Patent Application
20150059934
The invention relates to a high-strength aluminum alloy thin extruded shape and a method for producing the same. More specifically, the invention relates to an Al—Zn—Mg—Cu-based high-strength aluminum alloy thin extruded shape that may suitably be used for transport machines (e.g., airplane) and sporting goods (e.g., bat), and a method for producing the same.
A high-strength aluminum alloy (particularly an Al—Zn—Mg—Cu-based aluminum alloy) has been widely used as a material for transport machines (e.g., airplane, helicopter, and motorcycle) and sporting goods (e.g., bat), and development of an Al—Zn—Mg—Cu-based aluminum alloy thin extruded shape having a yield strength of 700 MPa or more has been desired in order to implement a further reduction in weight.
A method that produces an aluminum alloy by consolidating a rapidly solidified powder obtained using an atomization method has been proposed in order to improve the strength of an Al—Zn—Mg—Cu-based aluminum alloy extruded shape. For example, it has been known that the tensile strength can be increased to about 900 MPa by performing a T6 treatment on a formed body produced by a powder metallurgical process using an Al alloy rapidly solidified powder that includes 5 to 11% of Zn, 2 to 4.5% of Mg, 0.5 to 2% of Cu, and 0.01 to 0.5% of Ag, with the balance substantially being Al.
However, the industrial production process becomes complex, and the production cost increases when using a rapidly solidified powder. Therefore, a wrought material produced by rolling or extrusion has been used at the expense of strength. A round bar-like extruded shape for which high strength can be easily obtained may have a yield strength of 700 MPa or more. However, it has been difficult to obtain a high-strength thin extruded shape having a yield strength of 700 MPa or more.
When producing an Al—Zn—Mg—Cu-based aluminum alloy thin extruded shape having a thickness of 5 mm or less, the brass orientation tends to be predominant in the extrusion direction, and it is difficult to obtain a high strength of 700 MPa or more. Therefore, a high-strength thin extruded shape has been produced by extruding a round bar shape or a thick shape in which the P orientation is predominant and for which high strength can be obtained, and machining the resulting extruded shape.
The invention was conceived in view of the problems relating to an Al—Zn—Mg—Cu-based high-strength aluminum alloy thin extruded shape having a thickness of 5 mm or less. An object of the invention is to provide an Al—Zn—Mg—Cu-based high-strength aluminum alloy thin extruded shape having a yield strength of 700 MPa or more, and a method for producing the same.
According to a first aspect of the invention, a high-strength aluminum alloy thin extruded shape includes 9.0 to 13.0 mass % (hereinafter may be referred to as “%”) of Zn, 2.0 to 3.0% of Mg, 1.0 to 2.0% of Cu, and 0.05 to 0.3% of Zr, with the balance being Al and unavoidable impurities, fine precipitates having a circle equivalent diameter of 5 to 20 nm being dispersed in a crystal grain of the extruded shape in a number of 4000 to 6000 per μm2.
The high-strength aluminum alloy thin extruded shape according to the first aspect of the invention may have a yield strength of 700 MPa or more and an elongation of 9% or more.
According to a second aspect of the invention, a method for producing a high-strength aluminum alloy thin extruded shape includes casting an aluminum alloy to obtain an ingot, subjecting the ingot to a homogenization treatment and hot extrusion to obtain a hot-extruded shape, and subjecting the hot-extruded shape to a solution treatment, a quenching treatment, and an aging treatment, the aluminum alloy including 9.0 to 13.0% of Zn, 2.0 to 3.0% of Mg, 1.0 to 2.0% of Cu, and 0.05 to 0.3% of Zr, with the balance being Al and unavoidable impurities, and the aging treatment including a first-stage aging treatment, a second-stage aging treatment, and a third-stage aging treatment, the first-stage aging treatment holding the hot-extruded shape at 100 to 130° C. for 6 to 48 hours, and cooling the hot-extruded shape to room temperature, the second-stage aging treatment heating the hot-extruded shape to 160 to 180° C. at a temperature increase rate of 0.5° C./sec or more, holding the hot-extruded shape at 160 to 180° C. for 10X to 10Y minutes, and cooling the hot-extruded shape to room temperature at a cooling rate of 0.02° C./sec or more, and the third-stage aging treatment holding the hot-extruded shape at 100 to 130° C. for 6 to 48 hours, and cooling the hot-extruded shape to room temperature, provided that X=−0.03×holding temperature+5.11, and Y=−0.03×holding temperature+7.07.
The aspects of the invention thus provide an Al—Zn—Mg—Cu-based high-strength aluminum alloy thin extruded shape that make it possible to obtain a structure in which fine precipitates having a circle equivalent diameter of 5 to 20 nm are dispersed in the crystal grain in a number of 4000 to 6000 per μm2, and achieve a yield strength of 700 MPa or more and an elongation of 9% or more by performing the three-step aging treatment even when producing a thin extruded shape having a thickness of 5 mm or less in which the brass orientation tends to be predominant in the extrusion direction.
The effects of each alloy component (element) of the aluminum alloy that forms the high-strength aluminum alloy thin extruded shape, and the reasons for which the content range of each alloy component is limited as described above, are described below. Zn forms an η′ phase and MgZn2 together with Mg, and improves the strength of the aluminum alloy. The Zn content is preferably 9.0 to 13.0%. If the Zn content is less than 9.0%, the aluminum alloy may exhibit insufficient strength. If the Zn content exceeds 13.0%, a decrease in ductility may occur.
Mg forms an η′ phase and MgZn2 together with Zn, and improves the strength of the aluminum alloy. The Mg content is preferably 2.0 to 3.0%. If the Mg content is less than 2.0%, the aluminum alloy may exhibit insufficient strength. If the Mg content exceeds 3.0%, a decrease in ductility may occur.
Cu improves the strength of the aluminum alloy. The Cu content is preferably 1.0 to 2.0%. If the Cu content is less than 1.0%, the aluminum alloy may exhibit insufficient strength. If the Cu content exceeds 2.0%, a decrease in ductility may occur.
Zr precipitates as Al3Zr, and suppresses recrystallization. Zr forms a fibrous structure, and improves the strength of the aluminum alloy. The Zr content is preferably 0.05 to 0.3%. If the Zr content is less than 0.05%, a decrease in strength may occur. If the Zr content exceeds 0.3%, coarse crystallized products may be produced during casting, and a decrease in ductility may occur.
Note that the aluminum alloy may include 0.30% or less of Si, 0.30% or less of Fe, and the like as unavoidable impurities. The aluminum alloy may include 0.05% or less of Ti and 0.01% or less of B so that the cast structure is refined.
The high-strength aluminum alloy thin extruded shape is produced by casting an aluminum alloy having the above composition (preferably by semi-continuous casting) to obtain an extrusion billet, subjecting the billet to a homogenization treatment (using a normal method) and hot extrusion to obtain a hot-extruded shape, and subjecting the hot-extruded shape to a solution treatment, a quenching treatment, and an aging treatment. The aging treatment includes a first-stage aging treatment, a second-stage aging treatment, and a third-stage aging treatment.
The first-stage aging treatment holds the hot-extruded shape at 100 to 130° C. for 6 to 48 hours, and cools the hot-extruded shape to room temperature. Sufficient precipitation occurs during the first-stage aging treatment. If the first-stage aging temperature is less than 100° C., sufficient precipitation may not occur. If the first-stage aging temperature exceeds 130° C., an η phase may precipitate, and a decrease in strength may occur. Note that the cooling rate when cooling the hot-extruded shape to room temperature does not affect the advantageous effects of the invention, and is not particularly limited.
The second-stage aging treatment heats the hot-extruded shape to 160 to 180° C. at a temperature increase rate of 0.5° C./sec or more, holds the hot-extruded shape at 160 to 180° C. for 10X to 10Y minutes, and cools the hot-extruded shape to room temperature at a cooling rate of 0.02° C./sec or more. The second-stage aging treatment is performed in order to redissolve the intragranular precipitates in the matrix. Note that X=−0.03×holding temperature+5.11, and Y=−0.03×holding temperature+7.07. If the holding time is less than 10X minutes, the intragranular precipitates may not be sufficiently redissolved, and a decrease in strength may occur. If the holding time exceeds 10Y minutes, a coarse η phase may precipitate, and a decrease in strength and ductility may occur.
If the second-stage aging temperature is less than 160° C., the precipitates may not be sufficiently dissolved. If the second-stage aging temperature exceeds 180° C., the heat treatment time may decrease, and industrial production may be difficult. If the temperature increase rate when heating the hot-extruded shape to 160 to 180° C. is less than 0.5° C./sec, an η phase may precipitate during heating, and a decrease in strength and ductility may occur. If the cooling rate when cooling the hot-extruded shape to room temperature is less than 0.02° C./sec, the precipitates may grow during cooling, and a decrease in strength and ductility may occur.
The third-stage aging treatment holds the hot-extruded shape at 100 to 130° C. for 6 to 48 hours, and cools the hot-extruded shape to room temperature. The η′ phase remains at the grain boundaries during the second-stage aging treatment, and the intragranular precipitates are almost completely dissolved to obtain a single matrix phase. The third-stage aging treatment is performed in order to cause reprecipitation of the η′ phase by heating the matrix phase to improve the strength of the aluminum alloy. If the third-stage aging temperature is less than 100° C., sufficient precipitation may not occur. If the third-stage aging temperature exceeds 130° C., an η phase may precipitate, and a decrease in strength may occur. Note that the cooling rate when cooling the hot-extruded shape to room temperature does not affect the advantageous effects of the invention, and is not particularly limited.
The high-strength aluminum alloy thin extruded shape has a 0.2% yield strength specified by ASTM E9 of 700 MPa or more. The high-strength aluminum alloy thin extruded shape exhibits a strength necessary for a reduction in weight even when the brass orientation is predominant in the extrusion direction.
The invention is further described below by way of examples and comparative examples to demonstrate the advantageous effects of the invention. Note that the following examples are provided for illustration purposes only, and the invention is not limited to the following examples.
An aluminum alloy having the composition shown in Table 1 was melted, and cast using a semi-continuous casting method to obtain an extrusion billet having a diameter of 90 mm. The billet was homogenized at 470° C. for 10 hours, cooled from 470° C. to 250° C. in 48 minutes (average cooling rate: 250° C./h), and cooled to room temperature. The billet was heated to 420° C. in 5 minutes in an induction furnace, held for 1 minute, and hot-extruded to produce a sheet-like extruded shape having a thickness of 4 mm and a width of 60 mm. The exit-side extrusion speed during extrusion was set to 1 m/min.
The extruded shape was heated to 470° C. at a temperature increase rate of 50° C./h, held at 470° C. for 60 minutes, quenched in water at 20 to 30° C., held at 120° C. for 24 hours, and air-cooled to room temperature (cooling rate: 25° C./sec) (first-stage aging treatment). The extruded shape was then heated to 160° C. at a temperature increase rate of 3° C./sec, held at 160° C. for 120 minutes, and air-cooled to room temperature (cooling rate: 25° C./sec) (second-stage aging treatment). The extruded shape was held at 120° C. for 24 hours, and air-cooled to room temperature (cooling rate: 25° C./sec) (third-stage aging treatment) to obtain a specimen (Specimens 1 to 16). The holding time at the holding temperature during the second-stage aging treatment falls within 10X to 10Y minutes (X=−0.03×holding temperature+5.11, and Y=−0.03×holding temperature+7.07). In Table 1, the values that fall outside the scope of the invention are underlined.
The tensile properties and the number of fine precipitates were measured as described below using Specimens 1 to 16. The measurement results are shown in Table 2. In Table 2, the number of fine precipitates that falls outside the scope of the invention is underlined. The tensile properties that fall outside the scope of the invention are also underlined.
A tensile specimen was prepared from the specimen in accordance with ASTM E9, and the tensile strength, the yield strength, and the elongation were measured. A case where the yield strength was 700 MPa or more and the elongation was 9% or more was evaluated as “Acceptable”.
The center area (2 mm in the thickness direction, and 30 mm in the widthwise direction) of the cross section (perpendicular to the extrusion direction) of the specimen was observed using a TEM (“JEM-2010” manufactured by JEOL Ltd.) (magnification: 50,000). The number (density) (per μm2) of fine precipitates having a circle equivalent diameter of 5 to 20 nm that were observed as a dark contrast in a bright-field image was calculated. The specimen was observed in three fields of view (18*104 nm2/field of view), and the average value was employed.
As shown in Table 2, Specimens 1 to 8 that fall within the scope of the invention had a structure in which fine precipitates having a circle equivalent diameter of 5 to 20 nm were dispersed in the crystal grains in a number of 4000 to 6000 per μm2. Specimens 1 to 8 had a yield strength of 700 MPa or more and an elongation of 9% or more (i.e., exhibited excellent strength and ductility).
On the other hand, Specimens 9 to 16 that fall outside the scope of the invention had a yield strength of less than 700 MPa or an elongation of less than 9%. Specimen 9 had inferior yield strength since the Zn content was too low, and a sufficient strength improvement effect could not be obtained. Specimen 10 showed insufficient elongation since the Zn content was too high, and grain boundary precipitation occurred. Specimen 11 had inferior yield strength since the Mg content was too low, and a sufficient strength improvement effect could not be obtained. Specimen 12 showed insufficient elongation since the Mg content was too high, and grain boundary precipitation occurred.
Specimen 13 had inferior yield strength since the Cu content was too low, and a sufficient strength improvement effect could not be obtained. Specimen 14 showed insufficient elongation since the Cu content was too high, and grain boundary precipitation occurred. Specimen 15 had inferior yield strength since the Zr content was too low (i.e., a recrystallized structure was formed), and a sufficient strength improvement effect could not be obtained. Specimen 16 showed insufficient elongation since the Zr content was too high, and a decrease in ductility occurred due to coarse crystallized products.
An aluminum alloy having the composition shown in Table 3 was melted, and cast using a semi-continuous casting method to obtain an extrusion billet having a diameter of 90 mm. The billet was homogenized at 470° C. for 10 hours, cooled from 470° C. to 250° C. in 48 minutes (average cooling rate: 250° C./h), and cooled to room temperature. The billet was heated to 420° C. in 5 minutes in an induction furnace, held for 1 minute, and hot-extruded to produce a sheet-like extruded shape having a thickness of 4 mm and a width of 60 mm. The exit-side extrusion speed during extrusion was set to 1 m/min.
The extruded shape was heated to 470° C. at a temperature increase rate of 50° C./h, held at 470° C. for 60 minutes, quenched in water at 20 to 30° C., and subjected to the first-stage aging treatment, the second-stage aging treatment, and the third-stage aging treatment under the conditions (a1 to a13) shown in Table 4 to obtain a specimen (Specimens 17 to 29). In the first-stage aging treatment, the extruded shape was air-cooled from the holding temperature to room temperature (cooling rate: 25° C./sec). In the third-stage aging treatment, the extruded shape was air-cooled from the holding temperature to room temperature (cooling rate: 25° C./sec). The holding time at the holding temperature during the second-stage aging treatment falls within 10X to 10Y minutes (X=−0.03×holding temperature+5.11, and Y=−0.03×holding temperature+7.07).
The tensile properties and the number of fine precipitates were measured in the same manner as in Example 1 using Specimens 17 to 29. The measurement results are shown in Table 5.
As shown in Table 5, Specimens 17 to 29 that fall within the scope of the invention had a structure in which fine precipitates having a circle equivalent diameter of 5 to 20 nm were dispersed in the crystal grains in a number of 4000 to 6000 per μm2. Specimens 17 to 29 exhibited a yield strength of 700 MPa or more and an elongation of 9% or more (i.e., exhibited excellent strength and ductility).
A sheet-like extruded shape having a thickness of 4 mm and a width of 60 mm was produced in the same manner as in Example 2 using an aluminum alloy having the composition shown in Table 3. The extruded shape was heated to 470° C. at a temperature increase rate of 50° C./h, held at 470° C. for 60 minutes, quenched in water at 20 to 30° C., and subjected to the first-stage aging treatment, the second-stage aging treatment, and the third-stage aging treatment under the conditions (b1 to b26) shown in Table 6 to obtain a specimen (Specimens 30 to 55). In the first-stage aging treatment, the extruded shape was air-cooled from the holding temperature to room temperature (cooling rate: 25° C./sec). In the third-stage aging treatment, the extruded shape was air-cooled from the holding temperature to room temperature (cooling rate: 25° C./sec). In Table 6, the values that fall outside the scope of the invention are underlined.
The tensile properties and the number of fine precipitates were measured in the same manner as in Example 1 using Specimens 30 to 55. The measurement results are shown in Table 7. In Table 7, the number of fine precipitates that falls outside the scope of the invention is underlined. The tensile properties that fall outside the scope of the invention are also underlined.
150
150
240
120
0.4
72
72
52
52
140
140
140
140
As shown in Table 7, Specimens 30 to 55 that fall outside the scope of the invention had a yield strength of less than 700 MPa and/or an elongation of less than 9%. Specimens 30 and 31 had inferior yield strength since the second-stage aging temperature was low (i.e., fine precipitates were not sufficiently redissolved), and sufficient precipitation hardening did not occur during the third-stage aging treatment.
Specimens 32, 34, and 36 had inferior yield strength since the holding time during the second-stage aging treatment was short (i.e., redissolution of the η′ phase did not proceed), and sufficient precipitation hardening did not occur during the third-stage aging treatment. Specimens 33, 35, and 37 had inferior ductility since the holding time during the second-stage aging treatment was long (i.e., precipitation of a coarse η phase occurred during heating), and had inferior yield strength since sufficient precipitation hardening did not occur during the third-stage aging treatment.
Specimen 38 had inferior ductility since the temperature increase rate during the second-stage aging treatment was low (i.e., precipitation of a coarse η phase occurred during heating), and had inferior yield strength since sufficient precipitation hardening did not occur during the third-stage aging treatment. Specimen 39 had inferior ductility since the cooling rate during the second-stage aging treatment was low (i.e., precipitation of a coarse η phase occurred during cooling), and had inferior yield strength since sufficient precipitation hardening did not occur during the third-stage aging treatment.
Specimen 40 had inferior yield strength since the holding time during the first-stage aging treatment was short, and sufficient precipitation hardening did not occur. Specimen 42 had inferior yield strength since the holding time during the first-stage aging treatment was long, and a coarse phase was formed. Specimen 42 had inferior yield strength since the holding time during the third-stage aging treatment was short, and sufficient precipitation hardening did not occur. Specimen 43 had inferior yield strength since the holding time during the third-stage aging treatment was long, and a coarse η phase was formed.
Specimen 44 had inferior yield strength since the holding time during the first-stage aging treatment was short, and sufficient precipitation hardening did not occur. Specimen 45 had inferior yield strength since the holding time during the first-stage aging treatment was long, and a coarse η phase was formed. Specimen 46 had inferior yield strength since the holding time during the third-stage aging treatment was short, and sufficient precipitation hardening did not occur. Specimen 47 had inferior yield strength since the holding time during the third-stage aging treatment was long, and a coarse η phase was formed.
Specimens 48 and 49 had inferior yield strength since the holding temperature during the first-stage aging treatment was low, and sufficient precipitation hardening did not occur. Specimens 50 and 51 had inferior yield strength since the holding temperature during the third-stage aging treatment was low, and sufficient precipitation hardening did not occur. Specimens 52 and 53 had inferior yield strength since the holding temperature during the first-stage aging treatment was high, and sufficient precipitation hardening did not occur. Specimens 54 and 55 had inferior yield strength since the holding temperature during the third-stage aging treatment was high, and sufficient precipitation hardening did not occur.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2013-178902 | Aug 2013 | JP | national |
| 2014-138398 | Jul 2014 | JP | national |
