 (en) | 1989-11-02 |
| AU578828B2 (en) | 1988-11-03 |
| EP0229075A1 (en) | 1987-07-22 |
| JPS648066B2 (en) | 1989-02-13 |
| JPS62502295A (en) | 1987-09-03 |
| EP0229075B1 (en) | 1989-09-27 |
| WO1987000206A1 (en) | 1987-01-15 |
| AU5774986A (en) | 1987-01-30 |
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