JP6490608B2 - Method for producing Cu-Al-Mn alloy material - Google Patents

Description

The present invention relates to a method for producing a Cu-Al-Mn alloy material having excellent fracture resistance when subjected to repeated deformation.

  Shape memory alloys and superelastic alloys such as copper alloys show remarkable shape memory effects and superelastic properties accompanying the reverse transformation of the thermoelastic martensitic transformation, and have excellent functions near the living environment temperature. Have been put to practical use in various fields. Typical materials for shape memory alloys and superelastic alloys include TiNi alloys and copper (Cu) alloys. Copper-based shape memory alloys and superelastic alloys (hereinafter also referred to simply as copper-based alloys) are inferior to TiNi alloys in terms of repeatability, corrosion resistance, etc. There is a movement to expand its coverage for cheap. However, the copper-based alloy is advantageous in terms of cost, but has poor cold workability and low superelastic characteristics. For this reason, despite various researches, copper-based alloys are not necessarily sufficient for practical use.

  So far, various studies have been made on copper-based alloys. For example, the following patent documents 1 to 4 report Cu-Al-Mn shape memory alloys having a β single phase structure excellent in cold workability. In these examples, for example, with respect to the crystal orientation, a recrystallized texture in which the β single-phase metal structure is aligned in a specific direction such as <101>, <100>, etc., in a cold working direction such as rolling or wire drawing. ing.

Japanese Patent Laid-Open No. 7-62472 JP 2000-169920 A Japanese Patent Laid-Open No. 2001-20026 International Publication WO2011 / 152009A1 International Publication WO2015 / 137283A1

Proceedings of the Japan Society of Materials Science Lecture 45, pp. 169-170, 1996 “Superelastic cyclic behavior of CuAlMn shape memory alloy single crystals”

  The Cu—Al—Mn alloy produced by the method of Patent Document 1 does not have sufficient characteristics, particularly superelasticity, and the maximum strain showing a shape recovery of 90% or more is about 2-3%. The reason is considered to be that irreversible defects such as dislocations are introduced because a strong restraining force is generated between crystal grains during deformation due to random crystal orientation. Therefore, a large amount of residual strain is accumulated due to repeated deformation, and the superelastic characteristics are significantly deteriorated after repeated deformation. In addition, since the fracture resistance when repeatedly deformed is low and cannot be controlled, the fracture often occurs several tens of times.

  Further, the copper-based alloy of Patent Document 2 is a copper-based alloy having shape memory characteristics and superelastic characteristics and substantially composed of β single phase, and the crystal structure of the β single phase is β single phase. <101>, <100>, etc. have a recrystallized texture in which specific crystal orientations are aligned in the cold working direction such as rolling or wire drawing. The copper-based alloy is also referred to as an electron backscatter diffraction patterning method (hereinafter abbreviated as “EBSP”) (or an electron backscatter diffraction (hereinafter abbreviated as EBSD)). The cold working is performed at a total working rate after the final annealing such that the existence frequency of the specific crystal orientation of the β single phase in the working direction is 2.0 or more. Even in such a material, in the Cu-Al-Mn alloy, the orientation dependence of the transformation strain is large, so that it is still insufficient to obtain a stable and excellent superelastic property with high accuracy and uniformity. It is. In addition, there is a large amount of residual strain accumulated by repeated deformation, and the superelastic characteristics are significantly deteriorated after repeated deformation. Similarly to Patent Document 1, the resistance to rupture in the case of repeated deformation is low and cannot be controlled, and the rupture often occurs several tens of times.

  Further, in the copper-based alloys described in Patent Document 3 and Patent Document 4, there is still room for improvement in that the performance of the shape memory characteristics and the superelastic characteristics that are manifested is greatly uneven and these characteristics are not stable. Is a level. In addition, it is considered that texture control is indispensable for stabilizing shape memory characteristics and superelastic characteristics. However, in the method described in Patent Document 3, the degree of organization of the structure in the Cu—Al—Mn alloy is as follows. The low shape memory and superelastic properties are not yet stable enough. In Patent Document 3, in order to improve the shape memory characteristics and superelastic characteristics of a copper-based alloy, the crystal orientation to the β single phase is controlled, and if the average crystal grain size is a wire, it is set to more than half the wire diameter. Alternatively, in the case of a plate material, it is proposed that the thickness be equal to or greater than the plate thickness, and that the region having such a crystal grain size be 30% or more of the entire length of the wire or the total area of the plate material. Further, in Patent Document 4, in order to improve the shape memory characteristics of a copper-based alloy and to make a copper-based alloy having a large cross-sectional size applicable to a structure, a giant crystal grain having a maximum crystal grain size exceeding 8 mm is used. Propose to be an organization. However, in the methods described in Patent Document 3 and Patent Document 4, control of the particle size distribution of crystal grains having a predetermined large crystal grain size in the Cu—Al—Mn-based alloy is still insufficient. The effect and superelastic properties are not stable. In addition, there is a large amount of residual strain accumulated by repeated deformation, and the superelastic characteristics are significantly deteriorated after repeated deformation. In either case, the resistance to fracture when subjected to repeated deformation is low and cannot be controlled.

  In Patent Document 5, it is possible to control the residual strain amount after repeated deformation by defining the boundary between large crystal grains and small crystal grains and obtaining the existence frequency of smaller crystal grains. It is similar to the present invention in that the larger the grain size, the better the characteristics. However, a significant difference is that the present invention can control the fracture resistance to a high degree when it is repeatedly deformed. As a result, it has been clarified that the present invention makes it possible to control the crystal grain size higher than that of Patent Document 5, and exhibits particularly excellent characteristics.

  Non-Patent Document 1 aims to improve repetitive deformation. Here, as a manufacturing method, a test piece is manufactured and evaluated by an industrially difficult method called a vertical Bridgman method which requires a lot of manufacturing time. The size of the test piece is also as small as 2 mm × 2 mm × 4 mm, and there is a big problem that the applicable field is limited with this size no matter how excellent the repeated deformation is. The vertical Bridgman method is a manufacturing method in which a hot zone is constituted by resistance heating and a heat insulating material, and the temperature is gradually lowered by pulling down the crucible to cause crystallization in the crucible. However, there is a high possibility that impurities are mixed from the crucible, and this causes a different crystal orientation to grow and polycrystallize easily, so that there is a problem that the desired characteristics are often not obtained. Furthermore, since the reported composition is a composition that cannot be processed after production, the gist of the composition differs greatly from the industrial Cu—Al—Mn alloy required in the present invention.

  Thus, it is considered that accumulation of crystal orientations or having a predetermined large crystal grain size is effective in improving superelasticity in a Cu-Al-Mn alloy. However, the prior art has not been improved due to low fracture resistance when repeated deformation is performed. However, when this alloy is used as a medical instrument, a building member, or the like, deterioration of characteristics due to repeated deformation becomes a serious problem and improvement is required. Furthermore, in order to use a Cu-Al-Mn alloy material as a vehicle-mounted part, an aerospace equipment part, etc., there is an effect of further preventing the fracture even after a high cycle in repeated deformation and the deterioration of superelastic characteristics. It has been demanded.

  Then, this invention makes it a subject to provide the Cu-Al-Mn type-alloy material which was excellent in the fracture resistance at the time of repeating a deformation | transformation, and was excellent, and a use using the same.

  As a result of intensive studies to solve the above-mentioned problems, the inventors of the present invention controlled the number of crystal grains of the Cu-Al-Mn alloy material, thereby preventing fracture when repeated deformation was performed. It has been found that an excellent alloy material can be obtained. In addition, the control that enables such a crystal grain size is performed after a predetermined intermediate annealing and cold working, and also in a state of (α + β) phase in which the α phase precipitation amount is fixed in the first stage of the memory heat treatment. After heating to a temperature range that becomes a β single phase at a specific slow temperature increase rate, hold for a predetermined time at a predetermined temperature, and further, from a temperature range that becomes a β single phase to a temperature range that becomes an (α + β) phase Performing a storage heat treatment in which cooling at a slow temperature-decreasing rate and heating at a specific slow temperature-raising rate from the temperature range of the (α + β) phase to the temperature range of the β single phase are repeated at least four times; and (α It has been found that this can be achieved by controlling both of the temperature range to be particularly limited when the state is + β). The present invention has been completed based on these findings.

According to the present invention, the following means are provided .
(1 ) 3.0-10.0 mass% Al, 5.0-20.0 mass% Mn, Ni, Co, Fe, Ti, V, Cr, Si, Nb, Mo, W, Sn, 0.000 to 10.000 mass% in total of one or more selected from the group consisting of Mg, P, Be, Sb, Cd, As, Zr, Zn, B, C, Ag and Misch metal Here, the content of Ni and Fe is 0.000 to 3.000 mass%, the content of Co is 0.000 to 2.000 mass%, and the content of Ti is 0.00. The content of V, Nb, Mo, and Zr is 0.000 to 1.000% by mass, and the content of Cr is 0.000 to 2.000% by mass. , Si content is 0.000 to 2.000 mass%, W content is 0.00 0 to 1.000 mass%, the Sn content is 0.000 to 1.000 mass%, the Mg content is 0.000 to 0.500 mass%, and the P content is 0. The content of Be, Sb, Cd, and As is 0.000 to 1.000 mass%, and the content of Zn is 0.000 to 5.000 mass%. Yes, the contents of B and C are 0.000 to 0.500% by mass, the content of Ag is 0.000 to 2.000% by mass, and the content of misch metal is 0.000 to 5%. A step of melting and casting a material of a Cu-Al-Mn alloy having a composition of 0.000% by mass and the balance of Cu and inevitable impurities;
A hot working process;
A step of performing intermediate annealing at 400 to 680 ° C. for 1 to 120 minutes and cold working at a processing rate of 30% or more at least once each in this order;
After heating from room temperature to a temperature range of 400 to 650 ° C. that becomes an (α + β) phase, hold in the temperature range for 2 to 120 minutes, and from a temperature range that becomes an (α + β) phase to a temperature range that becomes a β single phase 0.1 to Heat at a rate of temperature increase of 20 ° C./min and hold in the temperature range for 5 to 480 minutes, and then change from a temperature range of β single phase to a temperature range of 400 to 650 ° C. of (α + β) phase. Cool at a rate of temperature decrease of 20 ° C./min and hold in the temperature range for 20 to 480 minutes, then 0.1 to 20 from the temperature range 400 to 650 ° C. in which the (α + β) phase is reached to the temperature range in which the β single phase is reached. Heated at a temperature increase rate of ° C./min and held in the temperature range for 5 to 480 minutes, then rapidly cooled,
Here, from the step of maintaining in the temperature range that becomes the β single phase, from the temperature range that becomes the β single phase to the temperature range 400 to 650 ° C. that becomes the (α + β) phase, 0.1 to 20 ° C./min. The temperature is increased at a rate of 0.1 to 20 ° C./min from the temperature range that becomes the (α + β) phase to the temperature range that becomes the β single phase through the process of cooling at the temperature decrease rate and holding in the temperature range for 20 to 480 minutes. The process is repeated at least 4 times or more until it is heated at 5 to 480 minutes in the temperature range.
The alloy material is an alloy material having a long shape with respect to a processing direction which is a rolling direction or a wire drawing direction,
The crystal grain length a of the alloy material in the direction perpendicular to the processing direction is equal to the width or thickness or diameter R of the alloy material, and X is the number of grain boundaries between crystal grains where a = R. The amount of X present is 1 or less,
Method for manufacturing a Cu-Al-Mn-based alloy material number until the fracture is more than 10 ^twice when performed repeatedly load and unloading of the stress applied 3% strain in the alloy material.
The following Cu—Al—Mn alloy material and its application will be described in more detail with respect to the above-described embodiment.
<1> 3.0 to 10.0% by mass of Al, 5.0 to 20.0% by mass of Mn, Ni, Co, Fe, Ti, V, Cr, Si, Nb, Mo, W, Sn, 0.000 to 10.000 mass% in total of one or more selected from the group consisting of Mg, P, Be, Sb, Cd, As, Zr, Zn, B, C, Ag and Misch metal Here, the content of Ni and Fe is 0.000 to 3.000 mass%, the content of Co is 0.000 to 2.000 mass%, and the content of Ti is 0.00. The content of V, Nb, Mo, and Zr is 0.000 to 1.000% by mass, and the content of Cr is 0.000 to 2.000% by mass. , Si content is 0.000 to 2.000 mass%, W content is 0.00 To 1.000 mass%, the Sn content is 0.000 to 1.000 mass%, the Mg content is 0.000 to 0.500 mass%, and the P content is 0.00. The content of Be, Sb, Cd, and As is 0.000 to 1.000 mass%, and the content of Zn is 0.000 to 5.000 mass%. , B, and C are each 0.000 to 0.500 mass%, Ag is 0.000 to 2.000 mass%, and the misch metal content is 0.000 to 5.0. A Cu-Al-Mn alloy material having a composition consisting of Cu and unavoidable impurities,
The alloy material is an alloy material having a long shape with respect to a processing direction which is a rolling direction or a wire drawing direction,
The crystal grain length a of the alloy material in the direction perpendicular to the processing direction is equal to the width or thickness or diameter R of the alloy material, and X is the number of grain boundaries between crystal grains where a = R. The amount of X present is 1 or less,
Cu-Al-Mn-based alloy material, wherein the number of until breakage when subjected repeatedly load and unloading of the stress applied 3% strain in the alloy material is more than 10 ² times.
<2> From Ni, Co, Fe, Ti, V, Cr, Si, Nb, Mo, W, Sn, Mg, P, Be, Sb, Cd, As, Zr, Zn, B, C, Ag, and Misch Metal 1 to 2 or more kinds selected from the group consisting of 0.001 to 10.000% by mass in total, where the contents of Ni and Fe are 0.001 to 3.000% by mass, The Co content is 0.001 to 2.000 mass%, the Ti content is 0.001 to 2.000 mass%, and the contents of V, Nb, Mo, and Zr are 0.001 to 0.001, respectively. 1.000 mass%, Cr content is 0.001 to 2.000 mass%, Si content is 0.001 to 2.000 mass%, and W content is 0.001. To 1.000 mass%, and the Sn content is 0.001 to 1.000 mass The Mg content is 0.001 to 0.500 mass%, the P content is 0.010 to 0.500 mass%, and the contents of Be, Sb, Cd, and As are each 0. 0.001 to 1.000% by mass, Zn content is 0.001 to 5.000% by mass, B and C contents are 0.001 to 0.500% by mass, respectively. The Cu—Al—Mn based alloy material according to <1>, wherein the content is 0.001 to 2.000 mass%, and the content of misch metal is 0.001 to 5.000 mass%.
<3> In the cyclic deformation in which stress loading and unloading corresponding to a 5% strain amount are repeated 100 times, the remaining strain amount is 1.5% or less, and Cu— according to <1> or <2> Al-Mn alloy material.
<4> A spring material made of the Cu—Al—Mn alloy material according to any one of <1> to <3>.
<5> A damper made of the Cu—Al—Mn alloy material according to any one of <1> to <3>.
<6> A brace made of the Cu—Al—Mn alloy material according to any one of <1> to <3>.
<7> A screw or bolt made of the Cu—Al—Mn alloy material according to any one of <1> to <3>.
<8> A member made of the Cu—Al—Mn alloy material according to any one of <1> to <3>, having a mechanism for repeating tension and compression.

Here, the number of times until fracture when repeated deformation is performed refers to the number of times of repetition when deformation is performed until a fracture occurs by repeating loading and unloading with a predetermined strain amount. In the present invention, the larger the number of times, the better. In the present invention, the excellent times until breakage when performing repeated deformation, the number of until it breaks is not less than 10 ² times the case of performing the load and unloading of the stress giving the distortion of 3% That means. Furthermore, it is preferable that this number does not vary.
Here, the fact that there is no variation is, for example, that N = 5 times of the same test piece for each manufacturing condition. ^three times or more (measured at 5000 times the end) is excellent especially if, determined over all 10 ^twice (the measurement of N = 5, the minimum value is 971 times, the maximum value is 5000 times) good if it can. On the other hand, a portion of the 5 measurements be unbroken at least 10 ³ times, if you break less than even one 10 ² times if there are variations in the number of times to reach the fracture are considered poor.

Cu-Al-Mn superelastic alloy materials can be used for various applications that require superelastic properties. For example, in addition to mobile phone antennas and eyeglass frames, orthodontic wires and guides can be used as medical products. Application to wires, stents, ingrown nail correctors (ingrown nail correctors), hallux valgus prostheses, connectors, and actuators is being studied. Among them, the Cu-Al-Mn superelastic alloy material of the present invention has a high fracture resistance when subjected to repeated deformation, and therefore is excellent in a member intended for damping or damping related to vibration, noise suppression or It is suitable for a member for the purpose of damping or a member for the purpose of self-restoration (self-centering). In particular, it has become possible to use in fields that have been difficult in the past, such as space equipment, aeronautical equipment, automobile parts, building parts, electronic parts, medical products and the like that require resistance to repeated deformation. For example, vibration is suitable as a damping material such as a bus bar or a building material. Moreover, a damping structure or the like can be constructed using this damping material or building material. Furthermore, it can also be used as a civil engineering building material that can prevent noise and vibration pollution by utilizing the characteristics of absorbing vibration as described above. Furthermore, in the case of aiming at the effect of noise attenuation, it can be applied in the field of transportation equipment. In any case, since it has an excellent self-restoring force, it can be used as a self-restoring material.
The above and other features and advantages of the present invention will become more apparent from the following description, with reference where appropriate to the accompanying drawings.

FIG. 1 is a schematic view of a Cu—Al—Mn alloy rod (wire) 1 according to the present invention, in which the grain length (a) in a direction perpendicular to the processing direction (RD) of the alloy material defined in the present invention is shown. FIG. 6 is a schematic diagram illustrating the relationship between the width and the diameter (R) of the alloy material and the number of boundaries (grain boundaries) (X) between grains including a. FIG. 2 is a flowchart showing all the steps in the manufacturing method of the present invention. The name of each process is shown together with the flowchart. FIG. 3 is a schematic diagram illustrating the definition of each physical property value exhibited by the Cu—Al—Mn alloy material of the present invention. FIG. 3A is a schematic diagram of an SS curve when deformed 5000 times by repeated stress loading and slow loading corresponding to 3% strain. FIG. 3 (b) shows the time when the first cycle was completed (solid line in the figure) and the 1000th cycle were completed after 1000 cycles of a test in which tension and compression at a stress equivalent to 3% strain were repeated. It is a schematic diagram of each SS curve at the time (dotted line in a figure). The residual strain at the completion of the 1000th cycle is shown in the figure. FIG. 3 (c) shows the time when the first cycle was completed (solid line in the figure) and the time when the 100th cycle was completed (in the figure) after the test of repeating tension and compression at a stress equivalent to 3% strain. It is a schematic diagram of each SS curve in a dotted line. The stress values at the time of 3% strain loading with respect to the 0.2% proof stress at the completion of the first cycle and the 100th cycle are shown in the figure. In addition, the rate of decrease in 0.2% yield strength (yield stress) at the completion of the 100th cycle is shown in the figure as compared to 0.2% yield strength (yield stress) at the completion of the first cycle. Here, the repetition of loading and unloading refers to a test method in which stress loading and unloading are performed once to make one cycle. FIG. 4 is a schematic diagram illustrating the definition of each physical property value exhibited by the Cu—Al—Mn alloy material of the present invention. FIG. 4 (a) shows the time at which the first cycle after the test in which 5% strain load unloading was repeated 100 cycles was completed (solid line in the figure) and at the time when the 100th cycle was completed (dotted line in the figure). It is a schematic diagram of an SS curve. The residual strains at the completion of the first time and the 100th time are shown in the figure. FIG. 4B is a schematic diagram of the SS curve after the 5% strain load unloading test. The “stress difference” of the stress value at the time of 5% strain loading with respect to 0.2% proof stress is shown in the figure. FIG. 5 is a schematic diagram of the shape and dimensions of a JIS No. 14 test piece described in “JIS Z2241 Metal Material Tensile Test Method”. In this invention, the test piece was produced based on the shape of a JIS No. 14 test piece, and the characteristic evaluation was implemented. 6A is a flowchart showing a manufacturing process in Example 1 (manufactured in process No. a) and FIG. 6B is a manufacturing process in Comparative Example 1 (manufactured in process No. A). The conditions of processing and heat treatment in each step and the number of repetitions are also shown. In Example 1 (process No. a), the number of repetitions [19] of the gradual temperature decrease [process 5-5] [13] and the gradual temperature increase [process 5-7] [16] in the memory heat treatment was 4 times. On the other hand, in Comparative Example 1 (Step No. A), the slow temperature decrease [Step 5-5] [13] and the gradually temperature increase [Step 5-7] [16] were repeated only twice in this memory heat treatment [19]. Not. In Example 1 (Step No. a), the temperatures of [Step 5-2] [8] and [Step 5-6] [14] are 450 ° C., whereas Comparative Example 1 (Step No. In A), the temperatures of [Step 5-2] [8] and [Step 5-6] [14] are 300 ° C. That is, the number of repetitions [19] and the (α + β) phase holding temperature [8] and [14] are also different. 7 is a schematic diagram of an SS curve obtained when the sample obtained in Example 1 (step No. a) in Table 5 described later is measured by a test in which tension and compression corresponding to 3% strain are repeated. It is. FIG. 7: is a schematic diagram of the SS curve obtained about Example 1 in order from the left after 1 cycle of a 3% tension cycle test and a tension compression test, after 100 cycles, and after 1000 cycles. Here, “repeat load and slow load” and “repeat tension and compression” are different test methods. The foregoing is a test method in which stress load (tensile) and slow load (not compressed) are repeated, and the test method in the following is repeated in tensile load and compressive load. In other words, “repeating tension and compression” is “a test method in which a load in the tensile direction and a load in the compression direction of the stress are performed once to repeat them as one cycle”.

  The Cu—Al—Mn alloy material of the present invention is subjected to a predetermined intermediate annealing and cold working, and further heating before the temperature range to become the first β single phase of the memory heat treatment [Step 5-3]. The temperature range 400 where the (α + β) phase is changed from the temperature range where the β phase is changed to the (α + β) phase after the α phase precipitation amount is fixed by holding in the temperature range 400 to 650 ° C. [step 5-2]. Cooling at a specific slow temperature decrease rate to ˜650 ° C. [Step 5-5] and a temperature range from 400 to 650 ° C. to a temperature range becoming a (α + β) phase to a temperature range becoming a β single phase at a specific slow temperature rising rate A storage heat treatment is performed in which the heating [Step 5-7] is repeated at least four times. Thus, the existence of X when the crystal grain length a in the direction perpendicular to the processing direction is equal to the width or diameter R of the alloy material and the number of grain boundaries between the crystal grains where a = R is X. The amount can be controlled, and an alloy material that exhibits good superelasticity and high fracture resistance even when repeatedly deformed.

  In addition, a process direction (RD, refer FIG. 1) refers to a wire drawing direction if it is a wire drawing process, and refers to a rolling direction if it is a rolling process. Usually, the rolling direction at the time of rolling processing of a plate material or the like is referred to as RD (Rolling Direction), but the wire drawing direction at the time of wire drawing of a bar material or the like may also be conventionally expressed as RD. Therefore, when it is referred to as RD in the present specification, the rolling direction and the wire drawing direction are collectively referred to and the processing direction of a plate material, a bar material (wire material) or the like is meant. In the present invention, the characteristics were evaluated as having a long shape with respect to the processing direction which is the rolling direction or the wire drawing direction.

<Composition of Cu-Al-Mn alloy material>
The copper-based alloy of the present invention having shape memory characteristics and superelasticity is an alloy containing Al and Mn. This alloy becomes a β-phase (body-centered cubic) single phase (also referred to simply as a β-single phase in this document) at high temperatures, and a two-phase structure (in this document, simply a β-phase and face-centered cubic) at low temperatures. (Also referred to as (α + β) phase). Although it depends on the alloy composition, the high temperature at which the β single phase is obtained is usually 700 ° C. or higher, and the low temperature at which the (α + β) phase is obtained is usually less than 700 ° C.

  The Cu—Al—Mn alloy material of the present invention contains 3.0 to 10.0% by mass of Al and 5.0 to 20.0% by mass of Mn, and is composed of the balance Cu and inevitable impurities. Have If the Al element content is too small, a β single phase cannot be formed, and if it is too much, the alloy material becomes brittle. Although content of Al element changes according to content of Mn element, content of preferable Al element is 6.0-10.0 mass%. By containing the Mn element, the existence range of the β phase is expanded to the low Al side, and the cold workability is remarkably improved, so that the forming process is facilitated. If the amount of Mn element added is too small, satisfactory processability cannot be obtained, and a β single phase region cannot be formed. If the amount of Mn element added is too large, sufficient shape recovery characteristics cannot be obtained. A preferable Mn content is 8.0 to 12.0 mass%. The Cu—Al—Mn alloy material having the above composition is rich in hot workability and cold workability, and it is possible to achieve a working rate of 20% to 90% or more in the cold state. In addition, it can be molded into ultrafine wires, foils, pipes and the like that have been difficult to process.

  In addition to the above-mentioned essential additive elements, the Cu—Al—Mn alloy material of the present invention further includes Ni, Co, Fe, Ti, V, Cr, Si, Nb, Mo, W, Sn as optional auxiliary additive elements. Mg, P, Be, Sb, Cd, As, Zr, Zn, B, C, Ag, and one or more selected from the group consisting of misch metal (Pr, Nd, etc.) can be contained. . These elements exhibit the effect of improving the strength of the Cu—Al—Mn alloy material while maintaining cold workability. The total content of these additive elements is preferably 0.001 to 10.000% by mass, particularly preferably 0.001 to 5.000% by mass. If the content of these elements is too large, the martensitic transformation temperature decreases and the β single phase structure becomes unstable.

  Ni, Co, Fe, and Sn are effective elements for strengthening the base structure. Co coarsens crystal grains due to the formation of a Co—Al intermetallic compound, but if excessive, it lowers the toughness of the alloy. The Co content is 0.001 to 2.000 mass%. The contents of Ni and Fe are 0.001 to 3.000% by mass, respectively. Sn content is 0.001-1.000 mass%.

  Ti combines with inhibitory elements N and O to form oxynitrides. Moreover, a boride is formed by combined addition with B, and intensity | strength is improved. The Ti content is 0.001 to 2.000 mass%.

  V, Nb, Mo, and Zr have the effect of increasing the hardness and improve the wear resistance. Moreover, since these elements hardly dissolve in the matrix, they are precipitated as a β phase (bcc crystal) to improve the strength. Content of V, Nb, Mo, and Zr is 0.001-1.000 mass%, respectively.

  Cr is an effective element for maintaining wear resistance and corrosion resistance. The Cr content is 0.001 to 2.000 mass%. Si has the effect of improving the corrosion resistance. The Si content is 0.001 to 2.000 mass%. Since W hardly dissolves in the base, there is an effect of precipitation strengthening. Content of W is 0.001-1.000 mass%.

  Mg has the effect of removing the inhibitory elements N and O, and fixes the inhibitory element S as a sulfide, which is effective in improving hot workability and toughness. Addition of a large amount causes segregation of grain boundaries and causes embrittlement. The content of Mg is 0.001 to 0.500 mass%.

  P acts as a deoxidizer and has the effect of improving toughness. Content of P is 0.01-0.50 mass%. Be, Sb, Cd, and As have the effect of strengthening the base organization. The contents of Be, Sb, Cd, and As are 0.001 to 1.000 mass%, respectively.

  Zn has the effect of increasing the shape memory processing temperature. The Zn content is 0.001 to 5.000% by mass. If B and C are suitable amounts, the pinning effect is obtained and the crystal grains are more coarsened. In particular, combined addition with Ti and Zr is preferable. Content of B and C is 0.001-0.500 mass%, respectively.

  Ag has the effect of improving cold workability. The content of Ag is 0.001 to 2.000 mass%. If the amount of misch metal is an appropriate amount, the pinning effect can be obtained, so that the crystal grains are further coarsened. The content of misch metal is 0.001 to 5.000% by mass. Misch metal refers to an alloy of rare earth elements such as La, Ce, and Nd that are difficult to separate.

<Metal structure of Cu-Al-Mn alloy material>
The Cu—Al—Mn alloy material of the present invention has a recrystallized structure. Moreover, the Cu—Al—Mn alloy material of the present invention has a recrystallized structure substantially consisting of a β single phase. Here, “having a recrystallized structure consisting essentially of a β single phase” means that the proportion of the β phase in the recrystallized structure is usually 90% or more, preferably 95% or more.

In the technical field of the present invention, the technical significance of the present invention is that the Cu—Al—Mn alloy material can be controlled to a predetermined crystal grain size. According to the present invention, not only stably exhibits superelastic properties, but also has a large number of deformations (the number of times until fracture occurs when repeated loading and unloading of a stress that gives a strain of 3% to the alloy material are performed. It is possible to withstand breakage even if it reaches 10 ² times or more (hereinafter simply referred to as the above-mentioned multiple times). Thus, a remarkable effect that cannot be expected from the conventional means is obtained.

In addition, although the bamboo structure was calculated | required by the prior art, only a large crystal grain was controllable and the small crystal grain was not controllable. Therefore, good superelasticity was exhibited in several repeated cycles, but residual strain increased in many cycles. This is because residual strain accumulates at the grain boundaries. For this reason, attempts have been made to reduce the number of small crystal grains as much as possible. By controlling the amount of small crystal grains, it is possible to reduce residual strain even in the above-described many deformations. However, it has been clarified that when there are many large crystal grains that are not small crystal grains and so-called bamboo structure, the number of times until fracture occurs. In other words, not only the presence of small crystal grains but also the bamboo structure has a significant effect on the number of times until it breaks as the amount thereof increases, and when it is repeatedly deformed, it breaks early. Therefore, in the present invention, the abundance of X is controlled to be 1 or less, thereby making it possible to reduce both the residual strain after the numerous repetitions and the increase in the number of times until breakage. Thus, a remarkable effect that cannot be expected from the conventional means is obtained.
The bamboo structure (organization) will be described. Assuming that the number of grain boundaries between crystal grains satisfying a = R (see FIG. 1) is X, if the amount of X present is large, the grain boundaries exist in the test body like bamboo nodes. Here, the bamboo structure or the bamboo structure is a bamboo structure or a bamboo structure in which a number of grain boundaries penetrating the diameter of the specimen exist.

<Definition and control of crystal grain size>
Only large crystal grains exist in the Cu—Al—Mn based copper alloy of the present invention. For example, in the case of a bar, when the crystal grain length a in the direction perpendicular to the processing direction (RD) with respect to the sample diameter R is R = a and the number of grain boundaries between the crystal grains is X, The abundance of X is 1 or less, preferably 0. Similarly, in the case of a plate material, the crystal grain length a in the direction perpendicular to the processing (rolling) direction (RD) with respect to the sample diameter R is R = a, and the number of grain boundaries between the crystal grains is X. In this case, the abundance of X is 1 or less, preferably 0. Here, the abundance of the crystal grain boundary X can be determined by the number of grain boundaries when the surface or cross section in the longitudinal direction of the Cu—Al—Mn alloy material is arbitrarily measured at four or more points. The crystal grain X in the present invention has a Cu-Al-Mn alloy material that is substantially higher in processing than the center due to the effects of additional shear stress and tool surface friction in the processing step, and the crystal grains tend to become finer. Evaluation may be performed on the surface. Or you may measure X of a parallel part after processing into the test piece shape shown in FIG. In the present invention, since the characteristic evaluation was performed as it is after the confirmation, it was processed into the state shown in FIG.
The shape and dimensions of the test piece shown in FIG. 5 are as follows. Original test point distance L ₀ (mm) = 5.66√S ₀ , parallel part length L _c (mm) = 5.5d _{0 to} 7d ₀ , shoulder radius R (mm) = 15 or more .
When the parallel part has a square cross section, L _c = 5.66d ₀ , and when the parallel part has a hexagonal cross section, L _c = 5.26d ₀ may be set. The length of the parallel portion is set to L _c = 7d ₀ as much as possible. The diameter of the grip portion of the test piece may be the same as the diameter of the parallel portion. In this case, the interval between the grips is set to L _c ≧ 8d ₀ .

<Method for producing Cu-Al-Mn alloy material>
In the Cu-Al-Mn alloy material of the present invention, the production conditions for obtaining a superelastic alloy material having the above-mentioned stable and excellent superelastic properties and excellent in resistance to repeated deformation are as follows. Can be mentioned. An example of a typical manufacturing process is shown in FIG. An example of a preferable manufacturing process is shown in FIG.
In the following description, the processing temperature and processing time (holding time) in each heat treatment indicated as “(for example)”, and the processing rate (cumulative processing rate) in cold processing are the same as those in Example 1, step. No. The values used in a are representatively shown, and the present invention is not limited to these.

  Especially in the whole manufacturing process, the heat treatment temperature [3] in the intermediate annealing [Step 3] is set in the range of 400 to 680 ° C., and cold working (specifically, cold rolling or cold drawing) [Step 4 is performed. Cu-Al-Mn alloy material that stably exhibits good superelastic characteristics by setting the cold rolling ratio or cold drawing ratio [5] in -1] to a range of 30% or more. can get. In addition to this, in the memory heat treatment [Step 5-1] to [Step 5-10], the temperature ranges [8] and [14] (400 to 650 ° C., depending on the alloy composition, which become the (α + β) phase, preferably From 450 ° C. to 550 ° C. to a β single phase [11] and [17] (depending on the alloy composition, usually 700 ° C. or higher, preferably 750 ° C. or higher, more preferably 900 ° C. to 950 ° C.) The heating rates [10] and [16] in [Step 5-3] and [Step 5-7] are both controlled within a predetermined slow range of 0.1 to 20 ° C./min. In addition to this, the cooling rate [13] in the cooling [step 5-5] from the temperature range [11] that becomes the β single phase to the temperature range [14] that becomes the (α + β) phase is 0.1 to 20 Control to a predetermined slow range of ° C / min. Furthermore, after the heating [step 5-3] from the temperature range [8] that becomes the (α + β) phase to the temperature range [11] that becomes the β single phase, in the temperature range [11] that becomes the β single phase. From the holding [step 5-4] for a predetermined time [12] to the subsequent temperature range [11] in which the β single phase is reached to the temperature range [14] in which the (α + β) phase is reached, 0.1 to 20 ° C./min. Cooling at the rate of temperature decrease [13] [step 5-5], passing through the temperature range [14] for a predetermined time [15] [step 5-6], and then the temperature range [14] that becomes the (α + β) phase. To a temperature range [17] from 0 to 0.1 [deg.] C./minute [step 5-7], and further in the temperature range [17] for a predetermined time [18]. The steps [Step 5-4] to [Step 5-8] until the holding [Step 5-8] are repeated at least four times ([Step 5-9]). After this, the final cooling is performed [step 5-10].

  Further, including [step 5-5] and temperature increase [step 5-7], [step 5-4] to [step 5-8] are repeated at least four times before [step 5-9]. Is heated to the temperature range [8] that becomes the (α + β) phase at the rate of temperature increase [7] [Step 5-1], and then held at this temperature range [8] for a certain holding time [9] [Step 5- 2] is preferable. As described above, by once maintaining [step 5-2] in the temperature range [8] that becomes the (α + β) phase [step 5-2] and then raising the temperature to the temperature range [11] that becomes the β single phase [step 5-3], α Since the precipitation amount and size of the phase are kept constant, it is easy to obtain an effect of increasing the crystal grains when the crystal grain coarsening process is performed by rapid cooling [Step 5-10] at the end of the memory heat treatment.

  Therefore, first, the temperature is raised to the temperature range [8] where the α + β phase is obtained [Step 5-1], and then the temperature range [8] where the (α + β) phase is obtained (for example, 450 ° C.) for 2 to 120 minutes [9]. ] [Step 5-2]. When heating in the heat treatment [Step 5-1], it is only necessary to reach the temperature range [8] that becomes the (α + β) phase by raising the temperature, and the rate of temperature rise [7 in this [Step 5-1]. ] Is not particularly limited, and does not need to be gradually increased in the present invention. The rate of temperature increase [7] can be set to, for example, 30 ° C./min, but may be faster or slower. In the holding [Step 5-2], the holding time [9] in the temperature range [8] in which the (α + β) phase is obtained is preferably 10 to 120 minutes. In addition, the precipitation amount of the α phase is fixed in [Step 5-2]. Since the precipitation amount of the α phase can be controlled in [Step 5-2], there is no problem even if the rate of temperature increase in [Step 5-1] is not specified. For this reason, the temperature increase rate of [Step 5-1] can be performed at a high speed, and the entire time required for manufacturing can be shortened. This is one of the merits in the manufacturing method of the present invention.

  Thereafter, the temperature is raised at a rate of temperature rise [10] from a temperature range [8] (for example, 450 ° C.) that becomes the (α + β) phase to a temperature range [11] (for example, 900 ° C.) that becomes the β single phase [Step 5- 3] and hold [12] for a predetermined time [step 5-4] in this temperature range [11]. Thereafter, the temperature is lowered [Step 5-5] at the temperature drop rate [13] to the temperature range [14] that becomes the (α + β) phase, [15] is held for a predetermined time [Step 5-6] in this temperature range [14], The temperature is raised again in the same manner as described above (in the second and subsequent temperature increases [step 5-7], the temperature increase rate [16]). This [Step 5-4] to [Step 5-8] is repeated [19] four or more times in total [Step 5-9]. Then, finally, rapid cooling [Step 5-10] and solution treatment is performed. Such an overall process is preferable.

  Here, the temperature increase rates [10] and [16] and the temperature decrease rate [13] in the memory heat treatment are slowed (in this document, this is also referred to as gradual temperature increase or gradual decrease temperature) and the temperature decrease [Step 5-5]. And by repeating the temperature rise [Step 5-7] four or more times, desired good superelasticity can be obtained even after repeated deformation, and the fracture resistance is further improved. The heating rates [10] and [16] and the cooling rate [13] are all 0.1 to 20 ° C./min, preferably 0.1 to 10 ° C./min, more preferably 0.1. ~ 3.3 ° C / min. As for the memory heat treatment, the last heat treatment (step 5-5) and the stepwise temperature rise [step 5-7] that are repeated at least four times or more (the rightmost in the figure in the illustrated example). After [Step 5-7] [16]), a solution treatment is performed by rapid cooling [Step 5-10] (so-called quenching). This rapid cooling can be performed, for example, by water cooling in which the Cu—Al—Mn alloy material subjected to the storage heat treatment up to the holding heating [step 5-8] in the β single phase is introduced into the cooling water.

Preferably, the following manufacturing processes are mentioned.
After performing melting / casting [Step 1] and hot rolling or hot forging [Step 2] by a conventional method, intermediate annealing [4] at 400 to 680 ° C. [3] for 1 to 120 minutes [4] Step 3] is followed by cold rolling or cold drawing [Step 4-1] with a working rate of 30% or more [5]. Here, the intermediate annealing [Step 3] and the cold working [Step 4-1] may be performed once in this order, and repeated in this order with the number of repetitions [6] of two or more [Step 4]. -2] May be performed. Thereafter, memory heat treatment [Step 5-1] to [Step 5-10] is performed.

The memory heat treatment [Step 5-1] to [Step 5-10] is performed in a temperature range (for example, 450 ° C.) [8] from a temperature range (α + β phase) [8] to a temperature range (for example, 900 ° C.) [β]. 11] up to 0.1-20 ° C./min, preferably 0.1-10 ° C./min, more preferably 0.1-3.3 ° C./min. 3], and the heating temperature [11] is maintained for 5 minutes to 480 minutes, preferably 10 to 360 minutes [12] [Step 5-4], and further becomes a temperature range in which the β single phase is obtained (for example, 900 ° C.) [11] to (α + β phase) temperature range (for example, 450 ° C.) [14] is 0.1 to 20 ° C./min, preferably 0.1 to 10 ° C./min, more preferably 0. Cooling at a temperature drop rate [13] of 1 to 3.3 ° C./min [Step 5-5], and the temperature [14] is set to 20 to 480 minutes, preferably 3 360 min [15] held to [Step 5-6]. Thereafter, heating is performed at a rate of temperature increase [16] of the above-mentioned slow temperature increase from a temperature range (for example, 450 ° C.) [14] that becomes (α + β phase) again to a temperature range (for example, 900 ° C.) [17] that becomes the β single phase [17]. Step 5-7] and hold [18] [Step 5-8] at the temperature [17] for 5 minutes to 480 minutes, preferably 10 to 360 minutes. [Step 5-9] of repeating the slow temperature decrease [13] [Step 5-5] and the gradual temperature increase [16] [Step 5-7] is performed at least four times [19]. Thereafter, rapid cooling [Step 5-10], for example, water cooling steps are included.
The temperature range that becomes α + β single phase and the temperature range defined by the present invention is 400 to 650 ° C, preferably 450 to 550 ° C.
The temperature range for the β single phase is 700 ° C. or higher, preferably 750 ° C. or higher, more preferably 900 to 950 ° C.

  After the memory heat treatment [Step 5-1] to [Step 5-10], it is preferable to perform an aging heat treatment [Step 6] at 100 to 200 ° C. [21] for 5 to 120 minutes [22]. If the aging temperature [21] is too low, the β phase is unstable, and if left at room temperature, the martensitic transformation temperature may change. On the other hand, when the aging temperature [21] is slightly high, bainite (metal structure) occurs, and when it is too high, precipitation of α phase occurs. In particular, the precipitation of α phase tends to significantly reduce shape memory characteristics and superelasticity.

  By repeating the intermediate annealing [Step 3] and the cold working [Step 4-1] [Step 4-2], the crystal orientations can be accumulated more preferably. The number of repetitions [6] of the intermediate annealing [Step 3] and the cold working [Step 4-1] may be one time, but is preferably two times or more, more preferably three times or more. This is because as the number of repetitions [6] of the intermediate annealing [Step 3] and the processing [Step 4-1] increases, the characteristics improve.

(Preferred conditions for each step)
The intermediate annealing [Step 3] is performed at 400 to 680 ° C. [3] for 1 minute to 120 minutes [4]. The intermediate annealing temperature [3] is preferably set to a lower temperature, preferably 400 to 550 ° C.
The cold working [Step 4-1] is a working rate of 30% or more [5]. Here, the processing rate is a value defined by the following equation.
Processing rate (%) = {(A ₁ −A ₂ ) / A ₁ } × 100
A ₁ is the cross-sectional area of the samples before cold working (cold rolling or cold drawing), A ₂ is the cross-sectional area of the sample after cold working.

The cumulative working rate ([6]) when the intermediate annealing [Step 3] and the cold working [Step 4-1] are repeated twice or more is preferably 30% or more, more preferably 45% or more. It is. Although there is no restriction | limiting in particular in the upper limit of a cumulative machining rate, Usually, it is 95% or less.
In the memory heat treatment [Step 5-1] to [Step 5-10], first, in [Step 5-1], the temperature is increased from room temperature to [7] (for example, 30 ° C./min) after the cold working. The temperature is raised to a temperature range (for example, 450 ° C.) [8] that becomes (α + β phase). Then, hold [9] [step 5-2] for 2 to 120 minutes, preferably 10 to 120 minutes [8] in a temperature range (for example, 450 ° C.) that becomes (α + β phase). Then, when heating [step 5-3] from the temperature range (for example, 450 ° C.) [8] that becomes the (α + β phase) to the temperature range (for example, 900 ° C.) [11] that becomes the β single phase, the temperature rises. The temperature rate [10] is set to 0.1 to 20 ° C./min, preferably 0.1 to 10 ° C./min, more preferably 0.1 to 3.3 ° C./min of the above-mentioned slow temperature increase. Then, hold [12] [step 5-4] in this temperature range [11] for 5 to 480 minutes, preferably 10 to 360 minutes. Thereafter, the temperature range from the temperature range (for example, 900 ° C.) [11] to the β single phase to the temperature range (for example, 450 ° C.) [14] to the (α + β phase) 0.1 to 20 ° C./min, preferably 0. 1 to 10 ° C./minute, more preferably 0.1 to 3.3 ° C./minute at a rate of temperature decrease [13] [Step 5-5], and this temperature range [14] for 20 to 480 minutes, preferably Hold [15] for 30 to 360 minutes [Step 5-6]. Thereafter, heating is performed at a temperature increase rate [16] of the gradual temperature increase from a temperature range (for example, 450 ° C.) [14] that becomes (α + β phase) again to a temperature range (for example, 900 ° C.) [17] that becomes a β single phase [17]. Step 5-7] and hold in this temperature range [17] for 5 to 480 minutes, preferably 10 to 360 minutes [18] [Step 5-8]. [Step 5-4] to [Step 5-8] (Conditions [11] to [18]) are repeated [Step 5-9] at least four times [19].
The cooling rate [20] during the rapid cooling [Step 5-10] is usually 30 ° C./second or more, preferably 100 ° C./second or more, more preferably 1000 ° C./second or more.
The last optional aging heat treatment [Step 6] is usually performed at 100 to 200 ° C. [21] for 5 to 120 minutes [22], preferably 120 to 200 ° C. [21] for 5 to 120 minutes [22].

<Physical properties>
The superelastic Cu—Al—Mn alloy material of the present invention has the following physical properties (characteristics).

In the case where the Cu-Al-Mn alloy material of the present invention is subjected to deformation (for example, see FIG. 3 (a)) in which stress load corresponding to 3% strain amount and slow load are repeated, it is 10 ² times or more. Can withstand repeated deformation. More preferably 10 ³ times (although repeated deformation number is. The end of the study in 5000 times) can withstand repeated deformation until rupture.
In addition to the above characteristics, it has been confirmed that the Cu—Al—Mn alloy material of the present invention also achieves the conventionally targeted characteristics (the following residual strain amount). For example, in repeated deformation (for example, see FIG. 4A) in which stress loading and unloading corresponding to a 5% strain amount are repeated 100 times, the remaining strain amount is 2% or less. This amount of residual strain is preferably 1.5% or less. Although there is no restriction | limiting in particular in the lower limit of this residual distortion amount, Usually, it is 0.1% or more.
Further, when the difference between the stress value of 0.2% proof stress and the stress value shown when 5% strain is applied is the difference in stress (for example, see FIG. 4B), the difference is 50 MPa or less. It is preferable. The difference in stress is more preferably 30 MPa or less. Although there is no restriction | limiting in particular in the lower limit of the difference of this stress, Usually, it is 0.1 Mpa or more. This difference in stress indicates the amount of change in the region (plateau region) where the stress shows a substantially constant value with respect to the increase in strain in the stress-strain curve of the shape memory alloy. If this stress difference is reduced within a specified range, even if a large force is applied, only a certain force is transmitted for the strain. For example, when used as a building material, the effect on the building should be reduced. Can do. If the difference in stress is small, the transformation between the parent phase and the martensite phase is easy, so that excellent superelasticity that can withstand repeated deformation and vibration is exhibited.
In addition, the Cu—Al—Mn alloy material of the present invention is different from the usual repeated stress loading and unloading deformation, and the tensile and compressive stress equivalent to 3% strain is 1000 In repetitive deformation (for example, see FIG. 3B), the remaining strain amount is less than 1%. Although there is no restriction | limiting in particular in the lower limit of this residual distortion amount, Usually, it is 0.1% or more.
In this way, unlike deformation that repeats normal stress loading and unloading, the evaluation of repeated tensile and compression results in excellent stability of the properties of this alloy, and tensile and compressive stress equivalent to 3% strain. In the case of deforming 100 cycles, it was confirmed that not only the residual strain but also the yield stress reduction rate (for example, see FIG. 3C) is effective. The rate of decrease is preferably less than 50% compared to the original yield stress, but more preferably less than 30%.

  As described above, the effect of the present invention material becomes even more remarkable by repeating tension and compression. In the present invention, each sample piece is repeatedly deformed by repeating the tension and compression of stress corresponding to 3% strain 1000 times, and the load and slow load of stress corresponding to 3% strain, which is a conventional measurement method, are 1000 times. Repeated evaluation was performed to confirm changes in characteristics. As a result, it was found that the difference was remarkable when the value after 1000 times was compared for the remaining strain amount, and the difference was remarkable when the value after 100 times was compared for the yield stress reduction rate. Therefore, in the present invention, the values after 1000 times and 100 times after repeated deformation are compared for each characteristic.

<Size and shape of superelastic Cu-Al-Mn alloy material>
The Cu—Al—Mn alloy material of the present invention is a shape that is elongated in the processing direction (RD). As described above, the processing direction (RD) is the rolling direction of rolling if the alloy material is a plate material, and the drawing direction of wire drawing if the alloy material is a bar material. Although the alloy material of this invention is extended | stretched with respect to the process direction (RD), the longitudinal direction of an alloy material does not necessarily need to correspond. When cutting or bending the Cu-Al-Mn alloy material of the present invention which is a long body, it is included in the present invention in consideration of the original processing direction of the alloy material. To determine whether or not In addition, there is no restriction | limiting in particular in the specific shape of the Cu-Al-Mn type alloy material of this invention, For example, it can be set as various shapes, such as a stick | rod (wire) and a board | plate (strip). Although there is no restriction | limiting in particular also in these sizes, For example, if it is a rod, it can be set as the diameter of 0.1-50 mm in diameter, or the size of 8-16 mm depending on a use, respectively. Moreover, if it is a board | plate material, the thickness may be 1 mm or more, for example, 1-15 mm. In the manufacturing method of the present invention, a plate material (strip material) can be obtained by performing a rolling process instead of the wire drawing process. Here, it is confirmed that X = 0 can be manufactured with a length of 400 mm or more. However, in the present invention, the abundance of X in the test of d ₀ = 10 (mm) and L _C = 70 (mm) was evaluated based on FIG. 5 JIS No. 14 test piece shape.

  Further, the bar of the present invention is not limited to a round bar (round wire), but may be a square bar (square wire) or a flat bar (flat wire). Here, in order to obtain a square bar (square wire), the round bar (round wire) obtained in advance by the above method is subjected to a conventional method, for example, cold working with a processing machine, cold working with a cassette roller die, press Then, rectangular wire processing such as drawing processing may be performed. Moreover, if the cross-sectional shape obtained in the flat wire processing is appropriately adjusted, a square bar (square wire) having a square cross-sectional shape and a rectangular bar (flat wire) having a rectangular cross-sectional shape can be separately formed. Furthermore, the rod (wire) of the present invention may be in the shape of a hollow tube having a tube wall.

<Vibration control materials and building materials>
The Cu-Al-Mn alloy material of the present invention is suitable for a member for vibration suppression / damping, a member for noise suppression or damping, or a member for self-restoration (self-centering). Can be used. These members are composed of the bar or plate material. Examples of the damping material and the building material are not particularly limited, and examples thereof include a brace, a fastener, and an anchor bolt. Furthermore, it can be used in fields that have been difficult in the past, such as space equipment, aeronautical equipment, automobile parts, building parts, electronic parts, medical products and the like that particularly require repeated deformation resistance. It can also be used as a civil engineering building material that can prevent noise and vibration pollution by utilizing the characteristics of absorbing vibration. Furthermore, in the case of aiming at the effect of noise attenuation, it can be applied in the field of transportation equipment. In any case, since it has an excellent self-restoring force, it can be used as a self-restoring material.
<Seismic control structure>
The Cu—Al—Mn alloy material of the present invention can be suitably used as a damping structure. This damping structure is constructed from the damping material. Examples of the vibration control structure are not particularly limited, and any structure may be used as long as the structure is configured using the braces, fasteners, anchor bolts, and the like.
<Civil engineering materials>
The Cu—Al—Mn alloy material of the present invention can also be used as a civil engineering building material capable of preventing noise and vibration pollution. For example, a composite material can be formed and used with concrete.
<Others>
The Cu—Al—Mn alloy material of the present invention can also be used as a vibration absorbing member and a self-restoring material for space equipment, aircraft and automobiles. It can also be applied to the field of transportation equipment for the purpose of noise attenuation.

  Hereinafter, the present invention will be described in more detail based on examples, but the present invention is not limited thereto.

(Examples 1-43 , Comparative Examples 1-31, 34, 36, 37, 39-41 )
A sample (test material) of a bar (wire) was produced under the following conditions.
Pure copper, pure Mn, pure Al, and, if necessary, raw materials for other sub-addition elements were melted in a high-frequency induction furnace as a material for the Cu—Al—Mn alloy giving the composition shown in Table 1. The melted Cu-Al-Mn alloy was cooled to obtain an ingot having an outer diameter of 80 mm and a length of 300 mm. After the obtained ingot was hot extruded at 800 ° C., in Example 1 of the present invention, the process No. shown in Table 2 was performed. a (a flowchart is shown in FIG. 6A), and in Comparative Example 1, the process No. shown in Table 2 was performed. A bar material of a JIS No. 14 test piece was prepared according to the processing processes shown in A (the flowchart was shown in FIG. 6B). Each of the Examples and Comparative Examples other than these was prepared in the same manner as in Example 1 and Comparative Example 1 except that the processing processes shown in Table 2 were changed.
In addition to Table 2, each step in each processing process shown in Table 3, Table 4-1, Table 4-2, and Table 5 described later is shown in FIG. 2, FIG. 6 (a), and FIG. 6 (b). Corresponding to the parenthesized numbers shown ([Step #]), the alloy composition corresponds to the numbers in Table 1. Further, various production conditions (numbers in parentheses ([#])) other than those shown in Table 2 are as follows, and all examples and comparative examples are not particularly described in Tables 2 and 3. The same conditions were used.

As described above, the melting and casting conditions of [1] were as follows.
The hot working temperature of [2] was 800 ° C.
The intermediate annealing temperature of [3] was 550 ° C.
The intermediate annealing time of [4] was 100 minutes.
The cold working rate of [5] was 30%.
[6] [3] to [5] were repeated three times, and the cumulative cold working rate was 65%.
The rate of temperature increase from room temperature in [7] to the temperature range of the (α + β) phase was 30 ° C./min.
The holding temperature in the temperature range of [8] in the (α + β) phase was 450 ° C.
The holding time in the temperature range of [9] in the (α + β) phase was 60 minutes.
[11] The holding temperature in the temperature range where the β single phase is obtained was set to 900 ° C.
[12] The holding time in the temperature range where the β single phase is obtained was 120 minutes.
[14] The holding temperature in the temperature range of the (α + β) phase was 450 ° C.
The holding time in the temperature range of [15] in the (α + β) phase was 60 minutes.
The temperature increase rate from the temperature range of [16] in the (α + β) phase to the temperature range of the β single phase was the same as in [10].
[17] The holding temperature in the temperature range where the β single phase is obtained was set to 900 ° C.
[18] The holding time in the temperature range where the β single phase is obtained was 120 minutes.
The rapid cooling rate from the temperature range in which [20] is the β single phase was 50 ° C./second.
The aging temperature of [21] was 150 ° C.
The aging time of [22] was 20 minutes.

The evaluation of the superelastic property is performed by repeating stress loading and unloading by a tensile test 5000 times or applying in the stress tension direction and applying 1000 times in the compression direction to obtain a stress-strain curve (SS curve). The residual strain and yield stress reduction rate, and the number of times until breakage were determined and evaluated. In the tensile test, five (N = 5) test pieces were cut out from one specimen and tested. In the following test results, the residual strain is an average value of five.
Table 3, Table 4-1, Table 4-2, and Table 5 show the results of tests and evaluations of Examples and Comparative Examples of the present invention, the types of alloy materials (see Table 1), and processing process conditions (Tables 2 and Tables). 3)).
Hereinafter, each test and evaluation method will be described in detail.

a. Crystal grain size of recrystallized structure Before the tensile test for evaluation of the superelastic repeated deformation resistance described later, the surface of the test piece was etched with an aqueous ferric chloride solution to confirm the crystal grain size. Although the total length of the test piece to be confirmed is not particularly defined, it is considered necessary to have a length equal to or greater than the gauge distance of the tensile test described later. _Therefore, in the present _invention, it was confirmed for the portion of the L C = 70 mm. The schematic diagram of the method for measuring the crystal grain size is as shown in FIG. In the present invention, it is necessary that the crystal grain length a in the direction perpendicular to the working direction (longitudinal direction) is equal to the width or thickness or diameter R of the alloy material, and a = R. In this way, when the number of grain boundaries between crystal grains satisfying a = R is X, the abundance of X is counted and the number of X is measured and evaluated.

b. Cyclic deformation resistance [5% strain load unloading-residual strain after 100 cycles]
A stress-strain curve (SS curve) was obtained by repeatedly performing stress loading and unloading giving a strain of 5% to obtain a residual strain after one cycle to a residual strain after 100 cycles (FIG. 4 ( a)).
Twenty test pieces were cut out from each test material and used for the test. 5% strain load unloading—residual strain after 100 cycles was determined from a stress-strain curve (SS curve). In each table, the residual strain after 100 cycles was shown as “post-cycle residual strain”.

The test condition was a tensile test of 100 times at a test speed of 5% / min, with a test mark distance of 70 mm and alternately repeating stress loading and unloading to obtain a strain amount of 5%. Evaluation was made according to the following criteria.
If the residual strain is 1.5% or less, the superelastic property is excellent, “◎”, and if the residual strain exceeds 1.5% and 2.0% or less, the superelastic property is good. As “◯”, and the case where the residual strain was larger than 2.0% was judged as “X” because the superelastic property was rejected, and is shown in each table.

c. Stress difference between 5% strain and 0.2% strain The stress value giving 5% strain and unloading were performed. From the stress-strain curve (SS curve), the stress value of 0.2% proof stress and 5 The difference of the stress value shown when% strain was loaded was calculated | required as "stress difference" (refer FIG.4 (b)). The “stress difference” is generated when the “stress difference” cannot be properly controlled, for example, when the processing becomes insufficient. For example, when the difference in stress is used as a building material, it is desired that the value of the stress transmitted to the building is small. Therefore, it can be said that the smaller the difference in stress, the better the characteristic. Therefore, when “stress difference” is measured by the above method, “◎” indicates that 30 MPa or less is excellent, “◯” indicates that the pressure exceeds 30 MPa and 50 MPa or less, and “×” indicates that the value exceeding 50 MPa is inferior. Are shown in each table.

d. 3% Strain Tensile Cycle Test 5000 times Stress was applied to give 3% strain and unloading was performed to determine the number of times until fracture (see FIG. 3 (a)). It can be said that the greater the number of times until breakage, the more it can withstand repeated deformation, the more it can suppress the collapse of buildings and the destruction of members. Therefore, when measuring the "repetition endurance to failure" for five those prepared in the above process the same manufacturing conditions, as excellent ones everything was 5000 measurements limit be more than 10 ² times " ◎ ", all 10 ^twice or more, however, as good as a minimum number of more than 900 times" ○ ", those or minimum number less than 10 ² times is unable variation greater control to 10 ¹ to 10 ³ times It was judged as “x” as inferior, and is shown in each table.

e. Residual strain after 1000 cycles of 3% tensile cycle test and tensile compression test From the stress-strain curve (SS curve) by performing a load unloading test of stress giving 3% strain or repeated deformation of tensile load and compression load Residual strain after 1000 cycles was obtained (see FIG. 3B). The above-mentioned “residual strain after 1000 cycles” can be said to be more excellent in self-restoring ability and easier to return to the original shape as the residual strain is smaller. Therefore, when “residual strain after 1000 times” is measured by the above method, “と し て” indicates that the one less than 1.0% is excellent, and “◯” indicates that one that is 1.0% or more and less than 2.0% is good. Those exceeding 2.0% were judged as “x” as inferior, and are shown in each table.
In the case of a tension / compression test, the residual strain is often different between the tension side and the compression side. This is considered to be a case where the initial central axis is deviated. Accordingly, the average value of the values obtained by subtracting the residual strain a on the tension side and the residual strain a ′ on the compression side and dividing by 2 was defined as the residual strain (see FIG. 3B).

f. 3% tensile cycle test and tensile compression test Yield stress reduction rate after 100 cycles Stress-strain curve (SS) by performing unloading test of stress giving 3% strain or repeated deformation of tensile load and compressive load The first 0.2% proof stress value and the 100th 0.2% proof stress value were determined from the curve), and the yield stress reduction rate due to repeated deformation was determined (see FIG. 3C). The above “yield stress reduction rate” is one of the indices indicating the stability of the self-restoring ability characteristic. Since a smaller reduction rate is desired, it can be said that the smaller the decrease rate, the better the characteristics. Therefore, when the “yield stress reduction rate” is measured by the above method, “◎” indicates that 10% or less is excellent, “◎” indicates that 10% or more exceeds 50%, and “◯” indicates that 50% or less is exceeded. It was judged as “x” as inferior, and is shown in each table. The yield stress reduction rate (%) is the difference between the stress value of 0.2% proof stress obtained from the first stress-strain curve and the stress value of 0.2% proof stress of the 100th time. It can be determined by a value obtained by dividing by a stress value of 0.2% proof stress determined from a strain curve × 100 (see FIG. 3C).

g. 5% Tensile Cycle Test Stress value that gave a 0.2% resistance value and 5% strain for the first time from a stress-strain curve (SS curve) after performing an unloading test of stress giving 5% strain This difference was determined as a “stress difference” (see FIG. 4B). The small “stress difference” indicates that the amount of change in the region (plateau region) where the stress shows a substantially constant value with respect to the increase in strain in the stress-strain curve of this alloy is small. That is, it becomes one of the indexes indicating the stability of the self-restoring ability characteristic in the superelastic alloy. Since it is desired that the inclination of the plateau region is small, it can be said that the smaller the plateau region, the better the characteristics. Therefore, when the “stress difference” is measured by the above method, “◎” indicates that 30% or less is excellent, “Good” indicates that 30% or more exceeds 50%, and “◯” indicates that it is inferior if 50% or less. It was judged as “x” and shown in each table.

  It becomes clear from the results shown above, and Examples 1 to 43 are excellent in fracture resistance when subjected to repeated deformation by satisfying the crystal grain size defined in the present invention. Furthermore, by reviewing the usage conditions of this alloy, it has been confirmed that the improvement in characteristics is remarkably improved in the usage method in which tensile load and compression load are repeated rather than the normal load unloading method, and the residual strain after 1000 cycles is confirmed. It was confirmed that it was possible to prevent the yield stress from being deteriorated after 100 cycles.

On the other hand, each comparative example resulted in inferior properties.
Among these, Comparative Examples 1 to 17 shown in Table 3 were inferior because they could not satisfy the crystal grain size defined in the present invention.

In addition, Comparative Examples 18 to 31, 34, 36, 37, and 39 to 41 shown in Table 4-2 do not satisfy the predetermined alloy composition defined by the present invention, or are manufactured according to the present invention. Each of the production itself was not possible because the conditions were not satisfied, or the super-elastic cyclic deformation resistance characteristic of the present invention (residual strain after repeating 100 cycles of stress loading and slow loading corresponding to 5% strain) ).

  In addition, although the description of the test results is omitted, in the case of a Cu—Al—Mn alloy material having a preferable alloy composition of the present invention other than those described in Table 1, a plate material (strip material) instead of a bar material (wire material) ), The same results as in the previous example were obtained.

  While this invention has been described in conjunction with its embodiments, we do not intend to limit our invention in any detail of the description unless otherwise specified and are contrary to the spirit and scope of the invention as set forth in the appended claims. I think it should be interpreted widely.

1 Cu-Al-Mn alloy rod of the present invention (wire)
a Grain length in the direction perpendicular to the machining direction X Grain boundary R Width of alloy material or diameter of rod (wire) RD Processing direction of alloy material (drawing direction of rod (wire))

Claims (1)

3.0-10.0 mass% Al, 5.0-20.0 mass% Mn, Ni, Co, Fe, Ti, V, Cr, Si, Nb, Mo, W, Sn, Mg, P , Be, Sb, Cd, As, Zr, Zn, B, C, Ag, and one or two or more selected from the group consisting of misch metal in a total amount of 0.000 to 10.000% by mass, Here, the content of Ni and Fe is 0.000 to 3.000% by mass, the content of Co is 0.000 to 2.000% by mass, and the content of Ti is 0.000 to 2%. 0.000% by mass, the contents of V, Nb, Mo and Zr are 0.000 to 1.000% by mass, the Cr content is 0.000 to 2.000% by mass, Content is 0.000-2.000 mass%, content of W is 0.000-1 000 mass%, Sn content is 0.000 to 1.000 mass%, Mg content is 0.000 to 0.500 mass%, and P content is 0.000 to 0. .500 mass%, the contents of Be, Sb, Cd, and As are 0.000 to 1.000 mass%, the Zn content is 0.000 to 5.000 mass%, B, The C content is 0.000 to 0.500% by mass, the Ag content is 0.000 to 2.000% by mass, and the misch metal content is 0.000 to 5.000% by mass. A process of melting and casting a Cu-Al-Mn alloy material having a composition consisting of Cu and inevitable impurities in the balance,
A hot working process;
A step of performing intermediate annealing at 400 to 680 ° C. for 1 to 120 minutes and cold working at a processing rate of 30% or more at least once each in this order;
After heating from room temperature to a temperature range of 400 to 650 ° C. that becomes an (α + β) phase, hold in the temperature range for 2 to 120 minutes, and from a temperature range that becomes an (α + β) phase to a temperature range that becomes a β single phase 0.1 to Heat at a rate of temperature increase of 20 ° C./min and hold in the temperature range for 5 to 480 minutes, and then change from a temperature range of β single phase to a temperature range of 400 to 650 ° C. of (α + β) phase. Cool at a rate of temperature decrease of 20 ° C./min and hold in the temperature range for 20 to 480 minutes, then 0.1 to 20 from the temperature range 400 to 650 ° C. in which the (α + β) phase is reached to the temperature range in which the β single phase is reached. Heated at a temperature increase rate of ° C./min and held in the temperature range for 5 to 480 minutes, then rapidly cooled,
Here, from the step of maintaining in the temperature range that becomes the β single phase, from the temperature range that becomes the β single phase to the temperature range 400 to 650 ° C. that becomes the (α + β) phase, 0.1 to 20 ° C./min. The temperature is increased at a rate of 0.1 to 20 ° C./min from the temperature range that becomes the (α + β) phase to the temperature range that becomes the β single phase through the process of cooling at the temperature decrease rate and holding in the temperature range for 20 to 480 minutes. The process is repeated at least 4 times or more until it is heated at 5 to 480 minutes in the temperature range.
The alloy material is an alloy material having a long shape with respect to a processing direction which is a rolling direction or a wire drawing direction,
The crystal grain length a of the alloy material in the direction perpendicular to the processing direction is equal to the width or thickness or diameter R of the alloy material, and X is the number of grain boundaries between crystal grains where a = R. The amount of X present is 1 or less,
Method for manufacturing a Cu-Al-Mn-based alloy material number until the fracture is more than 10 ^twice when performed repeatedly load and unloading of the stress applied 3% strain in the alloy material.

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