JPH09111426A - Method for producing high strength Al alloy - Google Patents

Abstract

(57) Abstract: To obtain a high-strength Al alloy. In manufacturing an Al alloy, a material having a precipitation type Al alloy composition is subjected to a solution treatment to obtain a primary intermediate A _1.
To obtain the true strain ε ≧ 0.4 in the primary intermediate A _1.
Subjected to cold working to produce, for example a primary intermediate A ₁ performing cold working through a die hole 9 of the bending angle phi = 135 ° 1 times, secondary intermediate achieved grain refinement The step of obtaining a body and the step of subjecting the secondary intermediate to an aging treatment to expose a precipitation phase are used. The refinement of the crystal grains and the precipitation strengthening achieve high strength of the Al alloy.

Description

Detailed Description of the Invention

[0001]

TECHNICAL FIELD The present invention relates to a method for producing a high-strength Al alloy.

[0002]

2. Description of the Related Art Conventionally, in the production of Al alloys, a hot extrusion method has been applied in order to increase the strength of the Al alloys in order to make the crystal grains finer (for example, JP-A-7-706).
88 publication). In this case, in particular, the finer the crystal grains are expected as the extrusion ratio becomes larger.

[0003]

However, the extrusion ratio has an upper limit value due to equipment restrictions such as a pressurizing mechanism, so that there is a limit to the miniaturization of crystal grains. For example, the average grain size of crystal grains in a billet is 200
In the case of μm, even if the maximum extrusion ratio is set to 15, the average grain size of crystal grains in the extruded material is at most about 120 μm.

[0004]

According to the present invention, by adopting a specific means, it is possible to obtain an Al alloy in which high strength is achieved by grain refinement and precipitation strengthening or grain dispersion strengthening. It is an object of the present invention to provide the above-mentioned manufacturing method that can be performed.

According to the present invention to achieve the above object,
A step of subjecting a material having a precipitation type Al alloy composition to a solution treatment to obtain a primary intermediate body; and a cold working for producing a true strain ε ≧ 0.4 of the primary intermediate body to carry out a secondary intermediate body A step of obtaining a body, and a step of subjecting the secondary intermediate to an aging treatment,
There is provided a method for producing a high strength Al alloy using.

By the solution treatment, a primary intermediate which is a supersaturated solid solution is obtained. When the cold working specified as above is applied to the primary intermediate, the average crystal grain size of the secondary intermediate becomes half or less of that of the primary intermediate. That is, the crystal grains can be made finer. Further, the aging treatment causes the precipitation phase to appear and be dispersed, whereby the precipitation strengthening is achieved.

Al, which has been made to have a high hardness and hence a high strength, by such a refinement of crystal grains and precipitation strengthening.
An alloy can be obtained. However, the true strain ε is ε <
At 0.4, grain refinement due to cold working is not achieved.

Further, according to the present invention, a matrix made of one of Al and Al alloy, having a hardness higher than that of the matrix, and having an average particle diameter D of 0.01 μm ≦ D.
≦ 10 μm and the volume fraction Vf is 3% ≦ Vf ≦ 20
%, The true strain ε ≧ 0.
A method for producing a high-strength Al alloy that is cold worked to produce No. 4 is provided.

When the dispersed particles specified as described above are dispersed in the matrix at the volume fraction Vf, an assisting action by the dispersed particles is obtained in the cold working, so that the crystal grains of the matrix are miniaturized. Is done efficiently. In addition, particle dispersion strengthening is achieved by the dispersed particles.

By refining the crystal grains of the matrix and strengthening the dispersion of the grains, it is possible to obtain an Al alloy having a high hardness and hence a high strength.

However, if the true strain ε is ε <0.4, the grain refinement of the matrix by cold working cannot be achieved. When the average particle diameter D of the dispersed particles is D> 10 μm, the material cracks during cold working, while D <0.01
It is difficult to manufacture dispersed particles of μm. Further, when the volume fraction Vf of the dispersed particles is Vf> 20%, the toughness of the material decreases, so that the material cracks during cold working, while
When Vf <3%, the particle dispersion strengthening by the dispersed particles cannot be achieved, and the assisting action for refining the crystal grains cannot be obtained.

Further, according to the present invention, a step of subjecting a material having a precipitation type Al alloy composition to a solution treatment to obtain a primary intermediate, and a step of subjecting the primary intermediate to an aging treatment to obtain a hardness higher than that of the matrix. A step of obtaining a secondary intermediate in which the appearance and dispersion of a precipitation phase having a thickness are achieved, and a step of subjecting the secondary intermediate to a cold working for causing a true strain ε ≧ 0.4 A method of manufacturing a high strength Al alloy is provided.

By the solution treatment, a primary intermediate which is a supersaturated solid solution is obtained. In the secondary intermediate obtained by subjecting the primary intermediate to the aging treatment, a precipitation phase having high hardness appears and is dispersed, whereby precipitation strengthening is achieved. Then, in the next cold working, the assisting action of the high hardness precipitation phase is obtained, so that the crystal grains of the matrix are efficiently refined.

By making the crystal grains of the matrix finer and strengthening the precipitation, it is possible to obtain an Al alloy having a high hardness and hence a high strength. However, if the true strain ε is ε <0.4, refining of the crystal grains of the matrix by cold working cannot be achieved.

[0015]

BEST MODE FOR CARRYING OUT THE INVENTION In FIGS. 1 to 3, a die 1 used for cold working is composed of halved first and second halves 2 and 3, and these halves 2 and 3 are a plurality of through bolts. 4 and nut 5 tighten.

On the mating surface 6 of the first half body 2 with the second half body 3, a concave groove 7 which is bent and extends vertically is formed, and the concave groove 7 and the second half body 3 are formed. A die hole 9 that bends in cooperation with the mating surface 8 is formed. The die hole 9
The inlet 10 is opened on the upper surface of the die 1, and the outlet 11 is opened on the lower surface of the die 1.

Both upper and lower hole portions 13 and 14 sandwiching the bent portion 12 of the die hole 9 have a square cross section in a direction orthogonal to the axes a and b, and their cross sectional areas are substantially the same. The same is set in the embodiment.

The bending angle between the upper and lower holes 13 and 14 of the bent portion 12, that is, the upper and lower holes 13 and 14,
Portions a ₁ and b ₁ extending from the intersection c of the axes a and b of 14
The angle φ between them is set to 90 ° ≦ φ ≦ 135 °. In the illustrated example, the bending angle φ is φ = 135 °.

In the die hole 1, the dimensions of the square cross section are length d = 10.5 mm and width e = 10.5 mm, and the length f from the inlet 10 to the intersection c of both axes a and b is f = 60.
mm, the length g from the intersection point c to the exit 11 is g = 55 mm
It is.

In the first half body 2, an insertion hole 15 which opens to the lower hole portion 14 is formed, and a knockout pin 16 is inserted into the insertion hole 15.

FIG. 4 shows the first half 2 of another die 1, in which case the bending angle φ at the bending portion 12 of the die hole 9 is set to φ = 90 °. The inlet 10 of the die hole 9 is open on the upper surface of the die 1, and the outlet 11 is the die 1.
Open to the side.

In the die hole 9, the dimension of the square cross section is d = 10.5 mm in length and e = 10.5 mm in width, as described above, and the length f from the inlet 10 to the intersection c of both axes a and b. Is f = 67.5 mm, and the length g from the intersection c to the exit 11 is g = 50 mm. Other configurations are substantially the same as those of the die 1. [Example 1] [I] Precipitation type Al alloy composition in the material Table 1 shows three types of precipitation type Al alloy composition A in the material,
B and C are shown.

[0023]

[Table 1]

In Table 1, the precipitation type Al alloy composition A is A
One of the 1-Cu-based alloy composition and the Al-Cu-Mg-based alloy composition, and a suitable example is the 2000-based Al alloy composition.

In each chemical component, Cu, Mg and M
n is an essential component, and Si is a selective component.

Cu forms a GP zone (precipitation phase) during the aging treatment, whereby precipitation strengthening is achieved. However, if the Cu content is Cu <1.5 wt%, there is no effect, while if Cu> 6.8 wt%, it becomes difficult to form a solution and the θ phase (stable phase Al ₂ Cu) becomes coarse.

Mg exhibits a solid solution strengthening effect. However,
When the Mg content is Mg <0.02% by weight, the above action is diminished, while when Mg> 1.8% by weight, an intermetallic compound is produced and coarsens.

Mn exhibits an effect of preventing stress corrosion cracking.
However, when the Mn content is Mn <0.2% by weight, the above-mentioned action is diminished, while when Mn> 1.2% by weight, an intermetallic compound is produced and coarsens.

Si forms a Mg ₂ Si phase (precipitation phase) during the aging treatment, whereby the improvement of high temperature aging property and the improvement of heat resistance are achieved. However, if the Si content is Si
If it is <0.2% by weight, there is no effect, while if Si> 1.2% by weight, Si crystals are crystallized and coarsen.

As selective components, in addition to the above Si, Cr and Zr added for the purpose of increasing the recrystallization temperature, Pb and Bi added for the purpose of improving machinability, and addition for the purpose of improving heat resistance. Ni to be mentioned can be mentioned.
In this case, the content of Cr and Zr is 0.1% by weight,
The content of Pb and Bi is 0.2% by weight ≤ Pb ≤
0.6% by weight, and the Ni content is 0.4% by weight ≦
Ni is set to 1.4% by weight, respectively.

In Table 1, the precipitation type Al alloy composition B is A
1-Mg-Si alloy composition, and a suitable example is 60
A 00 series Al alloy composition can be mentioned.

Of the chemical components, Mg and Si are essential components, and Cu and Cr are optional components.

Mg produces an Mg ₂ Si phase (precipitation phase) during the aging treatment, whereby age hardening is achieved. However, when the Mg content is Mg <0.35% by weight, there is no effect, while when Mg> 1.5% by weight, the AlMg-based intermetallic compound becomes coarse.

Si exhibits the same action as Mg. However, if the Si content is Si <0.2% by weight, there is no effect,
On the other hand, when Si> 1.2% by weight, Si crystals are crystallized and coarsened.

Cu forms a GP zone during the aging treatment and strengthens the precipitation. However, if the Cu content is Cu <
0.1% by weight has no effect, while Cu> 0.4% by weight deteriorates workability and corrosion resistance.

Cr contributes to the improvement of stress corrosion cracking resistance.
However, when the Cr content is Cr <0.03% by weight, there is no effect, while when Cr> 0.35% by weight, an intermetallic compound precipitates and coarsens.

In Table 1, the precipitation type Al alloy composition C is A
1-Zn-Mg-based alloy composition and Al-Zn-Mg-C
One of the u-based alloy compositions, and a suitable example is a 7000-based Al alloy composition.

In each chemical component, Zn and Mg are essential components, and Cu and Cr, and Mn are optional components.

Zn produces a MgZn ₂ phase (precipitation phase) in the aging treatment, whereby precipitation strengthening is achieved. However, when the Zn content is Zn <4.0% by weight, there is no effect, while when Zn> 6.7% by weight, stress corrosion cracking easily occurs.

Mg exhibits the same action as Zn. However, if the Mg content is Mg <0.5% by weight, there is no effect,
On the other hand, when Mg> 2.9% by weight, the AlMg-based intermetallic compound becomes coarse.

Cu forms a GP zone (precipitation phase) during the aging treatment, whereby precipitation strengthening is achieved. However, when the Cu content is Cu <0.1% by weight, there is no effect, while when Cu> 2.6% by weight, an intermetallic compound precipitates and coarsens.

Cr makes the recrystallized grains finer and contributes to the improvement of stress corrosion cracking resistance. However, if the Cr content is Cr <0.
04% by weight has no effect, while Cr> 0.3% by weight
Then, an intermetallic compound precipitates and coarsens.

Mn exhibits the same action as Cr. However, if the Mn content is Mn <0.1% by weight, there is no effect and
On the other hand, when Mn> 0.7% by weight, intermetallic compounds are precipitated,
It becomes coarse. [II] Production of high-strength Al alloy (1) Three types of ingots having precipitation-type Al alloy compositions A to C shown in Table 2 were machined to obtain 10 mm in length, 10 mm in width,
Three types of materials with a length of 40 mm were obtained.

[0044]

[Table 2]

The precipitation-type Al alloy composition A is 2024 alloy composition, B is 6061 alloy composition, and C is 70.
75 alloy composition. (2) Each material was subjected to solution treatment shown in Table 3 to obtain primary intermediates A ₁ , B ₁ and C ₁ . Primary intermediates A ₁ , B ₁ , C
₁ corresponds to the Al alloy compositions A, B, and C, respectively.

[0046]

[Table 3]

(3) For each of the primary intermediates A ₁ , B ₁ and C ₁ ,
Cold working was performed using the die 1 with a bending angle φ = 135 ° shown in FIG. 2, the die 1 with a bending angle φ = 90 ° shown in FIG. 4 and a die with a bending angle φ = 150 ° (not shown). To obtain a secondary intermediate.

In this cold working, for example, as shown in FIG. 2, the primary intermediate A ₁ is inserted into the upper hole portion 13 from the inlet 10 of the die hole 9, and then the pressure plunger 17 is inserted into the upper hole. The primary intermediate A ₁ is inserted into the portion 13 and is extruded under pressure from the upper hole portion 13 to the lower hole portion 14 to form the primary intermediate A 1.
Shear deformation occurs in ₁ . In this case, since the lower end surface of the pressure plunger 17 moves to the vicinity of the lower end surface position of the unprocessed primary intermediate body A ₁ in FIG. 2, the primary intermediate body A ₁
Shear deformation occurs in most parts except the tip part of. The tip portion that has not generated shear deformation is cut off.

The cold working operation is performed once or more to obtain a secondary intermediate. However, in the second and subsequent cold working operations, in the work before that, the primary intermediate A
_When the h-plane of ₁ and the k-plane of the die hole 9 are opposed to each other, the primary intermediate A ₁ is rotated around its axis to rotate the primary intermediate A ₁ such as the m-plane of the primary intermediate A ₁ and the die hole 9. The k-plane is made to face each other, whereby the primary intermediate A ₁ is uniformly sheared and deformed. (4) Each secondary intermediate was subjected to an aging treatment to obtain an Al alloy.

Table 4 shows an example (1) of an Al alloy obtained by the above method using a material having a precipitation type Al alloy composition A.
~ (7) and various data on examples (8) to (11) of Al alloys obtained by subjecting the material to T6 treatment without cold working.

In the table, the true strain ε is ε = (2 / √3) ×
It was calculated from (1 / tan 0.5φ). However, φ is a bending angle. This is the same hereinafter.

[0052]

[Table 4]

As is clear from Table 4, the true strain ε in cold working is the same as in the examples (3) to (7) of Al alloys.
Is set to ε ≧ 0.4, and in the embodiment, ε ≧ 0.48,
The crystal grains can be made finer. Further, since precipitation strengthening is carried out in the aging treatment, the hardness of the Al alloy is greatly enhanced in combination with the refinement of the crystal grains, whereby the strength of the Al alloy is increased.

In this case, when the true strain ε increases, the crystal grains become finer and the hardness also increases accordingly.

In the example (4) of Al alloy, FIG. 5 is an optical micrograph showing the metal structure of the primary intermediate, and FIG. 6 is an optical micrograph showing the metal structure of the secondary intermediate. Comparing FIG. 5 with FIG. 6, it is clear that the grain refinement is achieved by cold working.

FIG. 7 is a transmission electron micrograph showing the metal structure of the secondary intermediate in the Al alloy example (4), and FIG. 8 shows the metal structure of the secondary intermediate in the Al alloy example (3). It is a transmission electron micrograph shown.

In FIGS. 7 and 8, it can be seen that the straight and curved portions existing between the two opposing arrows indicate grain boundaries, and the crystal grains are divided by this.
Further, the mottled pattern shows that a large amount of strain (dislocation) is included, and it is understood that work hardening occurs.

Al whose true strain ε is set to ε <0.4
In the alloy examples (1) and (2), grain refinement is not achieved. However, due to work hardening due to cold working,
The hardness is slightly improved as compared with the examples (8) to (11) of the Al alloy by the T6 treatment.

Table 5 shows an example (1) of an Al alloy obtained by the above method using a material having a precipitation type Al alloy composition B.
~ (7) and various data on examples (8) to (10) of Al alloys obtained by subjecting the material to T6 treatment without cold working.

[0060]

[Table 5]

As is clear from Table 5, the true strain ε in cold working is the same as in Examples (3) to (7) of Al alloys.
Is set to ε ≧ 0.4, and in the embodiment, ε ≧ 0.48,
The crystal grains can be made finer. Further, since precipitation strengthening is carried out in the aging treatment, the hardness of the Al alloy is greatly enhanced in combination with the refinement of the crystal grains, whereby the strength of the Al alloy is increased.

Table 6 shows an example (1) of an Al alloy obtained by the above method using a material having a precipitation type Al alloy composition C.
~ (6) and various data on examples (7) and (8) of Al alloys obtained by subjecting the material to T6 treatment without cold working.

[0063]

[Table 6]

As is clear from Table 6, as in Examples (3) to (6) of Al alloys, the true strain ε in cold working is ε.
Is set to ε ≧ 0.4, and in the embodiment, ε ≧ 0.48,
The crystal grains can be made finer. Further, since precipitation strengthening is carried out in the aging treatment, the hardness of the Al alloy is greatly enhanced in combination with the refinement of the crystal grains, whereby the strength of the Al alloy is increased.

From the above facts, the bending angle φ between the upper and lower holes 13 and 14 in the bending portion 12 of the die 1 is 90 °.
It can be said that it is desirable that ≦ φ ≦ 135 °. When the bending angle φ is φ> 135 °, the true strain ε ≧ 0.4 cannot be achieved. On the other hand, when the bending angle φ is φ90 °, the pressing force must be excessive, and a die 1 that can withstand the stress can be obtained. Have difficulty.

In the aging treatment, the heating temperature T is preferably 90 ° C. ≦ T ≦ 175 ° C. in order to maintain the fine state of the crystal grains. In this case, if the heating temperature T is T <90 ° C., there is no effect, while if T> 175 ° C., the crystal grains and the precipitation phase may become coarse. However, the heating temperature T changes depending on the composition of the secondary intermediate and the true strain ε, and is lower than that in the case of the usual T8 treatment.

The heating time t depends on the size of the secondary intermediate, but is preferably 1h≤t≤24h. In this case, there is no effect when the heating time t is t <1h, while t>
If it is 24 h, the crystal grains and the precipitation phase may become coarse. [Example 2] [Example-I] Using a crucible, 200 g of pure Al was melted to prepare a molten metal at 790 ° C. Then add 0.3 to the melt
84 g of pure Fe was added, and 18.64 g of K ₂ T was added.
When iF ₆ powder and 23.00 g of KBF ₄ powder are added, Ti having an average particle diameter D = 0.05 μm and including Fe is added.
B particles (dispersed particles) were generated. After the reaction was completed, the molten metal was poured into a Cu mold to obtain an Al alloy ingot.

Then, the ingot is machined,
A raw material (1) having a length of 10 mm, a width of 10 mm and a length of 40 mm was obtained.
In this case, the Vickers hardness Hv of the Al matrix is 3
0, while the Vickers hardness Hv of TiB particles is 2
800. Therefore, the hardness of TiB particles is higher than that of Al matrix.

Pure Fe, K ₂ TiF ₆ powder and KB
Various materials (2) to (10) were obtained by carrying out the same method as above under the same conditions as above except that the addition amount of the F ₄ powder was changed. In each of the materials (2) to (10), the average particle diameter D of TiB particles was D = 0.05 μm.

Table 7 shows the compounding amounts of various chemical components relating to the respective raw materials (1) to (10).

[0071]

[Table 7]

Next, Example 1 was applied to each of the materials (1) to (10).
The same cold working as in (1) to (1)
0) was obtained. Al alloys (1) to (10) are materials (1) to
Each corresponds to (10).

Table 8 shows various data regarding each Al alloy (1) to (10). The dispersion state of TiB particles is A
The l-alloy was better than the material.

[0074]

[Table 8]

As is clear from Table 8, in the materials such as Al alloys (2), (4) to (10), T
The average particle diameter D of the iB particles is 0.01 μm ≦ D ≦ 10 μm,
In this example-I, D = 0.05 μm, the volume fraction Vf of the TiB particles was set to 3% ≦ Vf ≦ 20%, and in this example-I, 3% ≦ Vf ≦ 8%, respectively, True strain ε ≧ 0.4 in processing, ε ≧ 0.
When 48 is generated, the grain size of the Al matrix is made finer and the particle dispersion is strengthened by the TiB particles, so that the hardness of the Al alloy is greatly increased, whereby the strength of the Al alloy is increased. [Example-II] Using a crucible, 200 g of Al alloy (505
2 materials) were melt | dissolved and the 800 degreeC molten metal was prepared. Next, in the molten metal, 41 g of TiC having an average particle diameter D = 5 μm and
The particles (dispersed particles) were wrapped in an aluminum foil and added, and the molten metal was sufficiently stirred with an alumina rod to uniformly disperse the TiC particles in the molten metal. Then, the crucible was cooled to obtain an Al alloy ingot.

The ingot is machined to a length of 10 m.
A raw material (1) having an m, a width of 10 mm and a length of 40 mm was obtained. In this case, the Al alloy matrix has a Vickers hardness Hv of 48.
On the other hand, the Vickers hardness Hv of TiC particles is 30.
00. Therefore, the hardness of TiC particles is higher than that of Al alloy matrix.

Further, the same method as described above was carried out under the same conditions as described above except that the average particle diameter D of TiC particles and the addition amount thereof were changed, whereby various materials (2) to
(8) was obtained.

Table 9 shows the blending amounts of various chemical components for each of the materials (1) to (8).

[0079]

[Table 9]

Next, the respective materials (1) to (7) were cold-worked in the same manner as in Example 1 to obtain various Al alloys (1) to (7). The Al alloys (1) to (7) correspond to the materials (1) to (7), respectively. Material (8) cracked at a bending angle of φ = 150 ° during cold working.

Table 10 shows various data regarding each Al alloy (1) to (7) and the material (8). The dispersion state of TiC particles was better in the Al alloy than in the raw material.

[0082]

[Table 10]

As is clear from Table 10, the average particle diameter D of the TiC particles in the material is 0.01 μm ≦ D ≦ 10 μm as in the case of Al alloys (2) to (7).
Is 5 μm ≦ D ≦ 10 μm, and the volume fraction Vf of the TiC particles is 3% ≦ Vf ≦ 20%.
% ≦ Vf ≦ 20%, and true strain ε ≧ 0.4 in cold working, and in this Example-II, ε ≧ 0.4.
When No. 8 is generated, the grain size of the Al alloy matrix is refined and the particle dispersion is strengthened by the TiC particles, so that the hardness of the Al alloy is greatly increased, and thus the strength of the Al alloy is increased.

In Example 2, as the dispersed particles, in addition to titanium boride particles such as the TiB particles and titanium carbide particles such as the TiC particles, silicon carbide particles such as SiC particles and aluminum oxide particles are used. It is possible. These are used alone or as a mixture of two or more kinds. [Example 3] [I] Precipitation type Al alloy composition in the material Table 11 shows the precipitation type Al alloy composition in the material. In the table, alloy elements AE are Fe, Si, Ti, Cu and Mn.
At least one selected from the group consisting of:

[0085]

[Table 11]

Zr produces an Al ₃ Zr phase (precipitation phase) in the aging treatment, whereby precipitation strengthening is achieved. This Al ₃ Zr phase is dispersed in the Al alloy matrix and has a hardness higher than that of the Al alloy matrix. Also, the average particle diameter D of the Al ₃ Zr phase and the volume fraction Vf
Are the same as those of the second embodiment, and for the same reason,
0.01 μm ≦ D ≦ 10 μm and 3% ≦ Vf ≦ 20%
It is.

However, when the Zr content is Zr <0.2% by weight, there is no effect because the volume fraction Vf of the Al ₃ Zr phase is Vf <3%. On the other hand, when Zr> 0.5% by weight. Al
₃ Zr phase becomes coarse.

The alloy element AE has a function of strengthening the Al alloy matrix. However, if the content of the alloy element AE is AE <0.01% by weight, there is no effect, while AE>
If it is 0.3% by weight, the amount of solid solution increases, so that the precipitation phase becomes coarse and the thermal conductivity and electrical conductivity decrease. [II] Manufacture of high-strength Al alloy (1) Machined ingots having the precipitation-type Al alloy composition shown in Table 12 to obtain a length of 10 mm, a width of 10 mm, and a length of 40 mm.
Material was obtained.

[0089]

[Table 12]

(2) The material was solution-treated at 620 ° C. for 24 hours to obtain a primary intermediate. (3) The primary intermediate was subjected to an aging treatment at 420 ° C. for 1 hour to obtain a secondary intermediate. In this case, the Vickers hardness Hv of the Al alloy matrix was 30, while the Vickers hardness Hv of the spherical Al ₃ Zr phase was 560. Therefore, the hardness of the Al ₃ Zr phase is higher than that of the Al alloy matrix. Also, the average particle size D of the Al ₃ Zr phase
Was D = 0.01 μm. (4) The secondary intermediate was subjected to the same cold working as in Example 1 to obtain an Al alloy.

Table 13 shows various data regarding Al alloys.

[0092]

[Table 13]

As is clear from Table 13, in the secondary intermediate, the average particle diameter D of the Al ₃ Zr phase was 0.01 μm ≦ D ≦
10 μm, D = 0.01 μm in this embodiment, and the volume fraction Vf of the Al ₃ Zr phase is 3% ≦ Vf ≦ 20%,
In this embodiment, Vf = 4% is set, and the true strain ε ≧ 0.4 in cold working. In this embodiment, ε is ε ≧ 0.4.
= 2.30 is obtained, grain refinement of the Al alloy matrix and precipitation strengthening by the Al ₃ Zr phase are obtained, and A
The hardness of the 1-alloy is greatly enhanced, and the high strength of the Al alloy is thereby achieved.

Since this Al alloy has a high Al concentration, it is effective as a highly heat conductive material and a highly electrically conductive material.

[0095]

According to the present invention, a high-strength Al alloy can be obtained by adopting the means specified above.

[Brief description of the drawings]

FIG. 1 is a perspective view showing an example of a die.

FIG. 2 is a sectional view taken along line 2-2 of FIG.

FIG. 3 is a view taken in the direction of arrows 3-3 in FIG. 2;

FIG. 4 is a cross-sectional view showing another example of the die, which corresponds to FIG.

FIG. 5 is an optical micrograph showing a metal structure of a primary intermediate.

FIG. 6 is an optical micrograph showing a metal structure of an example of a secondary intermediate.

FIG. 7 is a transmission electron micrograph showing a metal structure of an example of a secondary intermediate.

FIG. 8 is a transmission electron micrograph showing a metal structure of another example of the secondary intermediate.

[Explanation of symbols]

1 die 9 die hole 17 pressing plunger A ₁ 1 primary intermediates φ bending angle

Front page continuation (72) Inventor, Izuru Kanaya, 1-4-1, Chuo, Wako-shi, Saitama, Ltd., Honda R & D Co., Ltd. (72) Inventor, Masashi Shinohara, 1-1, Chuo, Wako, Saitama Stock Company Inside the Honda R & D Laboratories (72) Inventor Yoshiya Fujiwara 1-4-1 Chuo, Wako-shi, Saitama Stock Company Honda R & D Laboratories (72) Inventor Seiichi Koike 1-4-1 Chuo, Wako-shi, Saitama Incorporated Honda Technical Research Institute (72) Inventor Masao Ichikawa 1-4-1 Chuo, Wako-shi, Saitama Stock Company Honda Technical Research Institute (72) Inventor Kensuke Honma 1-14-1 Chuo, Wako-shi, Saitama Stock Company Honda Technical Research Institute

Claims (14)

[Claims] 1. A step of subjecting a material having a precipitation type Al alloy composition to a solution treatment to obtain a primary intermediate, and a cold working for causing a true strain ε ≧ 0.4 to the primary intermediate. A method for producing a high-strength Al alloy, comprising the steps of: applying a secondary intermediate to obtain a secondary intermediate; and subjecting the secondary intermediate to an aging treatment. 2. A die having a bending die hole is prepared, and the cross-sectional areas of both holes sandwiching the bending portion of the die hole in the direction orthogonal to the axial direction are set to be substantially the same, and the cold working is performed. The high strength according to claim 1, wherein the primary intermediate is extruded from one hole of the die hole to the other hole to cause shear deformation of the primary intermediate at the bent portion. Manufacturing method of Al alloy. 3. A bending angle φ between both hole portions in the bent portion.
The method for producing a high-strength Al alloy according to claim 2, wherein is 90 ° ≦ φ ≦ 135 °. 4. The precipitation-type Al alloy composition is Al-Cu.
The method for producing a high-strength Al alloy according to claim 1, which is one of a system-based alloy composition and an Al-Cu-Mg-based alloy composition. 5. The precipitation-type Al alloy composition is Al-Mg.
The method for producing a high-strength Al alloy according to claim 1, 2 or 3, which has a -Si alloy composition. 6. The precipitation-type Al alloy composition is Al-Zn.
The method for producing a high-strength Al alloy according to claim 1, which is one of a —Mg alloy composition and an Al—Zn—Mg—Cu alloy composition. 7. A matrix made of one of Al and Al alloy and having a hardness higher than that of the matrix,
In addition, a cold strain that causes true strain ε ≧ 0.4 in a material composed of dispersed particles having an average particle diameter D of 0.01 μm ≦ D ≦ 10 μm and a volume fraction Vf of 3% ≦ Vf ≦ 20%. A method for producing a high-strength Al alloy, characterized by performing hot working. 8. A die having a bending die hole is prepared, and the cross-sectional areas of both holes sandwiching the bending portion of the die hole in a direction orthogonal to the axial direction are set to be substantially the same, and the cold working is performed. The high strength A according to claim 7, wherein the material is extruded from one hole portion of the die hole to the other hole portion to cause shear deformation of the material at the bent portion.
l Alloy manufacturing method. 9. A bending angle φ between both hole portions in the bending portion.
9. The method for producing a high-strength Al alloy according to claim 8, wherein is 90 ° ≦ φ ≦ 135 °. 10. The method for producing a high strength Al alloy according to claim 7, wherein the dispersed particles are at least one selected from titanium boride particles, titanium carbide particles, silicon carbide particles and aluminum oxide particles. . 11. A step of subjecting a material having a precipitation type Al alloy composition to a solution treatment to obtain a primary intermediate, and a step of subjecting the primary intermediate to an aging treatment to obtain a precipitation phase having a hardness higher than that of a matrix. Is obtained and a step of obtaining a secondary intermediate in which dispersion is achieved, and a step of subjecting the secondary intermediate to a cold working for producing a true strain ε ≧ 0.4 are used. Method for producing high strength Al alloy. 12. A die having a bending die hole is prepared, and the cross-sectional areas of both holes sandwiching the bending portion of the die hole in the direction orthogonal to the axial direction are set to be substantially the same, and the cold working is performed. In this case, the secondary intermediate body is extruded from one hole portion of the die hole to the other hole portion to cause shear deformation at the bent portion of the secondary intermediate body.
1. The method for producing a high-strength Al alloy according to 1. 13. The method for producing a high-strength Al alloy according to claim 12, wherein a bending angle φ between both hole portions in the bending portion is 90 ° ≦ φ ≦ 135 °. 14. In the composition of the precipitation type Al alloy, Z
The content of r is 0.2% by weight ≦ Zr ≦ 0.5% by weight, and the content of at least one alloy element AE selected from Fe, Si, Ti, Cu and Mn is 0.01.
% By weight ≦ AE ≦ 0.3% by weight, and the balance is Al
The high-strength Al according to claim 11, 12, or 13.
Alloy manufacturing method.