JP2014162958A - Aluminum alloy, and method of producing the same - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide an aluminum alloy which can have sufficient heat resistance even in a case of a large-sized component without increasing hardening sensitivity, and to provide a method of producing the same.SOLUTION: An aluminum alloy includes Cu of 2.0 mass% or more and 3.7 mass% or less, Mg of 1.3 mass% or more and 2.2 mass% or less, Fe of 0.8 mass% or more and 2.0 mass% or less, Ni of 0.8 mass% or more and 2.0 mass% or less, Si of 0.31 mass% or more and 0.90 mass% or less, and Ti of 0.01 mass% or more and 0.20 mass% or less, and a content of Mn is restricted to less than 0.1 mass%, a content of Cr to less than 0.1 mass%, a content of Zr to less than 0.1 mass%, a content of Sc to less than 0.1 mass%, and a content of V to less than 0.1 mass%. The rest includes Al and inevitably included impurities.

Description

  The present invention relates to an aluminum alloy and a method for producing the same.

  In the transportation field of automobiles, railway vehicles, ships, etc., lightweight aluminum alloy materials are often used for energy saving in mechanical parts such as engine parts and compressors. In the above applications, since heat resistance of the material is mainly required, 2000 series alloys (Al-Cu alloys) that are particularly excellent in heat resistance are used among aluminum alloy types.

  As represented by the recent tightening of fuel efficiency regulations in automobiles, transportation machinery such as automobiles and ships requires further energy savings, and accordingly the temperature conditions and loads under which aluminum alloy materials are used in this application. The demand for load conditions is increasing, and it is necessary to further improve heat resistance.

  As an aluminum alloy excellent in heat resistance in this application, JIS 2618 alloy (hereinafter referred to as 2618 alloy) defined in JIS standard has been widely used. The components of the 2618 alloy are Cu: 1.9 mass% to 2.7 mass%, Mg: 1.3 mass% to 1.8 mass%, Si: 0.10 mass% to 0.25 mass%, Fe: 0.9% by mass to 1.3% by mass, Ni: 0.9% by mass to 1.2% by mass, Ti: 0.04% by mass to 0.10% by mass, Al as defined by the balance of Al, -Cu-Mg-Si-Fe-Ni alloy, especially in the temperature range of 150-200 ° C, has superior heat resistance compared to other various 2000 series alloys.

  For example, Patent Documents 1 to 3 describe alloys in which the heat resistance is further improved by improving the 2618 alloy. All the alloys described in Patent Documents 1 to 3 are transition elements such as Mn, Cr, Zr, Sc, and V in addition to Al, Cu, Mg, Si, Fe, and Ni which are basic components of the 2618 alloy. The heat resistance is improved by positively adding as an essential element. However, when the heat resistance is improved by using a transition element as described in Patent Documents 1 to 3, as an incidental effect, the quenching sensitivity is increased, and the diameter and thickness of the member are increased. When the wall thickness increases and the cooling rate inevitably decreases during quenching after the solution treatment, the heat resistance improvement effect may not be obtained. In particular, in this application, since there are many members having a relatively large diameter, a thick plate, and a large thickness, members to which these techniques can be applied are often limited to relatively small members.

  On the other hand, Patent Document 4 is an Al—Cu—Mg—Si—Fe—Ni alloy having the same structure as the 2618 alloy, but the Fe content is limited to be as low as 0.1 to 0.7%. Although a technique for increasing the solid solution amount of Cu to increase heat resistance is described, by restricting the Fe content, the distribution density of crystallized particles composed of Fe and Ni decreases, so that these particles The contribution of the dispersion strengthening to the heat resistance improvement due to the decrease in the heat resistance, it was difficult to obtain a sufficient heat resistance improvement effect that is currently required.

  On the other hand, Patent Document 5 discloses an improvement in heat resistance by limiting the solution treatment temperature condition of the Al—Cu—Mg—Si—Fe—Ni alloy in approximately the same range as the 2618 alloy to the optimum condition. The technology to plan is described. However, even with this technique, it has been difficult to obtain a sufficient heat resistance improvement effect that is currently required.

JP 2011-122180 A JP 2010-18854 A JP 2008-202121 A Japanese Patent Laid-Open No. 7-242976 JP 2008-101264 A

  As described above, against the background of recent energy savings, there is an increasing demand for heat resistance of aluminum alloys, and conventional 2618 alloys are manufactured according to ordinary methods, and solution treatment conditions and the like are optimized for 2618 alloys. When manufacturing, it was difficult to satisfy the requirements for heat resistance. Further, as in the techniques described in Patent Documents 1 to 3, even when a transition element is appropriately added to the Al—Cu—Mg—Si—Fe—Ni alloy to improve heat resistance, the quenching sensitivity is improved. In order to increase, when it became a comparatively large member, heat resistance was often insufficient.

  The present invention has been made in view of the above circumstances, and provides an aluminum alloy that can have sufficient heat resistance even in the case of a relatively large member without increasing quenching sensitivity, and a method for producing the same. For the purpose.

The aluminum alloy according to the first aspect of the present invention is:
Cu with a content of 2.0% by mass or more and 3.7% by mass or less,
Mg with a content of 1.3% by mass or more and 2.2% by mass or less;
Fe with a content of 0.8% by mass or more and 2.0% by mass or less;
Ni with a content of 0.8% by mass or more and 2.0% by mass or less;
Si with a content of 0.31 mass% or more and 0.90 mass% or less,
Ti with a content of 0.01% by mass or more and 0.20% by mass or less;
Including
The content of Mn is less than 0.1% by mass,
Cr content of less than 0.1% by mass,
The content of Zr is less than 0.1% by mass,
Sc content is less than 0.1% by mass,
The content of V is less than 0.1% by mass,
Regulated by
The balance consists of Al and inevitable impurities,
It is characterized by that.

In the matrix, it has acicular precipitates containing Al, Cu and Mg,
The average length of the acicular precipitate is 150 nm or more and 250 nm or less,
The distribution density of the acicular precipitates may be 200 / μm ² or more.

The method for producing an aluminum alloy according to the second aspect of the present invention,
A method for producing the aluminum alloy, comprising:
In the homogenization process, after the aluminum alloy cast in the casting process is held at a temperature of 470 ° C. or more and a temperature below the solidus temperature of the cast aluminum alloy for 1 hour or more,
An aluminum alloy that has been hot-worked or hot-worked and cold-worked has been hot-worked or hot-worked and cold-worked at 470 ° C. or higher in the solution treatment step. Hold for at least 1 second at a temperature below the solidus temperature of the aluminum alloy,
In the artificial aging treatment step, the aluminum alloy solution treated in the solution treatment step is held at a temperature of 170 ° C. or higher and 210 ° C. or lower for 5 hours or more.
It is characterized by that.

  ADVANTAGE OF THE INVENTION According to this invention, even if it is the case of a comparatively large member, without raising quenching sensitivity, the aluminum alloy which can have sufficient heat resistance, and its manufacturing method can be provided.

It is a figure which shows typically the distribution form of the acicular precipitate S 'which concerns on the Example of this invention. It is a figure which shows typically the shape of the rotation bending fatigue test piece which concerns on the Example of this invention.

  The present inventor has found that heat resistance can be greatly improved by optimizing the amount of Cu, Mg, Si, Fe and Ni based on an Al—Cu—Mg—Si—Fe—Ni alloy. Invented the invention. In addition, in order to further enhance the heat resistance improvement effect by optimizing the content of the above elements, the present inventor performs a homogenization process in the manufacturing process, optimizes the homogenization conditions, and after processing. The inventors have found that it is effective to optimally perform the conditions for the solution treatment and the artificial aging treatment performed, and have found the method for producing an aluminum alloy according to the present invention.

(Component elements)
First, component elements of the aluminum alloy excellent in heat resistance according to the embodiment of the present invention will be described below. The balance of the aluminum alloy according to the embodiment of the present invention is made of aluminum and inevitable impurities.

(Cu, Mg)
Cu and Mg are main elements constituting the aluminum alloy having excellent heat resistance according to the embodiment of the present invention. Cu and Mg are once dissolved in the aluminum matrix when solution treatment is performed in the temperature range of 470 ° C. or higher for the aluminum alloy according to the embodiment of the present invention, and then rapidly cooled to near room temperature. By performing an artificial aging treatment at a predetermined temperature after making a supersaturated solid solution (called quenching), fine precipitates made of Al, Cu and Mg (hereinafter referred to as fine precipitates in this specification) (Referred to as S ′ precipitates) in the matrix contributes to the improvement in room temperature strength and heat resistance of the aluminum alloy according to the embodiment of the present invention. In the aluminum alloy according to the embodiment of the present invention, the lower limit of the Cu content in the aluminum alloy is 2.0 mass%, and the lower limit of the Mg content is 1.3 mass%. If the addition amount of each element of Cu and Mg is lower than this value, the amount of solid solution of Cu and Mg is insufficient, and the precipitation density of the S ′ precipitate is lowered, so that the aluminum alloy has sufficient heat resistance. Can't get. Moreover, in the aluminum alloy which concerns on embodiment of this invention, the upper limit of the content rate of Cu in an aluminum alloy was 3.7 mass%, and the upper limit of the content rate of Mg was 2.2 mass%. If the added amount of each element of Cu and Mg is higher than this value, the solid solubility limit of the elements of Cu and Mg will be exceeded, so that these elements can be completely dissolved by solution treatment. However, it remains in the matrix as a crystallized product composed of Al, Cu and Mg. Such residual crystallized particles not only impair the high-temperature ductility of the aluminum alloy material, but also deteriorate the fatigue properties at high temperatures. For the reasons described above, in the aluminum alloy according to the embodiment of the present invention, the Cu content in the aluminum alloy is 2.0 mass% or more and 3.7 mass% or less, and the Mg content is 1.3 mass% or more. 2.2% by mass or less.

(Fe, Ni)
Fe and Ni, as well as Cu and Mg, are main elements constituting the aluminum alloy having excellent heat resistance according to the embodiment of the present invention. Fe and Ni generate fine crystallized particles composed of Al, Fe and Ni in the matrix during melting or casting. The crystallization particles are further finely divided by hot processing and cold processing after melting and casting, and dispersion is strengthened by being relatively fine (1 μm to 5 μm size) and dispersed at high density in the matrix. It is made to express and it contributes to the room temperature intensity | strength and heat resistance improvement of the aluminum alloy which concerns on embodiment of this invention, and superimposes with the strengthening by the fine precipitation by the above-mentioned Cu and Mg. In addition, these fine crystallized particles made of Al, Fe and Ni are dispersed at a high density in the matrix, so that when they are held at a high temperature, such as when an aluminum alloy solution or aluminum alloy material is used. This contributes to the stabilization of the crystal grain structure in the aluminum alloy, suppresses the growth and coarsening of the crystal grains, and contributes to the improvement of the heat resistance of the aluminum alloy according to the embodiment of the present invention. When the content ratios of Fe and Ni in the aluminum alloy are lower than 0.8% by mass, the distribution density of crystallized particles made of Al, Fe and Ni becomes low, and sufficient dispersion strengthening cannot be obtained. Further, when the content of each of Fe and Ni in the aluminum alloy is higher than 2.0% by mass, some of the crystallized particles made of Al, Fe and Ni are significantly coarsened during melting or casting. Thereafter, very coarse crystallized particles (size of 100 μm or more) remain in the matrix even in the final state after hot working or cold working. In such a case, the coarse crystallized particles are the starting point of fracture, so the high temperature ductility of the aluminum alloy is significantly reduced and the fatigue properties at high temperatures are greatly reduced. Will drop significantly. For the reasons described above, in the aluminum alloy according to the embodiment of the present invention, the Fe content in the aluminum alloy is 0.8 mass% or more and 2.0 mass% or less, and the Ni content is 0.8 mass% or more. 2.0 mass% or less.

(Si)
Si is an important essential additive element that characterizes the aluminum alloy according to the embodiment of the present invention. In the prior art, Si is often regarded as an impurity element and the upper limit is regulated, or even if it is actively added, the content is less than 0.3% by mass in most cases. It has been said. In the 2618 alloy, the Si content in the alloy is in the range of 0.10% by mass to 0.25% by mass. On the other hand, in the aluminum alloy which concerns on embodiment of this invention, let the content rate of Si in an aluminum alloy be the range of 0.31 mass% or more and 0.90 mass% or less. The reason for this will be described below. In the stage before the solution treatment of the aluminum alloy, Si is present in the matrix as Mg ₂ Si crystallized matter or precipitate together with Mg, and as a solution treatment, 470 ° C. or more and below the solidus temperature of the aluminum alloy If it is kept in the range, Mg ₂ Si dissolves and Si atoms dissolve in the matrix. After that, when quenching treatment was performed in which the aluminum alloy was quenched to near room temperature, Si atoms formed pairs with atomic vacancies in the matrix (Si-vacancy pairs), and a large number existed in equilibrium at the solution treatment temperature. Atomic vacancies are stably brought to room temperature, and as a result, a large amount of excess vacancies are formed in the matrix after quenching. Thereafter, when an artificial aging treatment is performed in a temperature range of 180 ° C. or higher and 200 ° C. or lower for 5 hours or more, the Si-hole pair is decomposed and free holes are formed in the matrix. The free vacancies promote the nucleation of fine precipitates S ′ composed of Al, Cu and Mg formed during aging precipitation, resulting in a high density of precipitation of S ′ precipitates. The room temperature strength and heat resistance of aluminum alloys are greatly improved. A preferable Si content for obtaining such an effect is in the range of 0.31% by mass or more and 0.90% by mass or less. When the Si content is less than 0.31% by mass, the amount of Si-hole pairs formed during the solution treatment is insufficient, and the precipitation density of S ′ precipitates during artificial aging after rapid cooling to near room temperature is low. Therefore, the heat resistance obtained is insufficient. If the Si content exceeds 0.90% by mass, the amount of Mg ₂ Si formed with Mg in the state before the solution treatment is too large, and cannot be completely dissolved during the solution treatment. Coarse Mg ₂ Si particles remain in the matrix. The coarse particles greatly reduce the high temperature ductility of the aluminum alloy, lower the high temperature fatigue properties, and greatly reduce the properties of the aluminum alloy.

(Ti)
In melting and casting, Ti contributes to refinement of the ingot structure by functioning as a nucleation site for solidification when the aluminum alloy solidifies. As a result, a fine grain structure is formed during the hot working following the melting and casting and the solution treatment after the cold working, thereby contributing to the improvement of the heat resistance of the aluminum alloy. In the embodiment of the present invention, Ti may be added in its entirety, or a part or all of Ti may be added in the form of a Ti-B compound. In addition, when adding with the form of the compound of Ti-B, it is more preferable to make content rate of B co-added in an aluminum alloy into the range of 0.001 mass% or more and 0.05 mass%. When the Ti content in the aluminum alloy is less than 0.01% by mass, the effect of making the ingot structure fine is insufficient. Further, if the Ti content in the aluminum alloy exceeds 0.20 mass%, a coarse compound composed of Al-Ti crystallizes during casting, and the heat resistance characteristics such as high temperature ductility and high temperature fatigue of the material are greatly reduced. . Further, when B is co-added, the effect of making the ingot finer is further obtained when the content of B to be co-added is 0.001% by mass or more. Moreover, when the content rate of B is 0.05 mass% or less, crystallization of the coarse compound which consists of Ti-B at the time of casting is suppressed, and high temperature ductility and heat resistance can be maintained higher.

(Mn, Cr, Zr, Sc, V)
In the aluminum alloy according to the embodiment of the present invention, the Mn content in the aluminum alloy is less than 0.1% by mass (including 0% by mass), and the Cr content in the aluminum alloy is 0.1% by mass. % (Including 0% by mass), the Zr content in the aluminum alloy is less than 0.1% by mass (including 0% by mass), and the Sc content in the aluminum alloy is 0.1% by mass. % (Including 0% by mass), and the V content in the aluminum alloy is preferably less than 0.1% by mass (including 0% by mass). In the present specification, the content rate of “0% by mass” includes not only 0% by mass but also a minute content that is not more than the lower limit of detection of the element in the detection device. More preferable range of the content of the transition element is Mn: 0.001% by mass or more and less than 0.1% by mass, Cr: 0.001% by mass or more and less than 0.1% by mass, Zr: 0.001% by mass or more and 0 0.1% by mass or less, Sc: 0.001% by mass or more and less than 0.1% by mass, V: 0.001% by mass or more and less than 0.1% by mass. More preferable ranges of the content are Mn: 0.001% by mass or more and less than 0.05% by mass, Cr: 0.001% by mass or more and less than 0.05% by mass, Zr: 0.001% by mass or more and 0.05% by mass or less. Less than mass%, Sc: 0.001 mass% or more and less than 0.05 mass%, V: 0.001 mass% or more and less than 0.05 mass%. Further, the more preferable range of the content is Mn: 0.001% by mass or more and less than 0.035% by mass, Cr: 0.001% by mass or more and less than 0.035% by mass, Zr: 0.001% by mass or more and 0.0. Less than 035 mass%, Sc: 0.001 mass% or more and less than 0.035 mass%, V: 0.001 mass% or more and less than 0.035 mass%.

  When transition elements Mn, Cr, Zr, Sc, and V are added to an aluminum alloy, fine dispersed particles (Mn, Cr, Zr, Sc, V) and Al that are each composed of transition elements (Mn, Cr, Zr, Sc, and V) in the matrix ( The particle size is less than 1 μm), and contributes to further stabilization of the crystal grain structure in the aluminum alloy in the same manner as the crystallized particles made of Al, Fe and Ni described above. This has the effect of further stabilizing the crystal grain structure when the aluminum alloy is held at a high temperature, and suppresses the growth and coarsening of the crystal grains, thereby contributing to further improvement in heat resistance.

Here, when fine dispersed particles composed of these transition elements (Mn, Cr, Zr, Sc, V) and Al are present in the matrix, Cu, which is a main solute element, is quenched during quenching after solution treatment. The S phase (CuMgAl ₂ ), which is a stable phase composed of Mg, may be non-uniformly nucleated on the interface between the dispersed particles and the matrix to form relatively coarsely. As a result, the solute concentrations of Cu and Mg in the matrix after quenching decrease, the precipitation density of S ′ precipitates during the subsequent artificial aging treatment decreases, and the heat resistance of the aluminum alloy may decrease. . The precipitation of the S phase during cooling during quenching becomes more pronounced as the cooling rate is lower, and the strength of the aluminum alloy after the artificial aging treatment may be greatly reduced as the cooling rate decreases. As in this case, the tendency that the strength of the aluminum alloy after the artificial aging treatment is easily affected by the cooling rate is generally expressed as “high quenching sensitivity”. As the size increases, the cooling rate during quenching may inevitably decrease, so the strength decrease after artificial aging treatment increases, the heat resistance of the aluminum alloy decreases significantly, and relatively large members It may become impossible to ensure high heat resistance. However, by setting the content of the transition elements Mn, Cr, Zr, Sc, and V in the aluminum alloy within the above-mentioned range, excessive increase in quenching sensitivity is suppressed, especially during quenching after solution treatment. Even when the member size is large, a decrease in heat resistance can be suppressed.

  On the other hand, when the effect of stabilizing the crystal grain structure by the crystallized particles composed of Al, Fe and Ni described above is utilized, the crystallized particles and the matrix interface do not serve as sites for heterogeneous nucleation, so that the solution is formed. Inhomogeneous nucleation of the S phase does not occur during the subsequent quenching. Therefore, by setting the content of the transition element in the above range, even when the size of the aluminum alloy member is increased and the cooling rate during quenching is reduced, the strength reduction of the aluminum alloy after the artificial aging treatment is small. (Low quenching sensitivity) and low heat resistance degradation of aluminum alloy.

  In the aluminum alloy which concerns on embodiment of this invention, it is more preferable to make the content rate of Mn, Cr, Zr, Sc, and V which are transition elements into the above-mentioned range for the above reason. Even if these transition elements are not actively added to the aluminum alloy, they may be mixed in a large amount when a large amount of recycled secondary alloy is used. It is more preferable that the content of the transition element is confirmed to be within the above range.

(Other elements)
In addition to the above elements, a small amount of Zn may be added for the purpose of improving the corrosion resistance and the like. In the aluminum alloy which concerns on embodiment of this invention, it is more preferable to make content rate of Zn in an aluminum alloy into 0.001 mass% or more and less than 1.0 mass%.

(Production method)
Hereinafter, the manufacturing method of the aluminum alloy extended material (aluminum alloy) excellent in heat resistance which concerns on embodiment of this invention is demonstrated. The aluminum alloy wrought material according to the embodiment of the present invention can be manufactured in accordance with a conventional method of aluminum alloy wrought material such as rolling, extrusion, and forging. By appropriately controlling, high heat resistance can be imparted to the aluminum alloy according to the embodiment of the present invention.

  Hereinafter, the process of the manufacturing method of the aluminum alloy which concerns on embodiment of this invention including the manufacturing process performed in common with aluminum alloy extended materials, such as rolling, extrusion, and forging, is described briefly. First, melting and casting are performed to obtain an aluminum alloy ingot. Then, after performing a homogenization process, hot processing is performed. Thereafter, cold working is performed as necessary. Further, intermediate annealing may be performed as necessary before or during the cold working. Then, after performing solution treatment, after quenching to a temperature near room temperature, artificial aging treatment is finally performed.

  Hereinafter, a manufacturing method is explained in full detail for every main process.

(Melting process / Casting process)
An aluminum alloy ingot is manufactured by casting an aluminum alloy melt adjusted to be within the component range of the aluminum alloy according to the embodiment of the present invention, for example, by a semi-continuous casting method (DC casting method or hot top casting method). To do. After the casting, in preparation for the subsequent hot working, the surface of the ingot surface may be cut as necessary. The chamfering may be performed after a homogenization process described later.

(Homogenization process)
Next, a homogenization process is performed on the aluminum alloy for the purpose of eliminating concentration segregation characteristic of the solidified structure formed during casting. The homogenization treatment is performed under the condition that the aluminum alloy ingot is held at a temperature of 470 ° C. or higher for 1 hour or longer. By performing the homogenization treatment under these conditions, concentration segregation is eliminated, the temperature of the solution treatment performed thereafter can be set high, and higher heat resistance can be imparted to the aluminum alloy. Become. The upper limit of the homogenization treatment temperature is not particularly defined, but it is appropriate to set it below the solidus temperature of the aluminum alloy as it is cast (after the casting process) (for example, about 530 ° C. or less). By setting the temperature of the homogenization treatment to 470 ° C. or higher, it is possible to sufficiently eliminate concentration segregation of the solidified structure, and it is possible to suppress a decrease in heat resistance. By setting the homogenization treatment time to 1 hour or longer, it is possible to sufficiently eliminate concentration segregation of the solidified structure, and it is possible to suppress a decrease in heat resistance. The upper limit of the homogenization treatment time is set as appropriate within the range where the effects of the present invention are exerted, and is not limited to the following. However, the homogenization effect is saturated and becomes almost constant even after 10 hours or more. For the reason, for example, it is more preferable that the upper limit is 10 hours.

(Hot processing process)
After the homogenization treatment, the aluminum alloy is once cooled to room temperature and then heated again to perform hot working, or after the homogenization treatment, the hot working is performed by directly adjusting the temperature to the hot working temperature. The hot working may be performed by hot rolling when manufacturing a plate, by hot extrusion when manufacturing a tube or a bar, and may be performed by hot forging when processing into other shapes. The aluminum alloy according to the embodiment of the present invention can perform any hot working, and since the working conditions do not affect the heat resistance characteristics of the final product, the hot working conditions consider the hot workability of the material. And it sets suitably in the range with the effect of this invention.

(Cold working process / Intermediate annealing process)
After hot working, if necessary, cold work is performed to finish the shape of the final product with high accuracy. Before performing cold working, intermediate annealing may be performed as necessary. Moreover, when performing cold work in 2 steps or more, intermediate annealing may be appropriately performed between the cold work and the cold work. The cold working is performed by cold rolling when manufacturing a plate, cold drawing when manufacturing a tube or a rod, and cold forging when processing into other shapes. The conditions for cold working and intermediate annealing do not affect the final heat resistance characteristics, and therefore are appropriately selected within the scope of the effects of the present invention, and are not limited to the following. More preferably, it is carried out by keeping it in a temperature range of 450 ° C. for 1 hour or longer.

(Solution treatment process)
After hot working or cold working, the solution is formed into an aluminum alloy for the purpose of dissolving Cu and Mg in the matrix and decomposing and dissolving Mg ₂ Si to dissolve Si in the matrix. Processing is performed. It is more preferable that the solution treatment is performed by holding the aluminum alloy material in a temperature range of 470 ° C. or higher and lower than the solidus temperature of the aluminum alloy for 1 second or longer. A more preferable solution treatment temperature range is 510 ° C. or more and the solidus temperature of the aluminum alloy or less, and an even more preferred solution treatment temperature range is 530 ° C. or more and the aluminum alloy solidus temperature or less. It is a range. By setting the temperature of the solution treatment to 470 ° C. or higher, it is possible to ensure the amount of Cu, Mg and Si to be dissolved, and to obtain sufficient heat resistance. The higher the solution treatment temperature is within the above temperature range, the higher the equilibrium vacancy concentration in the aluminum alloy while it is held for solution treatment. Many Si-hole pairs can be formed, and as a result, the heat resistance of the aluminum alloy is improved. By setting the holding time to 1 second or longer, it is possible to secure time for Cu and Mg to dissolve, and to secure time for Mg ₂ Si to decompose and Si to dissolve. As a result, sufficient heat resistance can be obtained. The upper limit of the solution treatment time is appropriately set within the range where the effects of the present invention are exerted, and is not limited to the following. However, even if the solution treatment is performed for more than 10 hours, the homogenization effect is saturated and almost constant. Therefore, for economic reasons, for example, it is more preferably 6 hours or less, and even more preferably 4 hours or less.

  Of the wrought material manufacturing methods, in the case of hot working by hot extrusion, the above-mentioned hot working and solution treatment can be performed together. In this case, the solution treatment is also performed. Subsequent to the hot extrusion, a step called press quenching in which rapid cooling after the solution treatment described later is continuously performed by water quenching, mist spraying, or the like may be employed.

(Rapid cooling after solution treatment (quenching))
The rapid cooling after the solution treatment is performed, for example, by immersing the aluminum alloy material in, for example, cold water or hot water, spraying mist, or spraying cold air. The cooling rate is appropriately selected within the range where the effects of the present invention are exerted. For example, as a cooling rate that can substantially impart heat resistance to the aluminum alloy, an average cooling rate from the solution treatment temperature to 100 ° C. is 1 ° C. / More preferably, it is performed in seconds or more. By setting the cooling rate to 1 ° C./second or more, it is possible to suppress a decrease in heat resistance in the aluminum alloy according to the embodiment of the present invention having low quenching sensitivity. Further, the temperature at which the cooling is terminated is appropriately selected within the range in which the effects of the present invention are exhibited, and is not limited to the following, but for example, it is more preferable to terminate the cooling in the range of 100 ° C. to room temperature. The final heat resistance of the aluminum alloy can be kept high.

(Artificial aging treatment process)
After completing the solution treatment and the subsequent rapid cooling, an artificial aging treatment is performed to increase the material strength of the aluminum alloy and improve the heat resistance. It is more preferable that the artificial aging treatment is carried out in a temperature range of 170 ° C. or higher and 210 ° C. or lower and held for 5 hours or longer. By performing the treatment under these conditions, aging precipitation proceeds moderately and high heat resistance is obtained. By setting the artificial aging treatment temperature to 170 ° C. or higher, an aging effect can be obtained and sufficient heat resistance can be obtained. By setting the artificial aging treatment temperature to 210 ° C. or lower, it is possible to suppress the precipitates from becoming excessively coarse, maintain the strength of the aluminum alloy, and sufficiently obtain the heat resistance of the aluminum alloy. Moreover, by setting the artificial aging treatment time to 5 hours or longer, precipitation can be sufficiently advanced, and sufficient heat resistance of the aluminum alloy can be obtained. The upper limit of the artificial aging treatment time is appropriately selected within the range where the effects of the present invention are exerted, and is not limited to the following. For example, it is more preferable to select 30 hours when the strength almost reaches the upper limit.

  The aluminum alloy which concerns on embodiment of this invention is manufactured according to the above process.

(Microstructure)
The aluminum alloy according to the embodiment of the present invention has the following microstructural characteristics. That is, the average length of the acicular precipitates S ′ made of Al, Cu, and Mg is in the range of 150 nm to 250 nm, and the distribution density of the acicular precipitates S ′ is 200 / μm ² or more. It has the characteristics. When the average length of the acicular precipitate S ′ is 150 nm or more, since the acicular precipitate grows sufficiently, the strength is high and the heat resistance is sufficient. Further, when the average length of the acicular precipitate S ′ is 250 nm or less, the acicular precipitate is not excessively coarsened, so that the strength is high and the heat resistance is sufficient. When the distribution density of the acicular precipitates S ′ is 200 / μm ² or more, the distribution density of the acicular precipitates is high, so that the strength is high and the heat resistance is sufficient. Further, the upper limit of the distribution density of the acicular precipitates S ′ is considered to be 500 pieces / μm ² or less from the solid solution limit of Cu and Mg.

  Hereinafter, an example of a method for measuring the average size and distribution density of the above-described acicular precipitate S ′ will be described. First, using an aluminum alloy that has been subjected to artificial aging treatment, for example, using a FIB (Focused Ion Beam) apparatus, which is one of precision processing apparatuses, a thin film specimen for TEM (Transmission Electron Microscope) observation (for example, , 0.1 μm in thickness, 5 μm in length and 10 μm in width), and TEM observation is performed. With respect to this thin film sample, an electron beam is incident on the Al lattice from the direction of [100] by TEM, and is imaged at, for example, an imaging magnification of 20,000 times to obtain, for example, TEM images for five visual fields . In the obtained TEM image, S ′ precipitates can be observed as schematically shown in FIG. Here, the needle-like precipitates S ′ are distributed in a shape extending in two directions perpendicular to the incident direction and parallel to the incident direction of the electron beam. Among these, the length of the acicular precipitates S ′ (S′1) distributed perpendicularly to each other in the direction perpendicular to the incident direction is measured for, for example, 30 arbitrary precipitates, and the average thereof is measured. By determining the value, the average length of the acicular precipitate S ′ is measured. Further, among the observed acicular precipitates S ′, those that extend in a direction parallel to the incident direction (S′2: observed in the form of particles) are distributed in a form orthogonal to each other in the plane perpendicular to the incident direction. The distribution density of the needle-like precipitates S ′ is measured by dividing the number of the needle-like precipitates S ′ (S′1), for example, by the area of the field of view obtained by dividing the total number measured for five fields of view. can do.

  As described above, the aluminum alloy (aluminum alloy stretched material) according to the embodiment of the present invention has higher heat resistance characteristics than the conventional heat-resistant aluminum alloy, and the aluminum alloy according to the embodiment of the present invention. (Aluminum alloy wrought material) is a relatively large heat-resistant member that can maintain its high heat resistance with a small decrease in heat resistance even at a relatively low cooling rate in the rapid cooling after the solution treatment. It has the feature that it can be applied to. That is, the aluminum alloy according to the embodiment of the present invention uses the component range of the Al—Cu—Mg—Si—Fe—Ni alloy optimized by various studies and the optimized manufacturing process. Compared with 2618 alloy, an aluminum alloy having higher heat resistance can be obtained. The aluminum alloy according to the embodiment of the present invention is also characterized by low quenching sensitivity, and can be used as a relatively large member among members that require a heat-resistant alloy for energy saving. Is possible.

  In addition, this invention is not limited to the said embodiment, A various deformation | transformation and application are possible.

  EXAMPLES Hereinafter, although this invention is demonstrated in detail based on an Example, this invention is not limited to this.

(Example A)
Each aluminum alloy adjusted to the composition shown in Table 1 was melted and cast using a hot top casting method to produce an ingot (billet) having a size of φ300 mm × length 800 mm. The billet was heated to a heating and holding temperature of 500 ° C. and homogenized for 6 hours. After cooling to room temperature, 5 mm of skin was faced in the circumferential direction of the billet. These billets were heated and held at 480 ° C. and then subjected to hot extrusion to obtain a round bar of φ80 mm. For the aluminum alloy material in Example A, neither cold working nor intermediate annealing was performed, and as a solution treatment, the aluminum alloy material was held at 520 ° C. for 20 minutes and then quenched into 80 ° C. hot water. The average cooling rate from 520 ° C. to 100 ° C. in Example A was 5 ° C./second. Then, the process which hold | maintains at 190 degreeC for 12 hours as an artificial aging treatment was performed, and the aluminum alloy material of the alloy numbers 1-30 was produced. The above-described aluminum alloy material was used for the heat resistance evaluation test described in detail below and the TEM observation of the needle-like precipitate S ′. In Table 1, “-” indicates that the value was below the lower limit of detection.

(Heat resistance evaluation test)
About each aluminum alloy material manufactured by the above-mentioned process (alloy numbers 1 to 30), assuming a long-term use at a high temperature, after performing a heat treatment for 100 hours at 200 ° C., the longitudinal direction of the extruded rod is A round bar tensile test piece in the shape of a JIS No. 4 test piece was collected from the center of the extruded bar so as to be in the tensile direction. This round bar tensile test piece was subjected to a high-temperature tensile test in a test atmosphere at 200 ° C. under the condition of a tensile crosshead speed of 5 mm / min, and the tensile strength (TS) at that time was measured. Table 2 shows the measurement results.

  Further, the rotating bending fatigue test piece shown in FIG. 2 is sampled from each aluminum alloy material (alloy numbers 1 to 30) so that the longitudinal direction of the extruded bar is in the tension / compression direction, and the atmospheric temperature is 200 ° C. The rotating bending fatigue test was conducted. The test frequency of the rotating bending fatigue test was 20 Hz, and the number of repetitions until breakage was measured under the conditions of maximum and minimum stress ratio R = −1 and stress amplitude of 160 MPa. The measurement results are shown in Table 2.

(Evaluation of acicular precipitate S 'distribution)
About each aluminum alloy material (alloy numbers 1-30) manufactured at the above-mentioned process, the sample for TEM observation is produced with the following method, TEM observation is performed, and average length of acicular precipitate S ', And the distribution density was measured.

  Using a FIB precision processing apparatus (manufactured by JEOL Ltd., model name: JEM-9320FIB), a thin film test piece (thickness 0.1 μm, length 5 μm, width 10 μm size) was cut out, and a transmission electron microscope (Japan) TEM observation was performed as follows using the electronic company make, model name: JEM-2100F).

  TEM observation was performed as follows. About the above-mentioned thin film test piece, an electron beam was incident on the Al lattice from the direction of [100] by TEM and imaged at a photographing magnification of 20,000 times to obtain TEM images for five fields of view. In the obtained TEM image, S ′ precipitates were observed as schematically shown in FIG. The acicular precipitate S ′ was distributed in a shape extending in two directions perpendicular to each other in a direction parallel to the incident direction of the electron beam and a plane perpendicular to the incident direction. Among these, the length of the acicular precipitates S ′ (S′1) distributed in a form orthogonal to each other in the incident direction and the vertical plane is measured for 30 arbitrary precipitates, and the average value is obtained. By obtaining, the average length of the acicular precipitate S ′ was measured. Further, among the observed acicular precipitates S ′, those that extend in a direction parallel to the incident direction (S′2: observed in the form of particles) are distributed in a form orthogonal to each other in the plane perpendicular to the incident direction. The distribution density of the needle-like precipitates S ′ was measured by dividing the number of needle-like precipitates S ′ (S′1) by the area of the field of view obtained by dividing the total number measured in five fields.

  Table 2 shows the results of the heat resistance evaluation test and the distribution state evaluation results of the acicular precipitates S ′ for each sample (alloy numbers 1 to 30).

(Evaluation result of Example A)
The evaluation results of Examples 1 to 15 (alloy numbers 1 to 15) will be described. Alloy Nos. 1 to 15 have a Cu content of 2.0% by mass or more and 3.7% by mass or less, a Mg content of 1.3% by mass or more and 2.2% by mass or less, and a content rate of 0.8%. 8 mass% or more and 2.0 mass% or less of Fe, Ni content of 0.8 mass% or more and 2.0 mass% or less, and Si content of 0.31 mass% or more and 0.90 mass% or less And Ti with a content of 0.01% by mass or more and 0.20% by mass or less, with a Mn content of less than 0.1% by mass and a Cr content of less than 0.1% by mass. The Zr content was less than 0.1% by mass, the Sc content was less than 0.1% by mass, and the V content was less than 0.1% by mass.

Moreover, as for all the aluminum alloys of alloy numbers 1-15, the average length of the acicular precipitate was 150 nm or more and 250 nm or less, and the distribution density of the acicular precipitate was 200 pieces / micrometer < ² > or more.

  As shown in Table 2, Examples 1 to 15 (Alloy Nos. 1 to 15) all have the high temperature tensile strength (251 MPa) of Comparative Example 5 (Alloy No. 20) corresponding to the components of standard 2618 alloy, Compared with the number of repetitions until break (700,000 times), the high-temperature tensile strength (in each case 270 MPa or more) and the number of repetitions until high breaks (both 900,000 or more) were shown. From the above, it was found that the aluminum alloy materials of Examples 1 to 15 were superior in high-temperature strength and high-temperature fatigue properties to standard 2618 alloy.

On the other hand, Comparative Example 1 (Alloy No. 16) had a Cu content of less than 2.0 mass%. For this reason, the Cu solid solution amount is insufficient, the precipitation density of the acicular precipitate S ′ is lowered, and sufficient heat resistance cannot be obtained, and the number of repetitions until breakage in the high temperature tensile strength and high temperature fatigue tests. Was small.
In Comparative Example 2 (Alloy No. 17), the Cu content was more than 3.7% by mass. For this reason, it is considered that the amount of Cu exceeds the solid solution limit and cannot be completely dissolved by the solution treatment, and remains in the matrix as a crystallized product composed of Al, Cu and Mg. Such residual crystallized particles markedly lowered the high temperature fatigue properties, and the number of repetitions until fracture in the high temperature fatigue test in Comparative Example 2 was greatly reduced.

In Comparative Example 3 (Alloy No. 18), the Mg content was less than 1.3% by mass. For this reason, the Mg solid solution amount is insufficient, the precipitation density of the acicular precipitate S ′ is lowered, and sufficient heat resistance cannot be obtained, and the number of repetitions until breakage in the high-temperature tensile properties and the high-temperature fatigue test. Was small.
In Comparative Example 4 (Alloy No. 19), the Mg content was more than 2.2% by mass. For this reason, it is considered that the amount of Mg exceeds the solid solution limit and cannot be completely dissolved by the solution treatment, and remains in the matrix as a crystallized product composed of Al, Cu and Mg. Such residual crystallized particles markedly lowered the high temperature fatigue properties, and the number of repetitions until fracture in the high temperature fatigue test in Comparative Example 4 was greatly reduced.

In Comparative Example 5 (Alloy No. 20), the Si content was less than 0.31% by mass. Alloy No. 20 is a standard component of the 2618 alloy, but it is considered that the amount of Si-hole pairs formed during the solution treatment was insufficient because of the low Si content. For this reason, sufficient heat resistance was not obtained, and the number of repetitions until breakage in the high temperature tensile strength and high temperature fatigue test was small.
In Comparative Example 6 (Alloy No. 21), the Si content was more than 0.90 mass%. For this reason, the amount of Mg ₂ Si formed with Mg in the state before the solution treatment is excessively large, and cannot be completely dissolved at the time of the solution treatment, and coarse Mg ₂ Si particles remain in the matrix. It is thought that it has done. The coarse particles markedly lowered the high temperature fatigue properties, and the number of repetitions until breakage in the high temperature fatigue test in Comparative Example 6 was greatly reduced.

In Comparative Example 7 (Alloy No. 22), the Fe content was less than 0.8 mass%, and the Ni content was less than 0.8 mass%. For this reason, the distribution density of the crystallized particles composed of Al, Fe and Ni is low, and sufficient dispersion strengthening cannot be obtained by these particles, and sufficient heat resistance cannot be obtained. The number of repetitions until breakage was small.
In Comparative Example 8 (Alloy No. 23), the Fe content was more than 2.0% by mass, and the Ni content was more than 2.0% by mass. For this reason, some of the crystallized particles composed of Al, Fe, and Ni become extremely coarse during melting / casting, and very coarse crystallize even in the final state after subsequent hot working or cold working. It is thought that it remained as particles (100 μm or more). Such very coarse crystallized particles markedly lowered the high temperature fatigue characteristics, and the number of repetitions until breakage in the high temperature fatigue test in Comparative Example 8 was greatly reduced.

In Comparative Example 9 (Alloy No. 24), the Ti content was less than 0.01% by mass. For this reason, it is considered that the effect of refining the ingot structure is insufficient, the ingot structure is coarse, and as a result, the crystal grain structure after rapid cooling after the solution treatment is coarse. For this reason, sufficient heat resistance was not obtained, and the high temperature tensile properties and the number of repetitions until breakage in the high temperature fatigue test were small.
In Comparative Example 10 (Alloy No. 25), the Ti content was more than 0.20% by mass. For this reason, it is thought that the coarse compound which consists of Al and Ti crystallized at the time of casting, and remained as a coarse crystallized substance also in the final state. Such a coarse crystallized product significantly deteriorated the high temperature fatigue characteristics, and the number of repetitions until breakage in the high temperature fatigue test in Comparative Example 10 was greatly reduced.

  Comparative Example 11 (Alloy No. 26) has a Mn content of 0.1% by mass or more, and Comparative Example 12 (Alloy No. 27) has a Cr content of 0.1% by mass or more. Alloy No. 28) has a Zr content of 0.1% by mass or more, Comparative Example 14 (Alloy No. 29) has an Sc content of 0.1% by mass or more, and Comparative Example 15 (Alloy No. 30). The V content was 0.1% by mass or more. Therefore, the alloys of Alloy No. 26 to Alloy No. 30 have high quenching sensitivity, and when quenched at a relatively low cooling rate (4 ° C./second), the S phase of the stable phase consisting of Mg and Cu during the quenching. However, heterogeneous nucleation is likely to occur at the interface between the dispersed particles and the matrix, resulting in a decrease in the solute concentration of Cu and Mg, resulting in a decrease in the precipitation density of the acicular precipitate S ′, and sufficient heat resistance cannot be obtained. In addition, the number of repetitions until fracture in the high temperature tensile strength and high temperature fatigue tests was small.

(Example B)
About the aluminum alloy adjusted to the composition of the alloy number 1 shown in Table 1, the aluminum alloy material used as a sample was produced with the following manufacturing processes.

  First, an aluminum alloy was melted and cast by a hot top casting method to produce an ingot (billet) having a size of φ300 mm × length 800 mm. The billet was homogenized under the conditions shown in Table 3 and then cooled to room temperature, and then the 5 mm skin was chamfered in the circumferential direction of the billet. These billets were heated and held at 440 ° C., and then subjected to hot extrusion to extrude into a round bar shape having a size of φ50 mm. Thereafter, cold working was performed by cold drawing to obtain a round bar shape of φ40 mm size. The solidus temperature of this alloy in this state was 550 ° C. Each round bar was subjected to solution treatment under the conditions shown in Table 3, and then quenched with water at 25 ° C. as a rapid cooling. The cooling rate from each solution treatment temperature to 100 ° C. was about 10 ° C./second. Thereafter, artificial aging treatment was performed under the conditions shown in Table 3, and the heat resistance evaluation test and the TEM observation of the acicular precipitate S ′ were performed by the same method as that described in Example A. The obtained evaluation The results are shown in Table 4.

(Evaluation result of Example B)
First, the evaluation results of Examples 16 to 25 (condition numbers 1 to 10) will be described. Examples 16 to 25 (Condition Nos. 1 to 10) have the components of Alloy No. 1 and are held for 1 hour or more at a temperature of 470 ° C. or higher and the solidus temperature of the aluminum alloy or lower in the homogenization step In the solution treatment step, it is held for 1 second or more at a temperature not lower than 470 ° C. and below the solidus temperature of the aluminum alloy. In the artificial aging treatment step, it is maintained at a temperature of 170 ° C. or higher and 210 ° C. or lower for 5 hours or longer. Has been manufactured.
Moreover, the average length of the acicular precipitates of the aluminum alloys of Examples 16 to 25 was 150 nm to 250 nm, and the distribution density of the acicular precipitates was 200 / μm ² or more.
The aluminum alloys of Examples 16 to 25 all exhibit high high temperature tensile strength (270 MPa or more) and high repetition rate (900,000 times or more), and are excellent in high temperature strength and fatigue properties at high temperatures. I understood it.

On the other hand, the temperature of the homogenization process was less than 470 degreeC in the comparative example 16 (condition number 11). For this reason, concentration segregation of the solidified structure could not be sufficiently eliminated, sufficient heat resistance could not be obtained, and high temperature tensile properties and the number of repetitions until fracture in a high temperature fatigue test were small.
In Comparative Example 17 (condition number 12), the time for the homogenization treatment was less than 1 hour. For this reason, concentration segregation of the solidified structure could not be sufficiently eliminated, sufficient heat resistance could not be obtained, and high temperature tensile properties and the number of repetitions until fracture in a high temperature fatigue test were small.

In Comparative Example 18 (condition number 13), the solution treatment temperature was less than 470 ° C. For this reason, since there are few amounts of Cu, Mg, and Si which form a solid solution by solution treatment, sufficient heat resistance was not obtained, and the number of repetitions until the fracture | rupture in a high temperature tensile characteristic and a high temperature fatigue test was small.
In Comparative Example 19 (condition number 14), the retention time of the solution treatment was less than 1 second. For this reason, the time for solid solution of Cu, Mg and Si is insufficient, sufficient heat resistance cannot be obtained, and the number of repetitions until breakage in the high temperature tensile property and the high temperature fatigue test is small.

In Comparative Example 20 (condition number 15), the temperature of the artificial aging treatment was less than 170 ° C. For this reason, artificial aging hardly progressed, sufficient heat resistance was not obtained, and high temperature tensile properties and the number of repetitions until breakage in a high temperature fatigue test were small.
In Comparative Example 21 (condition number 16), the temperature of the artificial aging treatment exceeded 210 ° C. For this reason, artificial aging progressed too much, sufficient heat resistance could not be obtained, and the number of repetitions until fracture in high temperature tensile properties and high temperature fatigue tests was small.
In Comparative Example 22 (condition number 17), the time for the artificial aging treatment was less than 5 hours. For this reason, artificial aging did not progress sufficiently, sufficient heat resistance could not be obtained, and the number of repetitions until fracture in high temperature tensile properties and high temperature fatigue tests was small.

Claims (3)

Cu with a content of 2.0% by mass or more and 3.7% by mass or less;
Mg with a content of 1.3% by mass or more and 2.2% by mass or less;
Fe with a content of 0.8% by mass or more and 2.0% by mass or less;
Ni with a content of 0.8% by mass or more and 2.0% by mass or less;
Si with a content of 0.31 mass% or more and 0.90 mass% or less,
Ti with a content of 0.01% by mass or more and 0.20% by mass or less;
Including
The content of Mn is less than 0.1% by mass,
Cr content of less than 0.1% by mass,
The content of Zr is less than 0.1% by mass,
Sc content is less than 0.1% by mass,
The content of V is less than 0.1% by mass,
Regulated by
The balance consists of Al and inevitable impurities,
An aluminum alloy characterized by that. In the matrix, it has acicular precipitates containing Al, Cu and Mg,
The average length of the acicular precipitate is 150 nm or more and 250 nm or less,
The acicular precipitate has a distribution density of 200 pieces / μm ² or more.
The aluminum alloy according to claim 1. A method for producing an aluminum alloy according to claim 1 or 2,
In the homogenization process, after the aluminum alloy cast in the casting process is held at a temperature of 470 ° C. or more and a temperature below the solidus temperature of the cast aluminum alloy for 1 hour or more,
An aluminum alloy that has been hot-worked or hot-worked and cold-worked has been hot-worked or hot-worked and cold-worked at 470 ° C. or higher in the solution treatment step. Hold for at least 1 second at a temperature below the solidus temperature of the aluminum alloy,
In the artificial aging treatment step, the aluminum alloy solution treated in the solution treatment step is held at a temperature of 170 ° C. or higher and 210 ° C. or lower for 5 hours or more.
The manufacturing method of the aluminum alloy characterized by the above-mentioned.