EP2719784A1

EP2719784A1 - Aluminum alloy having excellent high-temperature characteristics

Info

Publication number: EP2719784A1
Application number: EP12796774.3A
Authority: EP
Inventors: Toshiyuki Tanaka; Yoshiki Takebayashi
Original assignee: Kobe Steel Ltd
Current assignee: Kobe Steel Ltd
Priority date: 2011-06-10
Filing date: 2012-05-14
Publication date: 2014-04-16
Also published as: CN103459630B; WO2012169317A1; JP5879181B2; CN103459630A; KR20150132604A; KR20180115350A; KR20140002063A; JP2013014835A; EP2719784A4

Abstract

An Al alloy having excellent high-temperature characteristics (heat resistance, high temperature fatigue strength, creep resistance under high temperatures, and high temperature yield strength) comprises more than 0.1 mass% and up to 1.0 mass% of Si, 3.0-7.0 mass% of Cu, 0.05-1.5 mass% of Mn, 0.01-2.0 mass% of Mg, 0.01-0.10 mass% of Ti, and 0.05-1.0 mass% of Ag, Zr being restricted to less than 0.1 mass%, and the remainder consisting of Al and unavoidable impurities.

Description

Technical Field

The present invention relates to an aluminum alloy having excellent high-temperature characteristics to be used for high-speed moving components to be rotated or slid at high speed.

Background Art

Aluminum has characteristics of low density, high strength, and easy working. As previously known, using such characteristics, aluminum is used for various applications requiring lightweight, high strength, and high working characteristics, including transport machinery such as railway vehicles, motor vehicles, and ships, various machine components, and engine components. Specifically, aluminum is used for high-speed moving components to be rotated or slid at high speed, such as a spinning rotor (small blade) and a spinning impeller (large blade) of a generator or a compressor, or a piston of an engine.
The high-speed moving components used for such applications are used in service environment at a high temperature of more than 100°C while being rotated or slid, and are therefore each required to have sufficient high-temperature characteristics (heat resistance, high-temperature fatigue strength, creep resistance under high temperatures, and high-temperature yield strength).
As described in PTL1 or PTL2, a known rotor, as one of the machine components, is improved in such high-temperature characteristics. The rotor is configured of an aluminum alloy (hereinafter, appropriately referred to as Al alloy) including aluminum as a matrix to which zirconium, manganese, iron, etc. are added as main components.
Such an Al alloy, which is an invention according to PTL1 or PTL2, contains Zr of 0.1 to 0.25 mass% as an additive alloy component, and has high static high-temperature strength, high dynamic high-temperature strength, and high creep characteristics. PTL1 or PTL2 describes that such a Zr component contributes to filling of Mn vacancies and thermal stability of the alloy.

Citation List

Patent Literature

PTL1: Japanese Unexamined Patent Application Publication (Translation of PCT Application) No. JP2009-535550 .
PTL2: Japanese Unexamined Patent Application Publication (Translation of PCT Application) No. JP2009-535551 .

Summary of the Invention

Problems that the Invention is to Solve

In such existing techniques, however, when the Zr content increases, quenching sensitivity of the Al alloy is increased. In particular, when a large material is subjected to a quenching step at a low cooling rate before artificial aging treatment, strength of the material is disadvantageously reduced. In addition, such a low cooling rate causes an increase in internal residual stress; hence, workability of the material is also problematic. In this way, particularly a large material has not been sufficiently improved in material characteristics including the high-temperature characteristics of the Al alloy.
In addition, the spinning rotor or the piston is recently used while being rotated or slid at a higher speed and/or at a higher temperature than in the past; hence, the Al alloy is required to be improved in metal fatigue strength.
An object of the invention, which has been made in light of the above-described problems, is to provide an Al alloy having excellent high-temperature characteristics. More preferably, an object of the invention is to provide an Al alloy improved in metal fatigue strength.

Means for Solving the Problems

An Al alloy having excellent high-temperature characteristics according to the present invention is characterized by containing Si of more than 0.1 to 1.0 mass%, Cu of 3.0 to 7.0 mass%, Mn of 0.05 to 1.5 mass%, Mg of 0.01 to 2.0 mass%, Ti of 0.01 to 0.10 mass%, Ag of 0.05 to 1.0 mass%, Zr in a limited content of less than 0.1 mass%, and the remainder consisting of Al and inevitable impurities.
According to the above-described composition, the Al alloy can have sufficient normal-temperature strength, sufficient high-temperature strength, sufficient creep characteristics at high temperatures, and sufficient normal-temperature yield strength, and sufficient high-temperature yield strength. This enables achievement of the properties required for the Al alloy, which is to be used particularly for high-speed moving components used at high temperatures, as the object of the invention.
Preferably, the Al alloy according to the present invention further contains V of 0.15 mass% or less.
V is further contained in the above-described composition, thereby Al-V-based dispersed particles can be precipitated in the Al alloy. Such dispersed particles have a function of hindering grain boundary migration after recrystallization, leading to prevention of coarsening of crystal grains. As a result, the normal-temperature strength, the high-temperature strength, and particularly the high-temperature metal fatigue strength of the Al alloy can be improved.
Furthermore, the Al alloy according to the present invention is preferably adjusted in largest circle-equivalent diameter of an intermetallic compound in the Al alloy to be 60 µm or less. The largest circle-equivalent diameter is described later.
The largest circle-equivalent diameter of the intermetallic compound is adjusted as described above. This allows a reduction in possibility of fracture starting from such an intermetallic compound due to metal fatigue caused by differences in material characteristics such as strength, hardness, and Young's modulus between the intermetallic compound and a matrix portion. Consequently, the metal fatigue strength is increased.

Advantageous Effects of the Invention

According to the present invention, an Al alloy having excellent high-temperature characteristics is provided. In addition, according to the present invention, an Al alloy improved in metal fatigue strength is provided.

Mode for Carrying Out the Invention

An Al alloy according to the present invention is now described in detail with an embodiment of the invention. The Al alloy has an element composition and a largest circle-equivalent diameter of an intermetallic compound as described below, and thus has sufficient high-temperature characteristics. Hence, when a high-speed moving component such as a rotor and a piston is manufactured using the Al alloy according to the invention, and when such a component is used within a product such as a generator, a compressor, and an engine, the Al alloy can withstand friction heat, which is generated by contact of the component with another component or (compressed) air while the component is rotated or slid at high speed, can withstand use under hot (compressed) air, and can withstand deforming stress exerted on the component. Consequently, the Al alloy according to the present invention can be optimally used for high-speed moving components to be rotated or slid at high speed, such as a spinning rotor (small blade) and a spinning impeller (large blade) of a generator or a compressor, or a piston of an engine.
The invention should not be limited to such an embodiment, and various modifications or alterations of the embodiment may be appropriately made within the scope without departing from the technical idea of the invention.

[Alloy Composition]

The Al alloy according to the present invention contains Si, Cu, Mn, Mg, Ti, and Ag, each being in a predetermined amount, Zr in a limited content of less than a predetermined amount, and the remainder consisting of Al and inevitable impurities.
The Al alloy further contains a predetermined amount or less of V as an optional component.
Furthermore, as an optional constitutional element, the Al alloy is limited in largest circle-equivalent diameter of an intermetallic compound to a predetermined value or less.
Description is now made on the reason for limiting the content of each component and the reason for specifying the largest circle-equivalent diameter of the intermetallic compound.

(Si: more than 0.1 to 1.0 mass%)

Si has a function of increasing strength of the Al alloy. Thus, addition of Si tends to increase precipitates effective for an increase in strength. Moreover, addition of Si is effective for suppression of formation of a dislocation loop in the alloy. As a result, addition of Si is effective for refinement of a precipitated phase and uniform precipitation of the phase.
If the content of Si is 0.1 mass% or less, such effects are small. If the content of Si exceeds 1.0 mass%, a coarse intermetallic compound is formed, causing bad forming, a reduction in metal fatigue strength, and failure of a high-speed moving component such as a spinning rotor, a spinning impeller, and a piston.
Consequently, the Si content is specified to be more than 0.1 to 1.0 mass%. Preferably, the Si content is more than 0.1 to 0.7 mass%. More preferably, the Si content is more than 0.1 to 0.6 mass%.

(Cu: 3.0 to 7.0 mass%)

Cu is an essential component of the Al alloy according to the present invention. The Al alloy according to the invention is used for machine components such as a spinning rotor, a spinning impeller, and a piston. Cu is an indispensable component for securing sufficient creep characteristics under high temperatures, sufficient normal-temperature yield strength, and sufficient high-temperature yield strength, which are mainly required for the Al alloy in applications of the invention, by both effects of solution strengthening and precipitation strengthening. In this way, Cu increases strength of the Al alloy by both effects of solution strengthening and precipitation strengthening. Specifically, Cu induces fine and dense precipitation of a θ' phase and a Ω phase on a (100) plane and a (111) plane of the Al alloy during the high-temperature artificial aging treatment, allowing strength of the Al alloy to be increased after the artificial aging treatment.
Such an effect is exhibited in the Cu content of 3.0 mass% or more. If the content of Cu is less than 3.0 mass%, the effect is small, so that the Al alloy does not have sufficient creep characteristics under high temperatures and sufficient high-temperature yield strength. If the content of Cu exceeds 7.0 mass%, strength excessively increases, leading to degradation in forgeability of the Al alloy. Moreover, intermetallic compounds in the Al alloy structure are readily coarsened, leading to a reduction in metal fatigue strength of the Al alloy.
Consequently, the Cu content is specified to be 3.0 to 7.0 mass%. Preferably, the Cu content is 4.0 to 7.0 mass%. More preferably, the Cu content is more than 4.5 to 7.0 mass%.

(Mn: 0.05 to 1.5 mass%)

Mn allows the microstructure of the Al alloy to be formed into a fibrous structure to increase the normal-temperature strength and the high-temperature strength of the Al alloy. In addition, Mn induces precipitation of Al-Mn-based dispersed particles, which are composed of a thermally stable compound, in the Al alloy matrix during the homogenization heating treatment. Such dispersed particles include Al₂₀Cu₂Mn₃, for example. Such dispersed particles have a function of hindering grain boundary migration after recrystallization, and therefore effectively prevent coarsening of crystal grains.
If the content of Mn is less than 0.05 mass%, such effects are small. If the content of Mn exceeds 1.5 mass%, a coarse insoluble intermetallic compound is readily formed during the melting and casting, causing bad forming and fracture of the Al alloy material.
Consequently, the Mn content is specified to be 0.05 to 1.5 mass%. Preferably, the Mn content is 0.05 to 1.0 mass%. More preferably, the Mn content is 0.05 to 0.8 mass%.

(Mg: 0.01 to 2.0 mass%)

As with Cu, Mg is an indispensable component for securing sufficient certain creep characteristics under high temperatures, sufficient normal-temperature yield strength, and sufficient high-temperature yield strength, which are mainly required for the Al alloy in applications of the invention, by both effects of solution strengthening and precipitation strengthening. Specifically, as with Cu, Mg also induces fine and dense precipitation of the θ' phase and the Ω phase on a (100) plane and a (111) plane of the Al alloy during the high-temperature artificial aging treatment, allowing strength of the Al alloy to be increased after the artificial aging treatment.
The effect is exhibited in the Mg content of 0.01 mass% or more. If the content of Mg is less than 0.01 mass%, the effect is small, so that the Al alloy does not have sufficient characteristics under high temperatures, sufficient normal-temperature yield strength, and sufficient high-temperature yield strength. If the content of Mg exceeds 2.0 mass%, strength excessively increases, leading to degradation in workability such as forgeability of the Al alloy.
Consequently, the Mg content is specified to be 0.01 to 2.0 mass%. Preferably, the Mg content is 0.01 to 1.5 mass%. More preferably, the Mg content is 0.01 to 1.0 mass%.

(Ti: 0.01 to 0.10 mass%)

Ti shows an effect of refining crystal grains during the casting.
If the content of Ti is less than 0.01 mass%, the effect is small. If the content of Ti exceeds 0.10 mass%, a coarse intermetallic compound is formed. Such a coarse intermetallic compound becomes a start point of fracture of the Al alloy material during the forming. Hence, addition of Ti of more than 0.10 mass% leads to degradation in formability of the Al alloy.
Consequently, the Ti content is specified to be 0.01 to 0.10 mass%.

(Ag: 0.05 to 1.0 mass%)

Ag forms a fine and uniform Ω phase in the Al alloy, and extremely reduces width of a solute-depleted precipitate free zone (PFZ), and thus improves the normal-temperature strength, the high-temperature strength, and the high-temperature creep characteristics of the Al alloy.
If the content of Ag is less than 0.05 mass%, the effect is small. If the content of Ag exceeds 1.0 mass%, the effect is saturated.
Consequently, the Ag content is specified to be 0.05 to 1.0 mass%. Preferably, the Ag content is 0.05 to 0.7 mass%.

(Zr: less than 0.1 mass% (including 0 mass%)

Zr induces precipitation of Al-Zr-based dispersed particles, which are composed of a thermally stable compound, in the Al alloy structure during the homogenization heat treatment. Such dispersed particles allow the microstructure of the Al alloy to be formed into a fibrous structure to increase the normal-temperature strength and the high-temperature strength of the Al alloy.
In a quenching step after a solution treatment step, however, in the case where average cooling rate between 400°C and 290°C is lowered to 500 °C/sec or less, and if Zr is contained in the amount of 0.1 mass% or more, a stable phase including AlCu₂ is coarsely precipitated in the periphery of each of the Al-Zr-based dispersed particles during the quenching treatment after the solution treatment. As a result, even if the artificial aging treatment is subsequently performed at high temperature, the high-temperature yield strength may be disadvantageously reduced.
Consequently, the Zr content is specified to be less than 0.1 mass% in order to reduce quenching sensitivity of the Al alloy.

(V: less than 0.15 mass%)

V as an optional component is precipitated in a form of an Al-V-based compound in the Al alloy structure and thus allows the high-temperature metal fatigue strength of the Al alloy to be increased. Moreover, V induces precipitation of Al-V-based dispersed particles, which are composed of a thermally stable compound, in the Al alloy structure during the homogenization heat treatment. Such dispersed particles have a function of hindering grain boundary migration after recrystallization, and therefore show an effect of preventing coarsening of crystal grains.
Through such an effect, V allows the microstructure of the Al alloy to be formed into a fibrous structure so that the Al alloy is increased in normal-temperature strength, high-temperature strength, and particularly high-temperature metal fatigue strength. In addition, V has a small function of allowing a stable phase to be coarsely precipitated compared with Zr, Cr, and Mn, and is therefore more preferred for improving the normal-temperature strength, the high-temperature strength, and the high-temperature metal fatigue strength of the Al alloy.
Consequently, the crystal grain size should be reduced to 500 µm or less to further secure the high-temperature characteristics of the Al alloy. To achieve this, V is preferably selectively contained in the amount of 0.15 mass% or less. If the content of V is less than 0.05 mass%, the effect is small. If the content of V exceeds 0.15 mass%, a coarse insoluble intermetallic compound is readily formed during the melting and casting, causing bad forming and fracture of the Al alloy.
Consequently, the V content is preferably 0.15 mass% or less. Although the V content may be 0 mass%, the lower limit value of the V content is preferably 0.05 mass%.

(Remainder: Al and Inevitable Impurities)

In addition, the composition of the Al alloy includes the remainder consisting of Al and inevitable impurities. Examples of the inevitable impurities include Ni, Zn, and B that are contained in ingot metal or intermediate alloy in the amount within a generally known range. Each of the impurities is allowed to be contained within the range without degrading the characteristics including the high-temperature characteristics according to the invention.

(Other Elements)

Other elements are each also allowed to be contained within the range without degrading the characteristics including the high-temperature characteristics according to the invention.
Fe has an effect of improving the high-temperature characteristics of the Al alloy, and may be mixed in from scraps etc. Thus, the content of Fe is specified to be 0.15 mass% or less.

[Intermetallic Compound: Largest Circle-Equivalent Diameter of 60 µm or less]

The largest circle-equivalent diameter of the intermetallic compound is preferably 60 µm or less in light of improvement in metal fatigue strength. More preferably, the largest circle-equivalent diameter is 50 µm or less. Most preferably, the largest circle-equivalent diameter is 40 µm or less. The largest circle-equivalent diameter refers to a diameter of a circle having area corresponding to area of a largest intermetallic compound. In a specific calculation method, the largest circle-equivalent diameter is estimated using extremal statistics described later.
If a large intermetallic compound exists in the Al alloy, fracture of the Al alloy material occurs starting from such an intermetallic compound due to differences in material characteristics such as strength, hardness, and Young's modulus between the intermetallic compound and the matrix, possibly resulting in a reduction in metal fatigue strength of the Al alloy material. Specifically, if an intermetallic compound has a large size, such an intermetallic compound highly possibly becomes a start point of fracture of the Al alloy material. Hence, it is desirable to control the largest circle-equivalent diameter (size) of the intermetallic compound contained in the Al alloy.
The largest circle-equivalent diameter (size) of the intermetallic compound is controlled to be smaller, thereby the metal fatigue strength of the Al alloy is improved.
The size of the intermetallic compound can be controlled through an appropriate combination of casting conditions (cooling rate, adjustment of casting size, etc.), homogenization heat treatment conditions (temperature, time, multistep temperature adjustment, etc.), forging conditions (a forging ratio, forging temperature, etc.), and solution treatment conditions (temperature, time, etc.).
In this embodiment, for example, the casting is performed at a cooling rate of 0.05 °C/sec or more, the homogenization heat treatment is performed within a temperature range of 500 to 545°C, and the hot forging is performed within a temperature range of 280 to 430°C, thereby the largest circle-equivalent diameter of the intermetallic compound is controlled to be 60 µm or less.

(Homogenization Heat Treatment)

Furthermore, the homogenization heat treatment is preferably performed at a temperature as high as possible within the above-described temperature range (500 to 545°C) of the homogenization heat treatment but within a temperature range in which eutectic melting does not occur. Through such a temperature condition, the intermetallic compound is effectively melted and diffused into a base metal. As a result, size of the intermetallic compound can be reduced.
Multistep homogenization heat treatment is advantageously performed in at least two steps depending on a type of the intermetallic compound, as a more effective process for reducing size of the intermetallic compound without eutectic melting of the intermetallic compound.
Such multistep homogenization heat treatment process is performed after setting of appropriate conditions (heating rate, homogenization temperature, and treatment time) corresponding to a type of the intermetallic compound.
For example, for any intermetallic compound, heat treatment is appropriately performed at a relatively low temperature within the above-described temperature range (500 to 545°C) of the homogenization heat treatment so that the intermetallic compounds are sufficiently melted and diffused. Subsequently, heat treatment is performed at a relatively high temperature within the temperature range of the homogenization heat treatment to reduce the size of each intermetallic compound. Such homogenization heat treatment, in which temperature is adjusted in multiple steps, is advantageously performed.
In addition, effects similar to those of the multistep homogenization heat treatment process are also given by another process. In such a process, the Al alloy is heated within a temperature range, in which eutectic melting of the intermetallic compound does not occur, at a relatively low heating rate up to the homogenization heat treatment temperature. This process may be performed in combination with the above-described multistep homogenization heat treatment. Such heating rate must be appropriately set depending on a type, size, amount, etc. of the intermetallic compound.
Such homogenization heat treatment processes each enable a reduction in size of the intermetallic compound while preventing eutectic melting of the intermetallic compound. Such a reduction in size of the intermetallic compound suppresses fatigue fracture starting from the intermetallic compound, leading to an increase in fatigue strength. Moreover, individual elements contained in the intermetallic compound are uniformly diffused into a base metal. This allows strength of the base metal to be increased by the effects of solution strengthening and precipitation strengthening. In addition, this also allows an increase in each of elongation, impact value, and metal fatigue strength of the Al alloy.
The largest circle-equivalent diameter (size) of the intermetallic compound is controlled to be smaller, thereby the metal fatigue strength is increased as above. Such an effect is shown not only in the range of the element composition according to the invention but also in a range of another element composition.
The value of the metal fatigue strength may be related not only with the size of the intermetallic compound but also with each of an aspect ratio, a shape, hardness, Young's modulus, and an amount of the intermetallic compound, and with an aerial proportion of the intermetallic compound in a metal structure.

(Measurement of Intermetallic Compound)

The size (largest circle-equivalent diameter) of the intermetallic compound can be calculated through estimation with extremal statistics as described in Y. Murakami "Metal Fatigue Influence of Micro-Defects and Inclusions" issued by Yokendo co. Ltd. (first edition of the OD version, pp. 233 to 250). The estimation with extremal statistics refers to a process of estimating an extremal value based on an extremal statistics graph being beforehand created.
In summary, a surface of a sample is polished, and then a region having predetermined inspecting standard area is photographed at a statistically sufficient number of points as by a microscope in such a manner that inspected portions are not overlapped with each other. Subsequently, an intermetallic compound having the largest area is selected in the individual inspecting standard area region, and the square root of the area of each largest intermetallic compound is calculated. Subsequently, a cumulative distribution function or cumulative frequency distribution and a scaling parameter are calculated. Subsequently, the root square value of the area of each largest intermetallic compound is plotted in a horizontal axis and the cumulative distribution function or the scaling parameter in a vertical axis. The root square value of the area of the largest intermetallic compound in the horizontal axis is substituted for the largest circle-equivalent diameter of the largest intermetallic compound, and a largest intermetallic compound distribution line is calculated.
Finally, such a largest intermetallic compound distribution line is used to estimate the largest circle-equivalent diameter of the largest intermetallic compound in predictive area. In this measurement; the largest circle-equivalent diameter of the largest intermetallic compound is estimated with the inspecting standard area of 0.37 mm², the number of inspecting portions of 40, and the predictive area of 100 mm².

[Method of Manufacturing Al Alloy]

A method of manufacturing the Al alloy according to the present invention is now described.
A manufacturing process itself of the Al alloy according to the invention is essentially the same as a manufacturing process of the existing Al alloy. Specifically, the manufacturing process of the Al alloy according to the invention includes a casting step, a homogenization heat treatment step, a hot forging step, a solution treatment step, a quenching treatment step, and an artificial aging treatment step. A cold compression (working) step may be included as necessary.
T61 temper and T6 temper described later are each performed through the solution treatment step, the quenching treatment step, and the artificial aging treatment step. Furthermore, T652 temper is performed through the solution treatment step, the quenching treatment step, the cold compression (working) step, and the artificial aging treatment step. Such types of temper are performed in an appropriately selected manner depending on size or an application of a member to be manufactured.

(Casting Step)

In the casting step, the Al alloy having the above-described composition is melted and casted to fabricate an ingot. Any of previously known casting processes may be used as the casting process without limitation. For example, the ingot can be fabricated by casting as follow. That is, a casting process is appropriately selected from typical melting and casting processes such as a continuous casting-rolling process and a semi-continuous casting process (DC casting process). Subsequently, the ingot is fabricated by the selected casting process using an Al alloy molten metal prepared into the composition range of the invention.

(Homogenization Heat Treatment Step)

Through the homogenization heat treatment, micro segregation caused by solidification is eliminated, a supersaturated solid solution element is precipitated, and a metastable phase is changed into an equilibrium phase. If the temperature of the homogenization heat treatment is less than 500°C, an intermetallic compound such as a crystallized grain in the ingot is not dissolved, leading to insufficient homogenization. If the temperature of the homogenization heat treatment exceeds 545°C, burning may occur at a high possibility. Consequently, the temperature of the homogenization heat treatment is specified to be within a range of 500 to 545°C.
In the case where the multistep homogenization heat treatment is performed, the heat treatment condition must be set in correspondence to a type of the intermetallic compound as described above. Even if the homogenization heat treatment is performed at a relatively low heating rate, the heat treatment condition must also be set in correspondence to a type of the intermetallic compound as described above.

(Hot Forging Step)

The temperature condition of the hot forging step, together with the forging ratio described later, is important to reproducibly manufacture the Al alloy with certain characteristics. Specifically, the temperature condition is important to allow the microstructure of the Al alloy subjected to the solution treatment step to be into isometric crystal grains. If the hot forging temperature is less than 280°C, a crack readily occurs in the Al alloy during the hot forging, and therefore forging itself is not adequately performed. If the hot forging temperature exceeds 430°C, coarse crystal grains are readily formed in the structure of the A1 alloy. As a result, the high-temperature characteristics of the Al alloy are degraded, and therefore Al alloy having excellent high-temperature characteristics cannot be reproducibly manufactured.
Consequently, the hot forging is performed at a temperature of 280 to 430°C.

(Forging Ratio)

The microstructure of the Al alloy subjected to the solution treatment is greatly affected by the forging ratio in the hot forging. Hence, the forging ratio is preferably adjusted to be 1.5 or more in order to allow the microstructure of the Al alloy subjected to the solution treatment to be into isometric crystal grains. If the forging ratio is less than 1.5, the structure of the Al alloy is readily formed into mixed grains. The forging is preferably performed in at least two different directions, in which the forging ratio is 1.5 or more in each direction, rather than in one direction.

(Solution Treatment Step and Quenching Treatment Step)

The solution treatment and the quenching treatment are now described. The solution treatment step and the quenching treatment step are each preferably performed within the conditions specified by JIS-H-4140 or AMS-H-6088 in order to redissolve a soluble intermetallic compound and suppress reprecipitation of the soluble intermetallic compound during the cooling as much as possible. However, even if the heat treatment is performed in accordance with the standard of, for example, AMS-H-6088, and if the solution treatment temperature is excessively high, burning occurs, leading to significant degradation in mechanical characteristics. Conversely, if the solution treatment temperature is less than the lower limit temperature, the yield strength of the Al alloy subjected to the artificial aging treatment is not high enough for the object of the invention, and the solution treatment itself is not adequately performed. Consequently, the upper limit and the lower limit of the solution treatment temperature are specified to be 545°C and 510°C, respectively.
A batch furnace, a continuous annealing furnace, a fused salt bath furnace, an oil furnace, etc. can be appropriately used as a furnace for the temper (heat treatment) such as the solution treatment and the quenching treatment. Various cooling techniques such as water immersion, warm-water immersion, boiled-water immersion, polymer-liquid immersion, water injection, and air injection can be appropriately selectively used for the quenching. Polymers used in the polymer-liquid immersion may include polyoxyethylene-propylene polyether, for example, UCON Quenchant^® from U.S. Union Carbide Corporation.

(Cold Compression (Working) Step)

After the quenching of the Al alloy, cold compression (working) may be performed with a cold roller, a stretcher, a cold forging machine, etc. for straightening after the quenching and improvement in high-temperature characteristics such as yield. strength and creep rupture strength of a final product.
A low compression (working) level of the cold compression (working) results in an insufficient effect of a reduction in residual stress. On the other hand, if the compression (working) level of the cold compression (working) is high, the precipitation amount of the θ' phase increases during the artificial aging treatment or during use of an Al alloy product at high temperature, leading to easy reduction in yield strength. Consequently, the cold compression (working) is preferably performed at a compression (working) rate of 1 to 5%.

(T6 Temper)

In applications of such as small components having a diameter of up to about 100 mm and a piston, even if residual stress is relatively large, the residual stress is not problematic in working, e.g., cutting, for some product. In a case of such a product, the Al alloy is desirably subjected to the artificial aging treatment after the solution treatment and the quenching treatment so as to be formed into a T6 tempered material. In such a case, the temperature of the quenching treatment is desirably 50°C or less in order to achieve high strength characteristics and high high-temperature characteristics even if the residual stress is relatively large.

(T61 Temper)

In a large product such as a rotor, cooling rate in the quenching treatment is greatly different between a surface and the center of the product. This leads to generation of a large residual stress of more than 10 kgf/mm² on a product surface. When such a large residual stress is generated, large strain is generated during the cutting of the product, which extremely makes it difficult to perform precise cutting. In the worst case, the Al alloy material may be fractured, e.g., cracked, due to the residual stress during the cutting. Even if the Al alloy material is not fractured (e.g., cracked), a microcrack may be initiated at an intermetallic compound such as a crystallized grain remaining in the material, or at a slight surface flaw that occurs during conveyance of a product. Such a microcrack is readily propagated and grown during the long-term usage of an Al alloy product, possibly leading to final failure of the Al alloy material.
Hence, in the case of the product such as the rotor for which residual stress is problematic, it is preferred that the Al alloy material is subjected to water quenching at a relatively high temperature of 90°C or more after the solution temperature in order to remove the residual stress or reduce the residual stress to 3.0 kgf/mm² or less, and then the Al alloy material is subjected to the artificial aging treatment to be into a T61 tempered material.

(T652 Temper)

In some application, a product is strictly controlled in residual stress regardless of size of the product. For such a product, preferably, the Al alloy is subjected to cold compression (working) to reduce the residual stress to the utmost, so that the residual stress is preferably removed or reduced to 3.0 kgf/mm² or less, and the Al alloy is then subjected to the artificial aging treatment to be into a T652 tempered material.
For the product, the Al alloy material is preferably subjected to quenching at a temperature of 50°C or less in order to remove the residual stress or reduce the residual stress to 3.0 kgf/mm² or less to achieve sufficient high high-temperature characteristics.
A low compression (working) level of the cold compression (working) results in an insufficient effect of a reduction in residual stress. On the other hand, if the compression (working) level of the cold compression (working) is high, the precipitation amount of the θ' phase increases during the artificial aging treatment or during use at high temperature, and therefore yield strength is readily reduced. Consequently, the cold compression (working) is preferably performed at a compression (working) rate of 1 to 5%.

(Artificial Aging Treatment Step)

The artificial aging treatment in each type of temper is performed to allow the Al alloy to have sufficient normal-temperature yield strength, sufficient high-temperature yield strength, sufficient high-temperature characteristics such as creep rupture strength, and sufficient metal fatigue characteristics. Through the artificial aging treatment, the Ω phase, which is precipitated on the (111) plane of the Al alloy, and the θ' phase, which is precipitated on the (100) plane thereof, can be precipitated, so that the Al alloy exhibits the above-described characteristics. Any of processes of the artificial aging treatment can be used without limitation, as long as the process allows the Al alloy as claimed in this application to include the Ω phase and the θ' phase each being in a precipitation state satisfying the specification of this application, and to have the normal-temperature yield strength, the high-temperature yield strength, the high-temperature characteristics such as creep rupture strength, and the metal fatigue characteristics satisfying the specification of this application.

[Examples]

The mode for carrying out the present invention has been described hereinbefore. Examples that reveal the effects of the invention are now specifically described in comparison with comparative examples that do not satisfy the requirements of the invention. It is assumed that a sample having a room-temperature yield strength of 400 MPa or more, a high-temperature yield strength of 300 MPa or more, and a high-temperature creep rupture strength of 150 MPa or more is acceptable, and other samples are unacceptable.
Checking results on yield strength and creep rupture strength and checking results on metal fatigue strength are now described as a first Example and a second Example, respectively, with inventive examples 1 to 18 and comparative examples 1 to 7.
The invention should not be limited to such Examples.

[First Example]

(Specimen)

Specimens were fabricated as described below to understand influence of a composition of the Al alloy on room-temperature yield strength, high-temperature yield strength, and high-temperature creep rupture strength. Al alloy ingots (diameter: 500 mm, length: 500 mm) having compositions shown in Table 1 were fabricated by casting, and then the ingots were subjected to homogenization heat treatment at 510°C for 20 hr, and were then subjected to hot forging (280 to 360°C, forging ratio: 1.5 or more) to be into forged materials 150 mm square and forged materials 80 mm square. Subsequently, all the forged materials were subjected to solution treatment at 528°C for 6 hr in an air furnace. In Table 1, when a value does not satisfy the range of the invention, the value is underlined.
Subsequently, the forged materials 150 mm square were subjected to warm-water quenching at 70 to 100°C after the solution treatment to reduce residual stress to 3.0 kgf/mm² or less, as a simulated application in which residual stress was problematic. Subsequently, the forged materials were subjected to artificial aging treatment at 175°C for 18 hr to be into T61 tempered materials. Specimens were then fabricated from the tempered materials (the inventive examples 1 to 7, the inventive examples 10 to 12, and the comparative examples 1 to 7).
Some of the forged materials 150 mm square were subjected to water quenching at 30 to 60°C after the solution treatment, and then were subjected to cold compression (working) at a cold compression (working) rate of 1.5% to reduce residual stress to 3.0 kgf/mm² or less, as a simulated application in which residual stress was problematic. Subsequently, such forged materials were subjected to artificial aging treatment at 175°C for 18 hr to be into T652 tempered materials.
Specimens were then fabricated from the tempered materials (the inventive example 9).
The forged materials 80 mm square were subjected to cold-water quenching at 30 to 45°C after the solution treatment, as a simulated application including small components and a piston, in which residual stress was allowed to be relatively large. Subsequently, the forged materials were subjected to artificial aging treatment at 175°C for 18 hr to be into T6 tempered materials. Specimens were then fabricated from the tempered materials (the inventive example 8).

(Room-Temperature Characteristics and High-Temperature Characteristics)

Measurement was made on 0.2% yield strength at room temperature as a room-temperature characteristic of each specimen, and on 0.2% yield strength after the specimen was held at high temperature, i.e., 180°C, for 100 hr as a high-temperature characteristic thereof. Such specimens were prepared with a diameter of a parallel portion of 6.35 mm and a gage length of 25 mm for the room-temperature test, and with a diameter of a parallel portion of 6 mm and a gage length of 30 mm for the 180°C test.
For the creep characteristic test, specimens were prepared with a diameter of a parallel portion of 6 mm and a gage length of 30 mm. Table 2 shows measurement results of tensile characteristics and creep characteristics of such specimens.

Furthermore, elongation of each specimen at room temperature was measured as a room-temperature characteristic of the specimen. The elongation was measured in the same way as the measurement of the 0.2% yield strength, where each specimen was prepared with a diameter of a parallel portion of 6.35 mm and a gage length of 25 mm. Measurement results of the elongation are described later.

Table 2

	Composition No.	Thickness (mm)	Temper	Room-temperature characteristics 0.2% Yield strength (MPa)	High-temperature characteristics
	Composition No.	Thickness (mm)	Temper	Room-temperature characteristics 0.2% Yield strength (MPa)	0.2% Yield strength (MPa)	Creep rupture (MPa)
Inventive example 1	1	150	T61	435	325	160
Inventive example 2	2	150	T61	430	330	160
Inventive example 3	3	150	T61	435	330	160
Inventive example 4	4	150	T61	420	315	155
Inventive example 5	5	150	T61	460	340	170
Inventive example 6	6	150	T61	440	330	155
Inventive example 7	7	150	T61	440	330	160
Inventive example 8	7	80	T6	470	355	165
Inventive example 9	7	150	T652	430	320	155
Inventive example 10	8	150	T61	445	335	160
Inventive example 11	9	150	T61	445	340	160
Inventive example 12	10	150	T61	455	345	165
Comparative example 1	11	150	T61	380	280	140
Comparative example 2	12	150	T61	390	290	140
Comparative example 3	13	150	T61	380	290	140
Comparative example 4	14	150	T61	390	295	145
Comparative example 5	15	150	T61	395	295	135
Comparative example 6	16	150	T61	350	260	130
Comparative example 7	17	150	T61	345	260	130

(Inventive Examples 1 to 12)

It was known that the inventive examples 1 to 12 each had excellent physical characteristics in any of measurement items of room-temperature yield strength, high-temperature yield strength, and high-temperature creep rupture strength, compared with the comparative examples 1 to 7. When the material of such an inventive example is used for high-speed moving components to be rotated or slid at high speed, such as a spinning rotor, a spinning impeller, or a piston, the material exhibits more excellent characteristics than those of any existing material.
Through comparison of room-temperature yield strength values between the inventive example 6 and the inventive example 7, it was known that when the yield strength values were equal to each other, the added amount of Cu in the Al alloy was able to be decreased by increasing the added amount of Si therein. Furthermore, the inventive examples 6 and 7 showed elongations at room temperature of 8.5% and 10.0%, respectively. These results show that Cu-related intermetallic compounds can be decreased while a certain yield strength value is maintained. Such a result further teaches that the Al alloy can be increased in elongation and in metal fatigue strength. Specifically, the amount of (the ratio between) Si and Cu is appropriately adjusted, thereby the metal fatigue strength and the elongation of the Al alloy can be further increased while a certain yield strength value is maintained.

(Comparative Examples 1, 5, and 6)

In the comparative examples 1, 5, and 6, the content of Si was small, i.e., 0.06 mass%; hence, the effect of refining precipitates and the effect of uniform precipitation were small.
Each of the comparative examples 1, 5, and 6 therefore did not satisfy any of the room-temperature yield strength of 400 MPa or more, the high-temperature yield strength of 300 MPa or more, and the high-temperature creep rupture strength of 150 MPa or more.

(Comparative Examples 3, 4, 6, and 7)

In the comparative examples 3, 4, 6, and 7, the content of Zr was large, i.e., 0.15 mass%; hence, a stable phase including AlCu₂ was coarsely precipitated in the periphery of each of the Al-Zr-based dispersed particles during the quenching treatment after the solution treatment. Each of the comparative examples 3, 4, 6, and 7 therefore did not satisfy any of the room-temperature yield strength of 400 MPa or more, the high-temperature yield strength of 300 MPa or more, and the high-temperature creep rupture strength of 150 MPa or more.

(Reference Example)

In some specimens, the content of an alloy component exceeded the upper limit of the specified value, and the following phenomena were shown.
In a specimen where the content of Si, as a component of a comparative example, exceeded 1.0 mass%, a coarse intermetallic compound was formed in the Al alloy, and thus metal fatigue strength was reduced. In a specimen where the content of Cu exceeded 7.0 mass%, strength of the Al alloy excessively increased, and thus forgeability of the Al alloy was degraded. Furthermore, in a specimen where the content of Mn exceeded 1.5 mass%, an insoluble intermetallic compound was formed during the melting and casting, and thus forgeability of the Al alloy was degraded. In a specimen where the content of Mg exceeded 2.0 mass%, strength of the Al alloy increased, and thus forgeability of the Al alloy was degraded.

[Second Example]

Specimens were fabricated according to a procedure described below in order to understand influence of a difference in size between the intermetallic compounds on the material characteristics of the Al alloy.

(Specimen)

An Al alloy ingot having a composition shown in the composition 7 in Table 1 was fabricated by casting such as sand casting, copper mold casting, or continuous casting while conditions such as casting rate were adjusted. Subsequently, the Al alloy ingot was subjected to homogenization heat treatment at 510°C for 20 hr, and was then subjected to hot forging (280 to 360°C, forging ratio: 1.5 or more) to be into forged materials 150 mm square.
Such fabricated forged materials 150 mm square were subjected to solution treatment at 528°C for 6 hr in an air furnace, and were then subjected to warm-water quenching at 70 to 100°C. Subsequently, the forged materials were subjected to artificial aging treatment at 175°C for 18 hr to be into T61 tempered materials.

(Size of Intermetallic Compound)

Size of the intermetallic compound was analyzed using extremal statistics analysis as described before. Specifically, 40 optical photomicrographs of 100 magnifications were taken, and a largest intermetallic compound was extracted in a region of the inspecting standard area, 0.37 mm², for each photomicrograph. The resultant 40 data were subjected to the statistics processing, and the largest circle-equivalent diameter of the intermetallic compound in set predictive area was estimated for each material. In this measurement, the predictive area was set as 100 mm².

(Rotary Bending Fatigue Strength Test)

Test pieces were fabricated from the above-described T61 tempered materials, and the test pieces were subjected to a metal fatigue strength test at a high temperature of 150°C (maximum stress: 130 MPa, stress ratio: -1).
In this metal fatigue test, round-bar test pieces were subjected to a rotary bending fatigue strength test, each round-bar test piece having a diameter of a parallel portion of 6 mm and a length of the parallel portion of 13.55 mm, and being finished with emery paper #1000.

Table 3 shows a relationship between the largest circle-equivalent diameter and a rupture repetition number as a measurement result of the rotary bending fatigue strength test. The rupture repetition number refers to a repetition number before rupture in the rotary bending fatigue test. The rupture repetition number is represented by a relative value assuming that a value at a largest circle-equivalent diameter of 90 µm is 1.

Table 3

	Composition No.	Thickness (mm)	Temper	Rotary bending fatigue strength test (150°C)
	Composition No.	Thickness (mm)	Temper	Largest circle-equivalent diameter (µm)	Rupture repetition number *
Inventive example 13	7	150	T61	90	1.0
Inventive example 14	7	150	T61	81	1.9
Inventive example 15	7	150	T61	60	4.0
Inventive example 16	7	150	T61	48	7.5
Inventive example 17	7	150	T61	39	9.2
Inventive example 18	7	150	T61	25	19.1

* A value at a largest circle-equivalent diameter of 90 µm is assumed to be 1.

(Metal Fatigue Strength)

Materials, which were adjusted in largest circle-equivalent diameter of the intermetallic compound, were subjected to the rotary bending fatigue strength test. As a result, it was known that in the case of alloys (test pieces) having the same composition (composition 7 in Table 1), the largest circle-equivalent diameter was adjusted to be 60µm or less in this predictive area (100 mm²), thereby the metal fatigue strength was increased. In addition, the metal fatigue strength tended to be increased with a further reduction in largest circle-equivalent diameter.

(Inventive Examples 13 to 18)

As shown in Table 3, it is known that the test pieces fabricated from the Al alloy ingot having the composition shown in the composition 7 in Table 1 are adjusted in largest circle-equivalent diameter to be 60 µm or less (the inventive examples 15 to 18), thereby the rupture repetition number becomes four times or more as large as that of the test piece having the largest circle-equivalent diameter of 90 µm (the inventive example 13).
In addition, the largest circle-equivalent diameter is adjusted to be 40 µm or less (the inventive example 17), thereby the rupture repetition number becomes nine times or more as large as that of the test piece having the largest circle-equivalent diameter of 90 µm (the inventive example 13), allowing further increase in metal fatigue strength of the Al alloy. In addition, the largest circle-equivalent diameter is adjusted to be 25 µm or less (the inventive example 18), thereby the rupture repetition number becomes 19 times or more as large as that of the test piece having the largest circle-equivalent diameter of 90 µm (the inventive example 13), allowing remarkable increase in metal fatigue strength of the Al alloy.
In this way, it is found that each of the test pieces having the largest circle-equivalent diameter of 60 µm or less (the inventive examples 15 to 18) has excellent physical characteristics compared with the test piece having the largest circle-equivalent diameter of 90 µm (the inventive example 13). When such materials are each used for high-speed moving components, on which intermittent stress is exerted, such as a spinning rotor, a spinning impeller, or a piston, the materials each exhibit more excellent characteristics than those of any existing material.

Industrial Applicability

The Al alloy of the present invention is excellent in heat resistance, high-temperature fatigue strength, creep resistance under high temperatures, and high-temperature yield strength, and is applicable particularly for Al alloy components to be used in service environment at a high temperature of more than 100°C, such as rockets, aircrafts, and space equipment, transport equipment such as railway vehicles, motor vehicles, and ships, engine components, components of machines such as a compressor, and a spinning rotor, a spinning impeller, or a piston.

Claims

An aluminum alloy having excellent high-temperature characteristics, comprising:
Si of more than 0.1 to 1.0 mass%;

Cu of 3.0 to 7.0 mass%;

Mn of 0.05 to 1.5 mass%;

Mg of 0.01 to 2.0 mass%;

Ti of 0.01 to 0.10 mass%;

Ag of 0.05 to 1.0 mass%;

Zr in an limited content of less than 0.1 mass%; and

the remainder consisting of Al and inevitable impurities.
The aluminum alloy according to claim 1, further comprising V of 0.15 mass% or less.
The aluminum alloy according to claim 1 or 2, characterized in that a largest circle-equivalent diameter of an intermetallic compound contained in the aluminum alloy is 60 µm or less.