Patent Application
20160376688
The present invention relates to a Cu—Al—Mn-based alloy material having excellent resistance to repeated deformations, a method of producing the alloy material, and a rod material or a sheet material using the alloy material.
Shape memory alloys/superelastic alloys, such as copper alloys, exhibit a remarkable shape memory effect and superelastic characteristics concomitantly to reverse transformation of the thermoelastic martensite transformation, and have excellent functions near the living environment temperature. Accordingly, these alloys have been put to practical use in various fields. Representative alloys of the shape memory alloys/superelastic alloys include TiNi alloys and copper (Cu)-based alloys. Copper-based shape memory alloys/superelastic alloys (hereinafter, these are also collectively refer to, simply, copper-based alloys) have characteristics inferior to those of TiNi alloys in terms of repetition characteristics, corrosion resistance, and the like. On the other hand, since the cost is inexpensive, there has been a movement to extend the application range of copper-based alloys. However, although copper-based alloys are advantageous in terms of cost, those alloys are poor in cold workability and inferior in superelastic characteristics. For this reason, despite that a variety of studies are being conducted, it is the current situation that practicalization of copper-based alloys has not been necessarily sufficiently progressed.
Heretofore, various investigations have been conducted on copper-based alloys. For example, Cu—Al—Mn-based shape memory alloys having a β single phase structure with excellent cold workability, have been reported in Patent Literatures 1 to 4 described below. In those examples, for example, regarding a crystalline orientation, the copper-based alloys have a recrystallized texture in which particular orientations, such as <101> and <100>, of a β single phase metallic texture, are aligned in the direction of cold-working, such as rolling or wire-drawing.
Patent Literature 1: JP-A-7-62472 (“JP-A” means unexamined published Japanese patent application)
Patent Literature 2: JP-A-2000-169920
Patent Literature 3: JP-A-2001-20026
Patent Literature 4: WO 2011/152009 A1
A Cu—Al—Mn-based alloy produced by the method of Patent Literature 1 does not have satisfactory characteristics, particularly superelastic characteristics, and the maximum given strain that exhibits shape recovery of 90% or more is about 2 to 3%. Regarding the reason for this, it is speculated that because a strong restraining force is generated among grains at the time of deformation due to reasons, such as the crystalline orientation being random, irreversible defects, such as transition, are introduced. Thus, residual strain that is accumulated due to repeated deformations occurs to a large extent, and after repeated deformations, deterioration of superelastic characteristics also becomes noticeable.
Further, the copper-based alloy of Patent Literature 2 is a copper-based alloy which has shape memory characteristics and superelastic characteristics and which is substantially formed of a β single phase, and the crystal structure is a recrystallized texture in which in the crystalline orientation of the β single phase, particular crystalline orientations, such as <101> and <100>, of the β single phase are aligned in the direction of cold-working, such as rolling or wire-drawing. In the above-described copper-based alloy, the cold-working is performed at a total working ratio after final annealing, at which the frequency of existence of a particular crystalline orientation of the β single phase in the working direction measured by Electron Back-Scatter Diffraction Patterning (hereinafter, may be abbreviated to “EBSP”) (alternatively, also referred to as Electron BackScatter Diffraction (hereinafter, also abbreviated as EBSD)) is 2.0 or higher. Even if the alloy is such a material as described above, since the amount of transformation strain is highly dependent on orientation in Cu—Al—Mn-based alloys, it was insufficient to stably obtain satisfactory superelastic characteristics precisely and uniformly. Further, residual strain that is accumulated due to repeated deformations occurs to a large extent, and after repeated deformations, deterioration of superelastic characteristics also becomes noticeable.
Further, in regard to the copper-based alloys described in Patent Literature 3 and Patent Literature 4, from the viewpoint that the shape memory characteristics and the superelastic characteristics exhibited thereby have large variations in the performance, and these characteristics are not stabilized, there is room for further improvement. Further, it may be considered that in order to stabilize the shape memory characteristics and the superelastic characteristics, texture control is indispensable. However, in the method described in Patent Literature 3, the degree of integration of the texture in the Cu—Al—Mn-based alloy is low, and the shape memory characteristics and the superelastic characteristics are not yet sufficiently stabilized. In Patent Literature 3, it is proposed that the crystalline orientation of the β single phase is controlled in order to enhance the shape memory characteristics and the superelastic characteristics of the copper-based alloy, and also, the average grain size is adjusted to a value equivalent to a half or greater of the wire diameter in the case of a wire material, or to a value equivalent to the sheet thickness or greater in the case of a sheet material, while the area of a region having such a grain size is adjusted to 30% or more of the entire length of the wire material or the entire area of the sheet material. Further, in Patent Literature 4, in order to enhance the shape memory characteristics of the copper-based alloy, and to obtain a copper-based alloy having a cross-section size applicable to structures, it is proposed to produce a macrocrystalline grain structure having a maximum grain size of more than 8 mm. However, in the methods described in Patent Literature 3 and Patent Literature 4, since the control of the grain size distribution of grains having predetermined large grain sizes is more unsatisfactory in a Cu—Al—Mn-based alloy, the shape memory effect or the superelastic characteristics are not stabilized. Further, residual strain that is accumulated due to repeated deformations occurs to a large extent, and after repeated deformations, deterioration of superelastic characteristics also becomes noticeable.
As such, it is considered that integration of the crystalline orientation and having a predetermined large grain size are effective for an enhancement of superelasticity in Cu—Al—Mn-based alloys. However, in the conventional art, no improvement has been made in connection with deterioration of the superelastic characteristics in repeated deformations. However, in a case where these alloys are used for a medical tool, a construction member or the like, deterioration of the characteristics caused by repeated deformations becomes a serious problem, and there is a demand for improvement.
The present invention is implemented for providing a Cu—Al—Mn-based alloy material which has excellent resistance to repeated deformations, for providing a method of producing the same, and for providing a rod material or a sheet material using the alloy material.
The inventors of the present invention conducted a thorough investigation in order to solve the problems described above. As a result, the inventors have found that when the grain size of a Cu—Al—Mn-based alloy material is controlled while the crystalline orientation of the alloy material is controlled, and when the amount of existence (existence proportion) of small grains that do not grow to a predetermined size or larger is controlled, the amount of residual strain after repeated deformations can be reduced. Further, the inventors have found that the control that enables such a balance between the grain size and the texture to be achieved, can be achieved by performing: a shape memory heat treatment, in which a Cu—Al—Mn-based alloy material is subjected to predetermined intermediate annealing and cold-working, then the alloy material is heated in the initial stage of a shape memory heat treatment to a temperature range, in which a state of an (α+β) phase with a fixed amount of a phase precipitation is converted to a β single phase at a particular slow speed of temperature raising, then the alloy material is maintained at a predetermined temperature for a predetermined time, and repeating at least two times of: cooling from a temperature range for forming a β single phase to the temperature range for forming an (α+β) phase at a particular slow speed of temperature lowering; and heating from the temperature range for forming an (α+β) phase to the temperature range for forming a β single phase at a particular slow speed of temperature raising. The present invention was completed based on these findings.
That is, the present invention is to provide the following means:
(1) A Cu—Al—Mn-based alloy material having a composition containing 3.0 to 10.0 mass % of Al, 5.0 to 20.0 mass % of Mn, and 0.000 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.000 to 3.000 mass %; the content of Co is 0.000 to 2.000 mass %; the content of Ti is 0.000 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.000 to 1.000 mass %; the content of Cr is 0.000 to 2.000 mass %; the content of Si is 0.000 to 2.000 mass %; the content of W is 0.000 to 1.000 mass %; the content of Sn is 0.000 to 1.000 mass %; the content of Mg is 0.000 to 0.500 mass %; the content of P is 0.000 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.000 to 1.000 mass %; the content of Zn is 0.000 to 5.000 mass %; the contents of B and C are each 0.000 to 0.500 mass %; the content of Ag is 0.000 to 2.000 mass %; and the content of misch metal is 0.000 to 5.000 mass %; with the balance being Cu and unavoidable impurities,
wherein the alloy material is an alloy material having a shape that is elongated in the working direction, which is the rolling direction or the wire-drawing direction,
wherein in regard to a grain X for which the grain length ax in the working direction of the alloy material is R/2 or less with respect to the width or diameter R of the alloy material, and for which the grain length bx in a direction perpendicular to the working direction is R/4 or less, the amount of existence of the grains X is 15% or less of the total amount of the alloy material, and
wherein in regard to a grain Y′, for which the grain length a in the working direction and the grain length b in the direction perpendicular to the working direction satisfy the relationships of a≧b, and for which the angle formed by the normal line of the (111) plane of that crystal and the working direction is 15° or larger, the amount of existence of the grains Y′ is 85% or more of the total amount of the alloy material.
(2) The Cu—Al—Mn-based alloy material described in the item (1), wherein the Cu—Al—Mn alloy material has the composition containing 0.001 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.001 to 3.000 mass %; the content of Co is 0.001 to 2.000 mass %; the content of Ti is 0.001 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.001 to 1.000 mass %; the content of Cr is 0.001 to 2.000 mass %; the content of Si is 0.001 to 2.000 mass %; the content of W is 0.001 to 1.000 mass %; the content of Sn is 0.001 to 1.000 mass %; the content of Mg is 0.001 to 0.500 mass %; the content of P is 0.010 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.001 to 1.000 mass %; the content of Zn is 0.001 to 5.000 mass %; the contents of B and C are each 0.001 to 0.500 mass %; the content of Ag is 0.001 to 2.000 mass %; and the content of misch metal is 0.001 to 5.000 mass %.
(3) A Cu—Al—Mn-based alloy material having a composition containing 3.0 to 10.0 mass % of Al, 5.0 to 20.0 mass % of Mn, and 0.000 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.000 to 3.000 mass %; the content of Co is 0.000 to 2.000 mass %; the content of Ti is 0.000 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.000 to 1.000 mass %; the content of Cr is 0.000 to 2.000 mass %; the content of Si is 0.000 to 2.000 mass %; the content of W is 0.000 to 1.000 mass %; the content of Sn is 0.000 to 1.000 mass %; the content of Mg is 0.000 to 0.500 mass %; the content of P is 0.000 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.000 to 1.000 mass %; the content of Zn is 0.000 to 5.000 mass %; the contents of B and C are each 0.000 to 0.500 mass %; the content of Ag is 0.000 to 2.000 mass %; and the content of misch metal is 0.000 to 5.000 mass %; with the balance being Cu and unavoidable impurities,
wherein the value of the difference between the stress value of 0.2% proof stress in the case of performing loading and unloading of stress that gives a strain of 5%, and the stress value obtainable when a strain of 5% is loaded, as determined from a stress-strain curve, is 50 MPa or less, and the amount of residual strain obtainable when loading and unloading of the stress that gives a strain of 5% is repeated 100 times, is 2.0% or less.
(4) The Cu—Al—Mn-based alloy material described in the item (3), wherein the Cu—Al—Mn alloy material has the composition containing 0.001 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.001 to 3.000 mass %; the content of Co is 0.001 to 2.000 mass %; the content of Ti is 0.001 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.001 to 1.000 mass %; the content of Cr is 0.001 to 2.000 mass %; the content of Si is 0.001 to 2.000 mass %; the content of W is 0.001 to 1.000 mass %; the content of Sn is 0.001 to 1.000 mass %; the content of Mg is 0.001 to 0.500 mass %; the content of P is 0.010 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.001 to 1.000 mass %; the content of Zn is 0.001 to 5.000 mass %; the contents of B and C are each 0.001 to 0.500 mass %; the content of Ag is 0.001 to 2.000 mass %; and the content of misch metal is 0.001 to 5.000 mass %.
(5) The Cu—Al—Mn-based alloy material described in any one of the items (1) to (4), wherein among the grains Y′, in regard to a grain Z′ in which the angle formed by the normal line of the (101) plane of the crystal and the working direction is 20° or less, the amount of existence of the grains Z′ is 50% or more of the total amount of the alloy material.
(6) A method of producing a Cu—Al—Mn-based alloy material, comprising the steps of:
melting and casting of a raw material of a Cu—Al—Mn-based alloy material having a composition containing 3.0 to 10.0 mass % of Al, 5.0 to 20.0 mass % of Mn, and 0.000 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.000 to 3.000 mass %; the content of Co is 0.000 to 2.000 mass %; the content of Ti is 0.000 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.000 to 1.000 mass %; the content of Cr is 0.000 to 2.000 mass %; the content of Si is 0.000 to 2.000 mass %; the content of W is 0.000 to 1.000 mass %; the content of Sn is 0.000 to 1.000 mass %; the content of Mg is 0.000 to 0.500 mass %; the content of P is 0.000 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.000 to 1.000 mass %; the content of Zn is 0.000 to 5.000 mass %; the contents of B and C are each 0.000 to 0.500 mass %; the content of Ag is 0.000 to 2.000 mass %; and the content of misch metal is 0.000 to 5.000 mass %; with the balance being Cu and unavoidable impurities;
performing hot-working;
performing at least once of intermediate annealing at 400° C. to 680° C. for 1 to 120 minutes and cold-working at a working ratio of 30% or more, in this order; and
heating from room temperature to a temperature range for obtaining an (α+β) phase, then maintaining in this temperature range for 2 to 120 minutes, heating from the temperature range for obtaining the (α+β) phase to a temperature range for obtaining a β single phase at a speed of temperature raising of 0.1° C./min to 20° C./min, maintaining in this temperature range for 5 to 480 minutes, then cooling from the temperature range for obtaining the β single phase to the temperature range for obtaining the (α+β) phase at a speed of temperature lowering of 0.1° C./min to 20° C./min, maintaining in this temperature range for 20 to 480 minutes, then heating from the temperature range for obtaining the (α+β) phase to the temperature range for obtaining the β single phase at a speed of temperature raising of 0.1° C./min to 20° C./min, and maintaining in this temperature range for 5 to 480 minutes, and then rapidly cooling;
wherein the series of steps: from maintaining in the temperature range for obtaining a β single phase, then cooling from the temperature range for obtaining a β single phase to the temperature range for obtaining an (α+β) phase at a speed of temperature lowering of 0.1° C./min to 20° C./min, and maintaining in this temperature range for 20 to 480 minutes; to heating from the temperature range for obtaining an (α+β) phase to the temperature range for obtaining a β single phase at a speed of temperature raising of 0.1° C./min to 20° C./min, and maintaining in this temperature range for 5 to 480 minutes, is repeated at least two times.
(7) The method of producing a Cu—Al—Mn-based alloy material described in the item (6), wherein the Cu—Al—Mn alloy material has the composition containing 0.001 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.001 to 3.000 mass %; the content of Co is 0.001 to 2.000 mass %; the content of Ti is 0.001 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.001 to 1.000 mass %; the content of Cr is 0.001 to 2.000 mass %; the content of Si is 0.001 to 2.000 mass %; the content of W is 0.001 to 1.000 mass %; the content of Sn is 0.001 to 1.000 mass %; the content of Mg is 0.001 to 0.500 mass %; the content of P is 0.010 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.001 to 1.000 mass %; the content of Zn is 0.001 to 5.000 mass %; the contents of B and C are each 0.001 to 0.500 mass %; the content of Ag is 0.001 to 2.000 mass %; and the content of misch metal is 0.001 to 5.000 mass %.
(8) A method of producing a Cu—Al—Mn-based alloy material, which has a composition containing 3.0 to 10.0 mass % of Al, 5.0 to 20.0 mass % of Mn, and 0.000 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.000 to 3.000 mass %; the content of Co is 0.000 to 2.000 mass %; the content of Ti is 0.000 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.000 to 1.000 mass %; the content of Cr is 0.000 to 2.000 mass %; the content of Si is 0.000 to 2.000 mass %; the content of W is 0.000 to 1.000 mass %; the content of Sn is 0.000 to 1.000 mass %; the content of Mg is 0.000 to 0.500 mass %; the content of P is 0.000 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.000 to 1.000 mass %; the content of Zn is 0.000 to 5.000 mass %; the contents of B and C are each 0.000 to 0.500 mass %; the content of Ag is 0.000 to 2.000 mass %; and the content of misch metal is 0.000 to 5.000 mass %; with the balance being Cu and unavoidable impurities;
wherein the alloy material is an alloy material having a shape that is elongated in the working direction, which is the rolling direction or the wire-drawing direction,
wherein in regard to a grain X for which the grain length ax in the working direction of the alloy material is R/2 or less with respect to the width or diameter R of the alloy material, and for which the grain length bx in a direction perpendicular to the working direction is R/4 or less, the amount of existence of the grains X is 15% or less of the total amount of the alloy material, and
wherein in regard to a grain Y, for which the grain length a in the working direction and the grain length b in the direction perpendicular to the working direction satisfy the relationships of a≧b, and for which the angle formed by the normal line of the (111) plane of that crystal and the working direction is 15° or larger, the amount of existence of the grains Y is 85% or more of the total amount of the alloy material.
(9) The method of producing a Cu—Al—Mn-based alloy material described in the item (8), wherein the Cu—Al—Mn alloy has the composition containing 0.001 to 10.000 mass % in total of at least one selected from the group consisting of 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, where the contents of Ni and Fe are each 0.001 to 3.000 mass %; the content of Co is 0.001 to 2.000 mass %; the content of Ti is 0.001 to 2.000 mass %; the contents of V, Nb, Mo, and Zr are each 0.001 to 1.000 mass %; the content of Cr is 0.001 to 2.000 mass %; the content of Si is 0.001 to 2.000 mass %; the content of W is 0.001 to 1.000 mass %; the content of Sn is 0.001 to 1.000 mass %; the content of Mg is 0.001 to 0.500 mass %; the content of P is 0.010 to 0.500 mass %; the contents of Be, Sb, Cd, and As are each 0.001 to 1.000 mass %; the content of Zn is 0.001 to 5.000 mass %; the contents of B and C are each 0.001 to 0.500 mass %; the content of Ag is 0.001 to 2.000 mass %; and the content of misch metal is 0.001 to 5.000 mass %.
(10) The method of producing a Cu—Al—Mn-based alloy material described in any one of the items (6) to (9), wherein after the rapid cooling, an aging heat treatment is carried out at 70° C. to 300° C. for 5 to 120 minutes.
(11) A rod material or a sheet material, which is formed from the Cu—Al—Mn-based alloy material described in any one of the items (1) to (5).
Herein, ‘having excellent resistance to repeated deformations’ means that the amount of residual strain obtainable after loading and unloading at a predetermined amount of strain is repeated at predetermined times, is small, and it is more desirable if this residual strain is smaller. According to the present invention, it means that in regard to repeated deformations, by which loading and unloading of a strain equivalent to an amount of strain of 5% is repeated 100 times, the amount of residual strain is 2.0% or less, and preferably 1.5% or less.
The Cu—Al—Mn-based superelastic alloy material of the present invention can be used in various applications where superelastic characteristics are required, and for example, applications to antennae of mobile telephones, spectacle frames; medical products, such as orthodontic wires, guide wires, stents, ingrown nail correctors (onychocryptosis correctors), and hallux valgus orthoses; as well as connectors and actuators, are expected. Further, the Cu—Al—Mn-based superelastic alloy material of the present invention is preferable as a vibration damping material, such as a bus bar, or as a construction material, due to its excellent resistance to repeated deformations. Further, vibration damping structures and the like can be constructed, using this vibration damping material or construction material. In addition, the alloy material can also be utilized as a civil engineering and construction material enabling prevention of pollutions, such as noises and vibrations, by utilizing the characteristics of absorbing vibrations as described above. The alloy material can also be used as a vibration-absorbing member for aircrafts or automobiles. The alloy material can also be applied in the field of transportation equipment intended for an effect of noise reduction.
Other and further features and advantages of the invention will appear more fully from the following description, appropriately referring to the accompanying drawings.
The Cu—Al—Mn-based alloy material of the present invention is subjected through predetermined intermediate annealing and cold-working, and further via maintaining [Step 5-2] in a temperature range for obtaining an (α+β) phase, which is carried out before the heating [Step 5-3] to a temperature range for obtaining a β single phase that is initially obtained by a shape memory heat treatment, so that the amount of a phase precipitation is fixed thereby. Then, the Cu—Al—Mn-based alloy material is subjected to the shape memory heat treatment, in which cooling [Step 5-5] from the temperature range for obtaining the β single phase to the temperature range for obtaining the (α+β) phase at a particular slow speed of temperature lowering, and heating [Step 5-7] from the temperature range for obtaining the (α+β) phase to the temperature range for obtaining the β single phase at a particular slow speed of temperature raising, are repeated at least two times. Thereby, while the crystalline orientation is controlled in a texture that is oriented to a direction other than the <111> direction, which is a crystalline orientation with high induced stress (that is, the amount of existence of grains in which the angle formed by the normal line of the (111) plane and the working direction (RD) is as small as less than 15° is small), the grain size of grains having a large grain size (the grains Y′ and Z′ in the final state, or the grains Y and Z in the state of the mid course) is controlled to be large in the grain size thereof, and the amount of existence of the grains is controlled to be large. Concomitantly, the amount of existence of small grains that do not grow to a predetermined size or larger (the grain X) can be appropriately controlled to be small. Thus, an alloy material that provides satisfactory superelasticity even if subjected to repeated deformations, is obtained.
The working direction (RD, see
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 (in the present specification, which may be simply referred to as β single phase) at high temperature, and becomes a two-phase texture of a β phase and an a phase (face-centered cubic) (in the present specification, may be simply referred to as (α+β) phase) at low temperature. The temperatures ranges may vary depending on the alloy composition, but 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-based alloy material of the present invention has a composition containing 3.0 to 10.0 mass % of Al and 5.0 to 20.0 mass % of Mn, with the balance being Cu and unavoidable impurities. If the content of elemental Al is too small, the β single phase cannot be formed, and if the content is too large, the alloy material becomes brittle. The content of elemental Al may vary depending onto the content of elemental Mn, but a preferred content of elemental Al is 6.0 to 10.0 mass %. When the alloy material contains elemental Mn, the range of existence of the β phase extends to a lower Al-content side, and cold workability is markedly enhanced. Thus, forming work is made easier. If the amount of addition of elemental Mn is too small, satisfactory workability is not obtained, and the region of a β single phase cannot be formed. Also, if the amount of addition of elemental Mn is too large, sufficient shape recovery characteristics are not obtained. A preferred content of Mn is 8.0 to 12.0 mass %. The Cu—Al—Mn alloy material having the above-described composition has high hot workability and cold workability, and enables to obtain a working ratio of 20 to 90% or higher in cold-working. Thus, the alloy material can be worked by forming into rods (wires) and sheets (strips), as well as fine wires, foils, pipes and the like that have been conventionally difficult to work.
In addition to the essential alloying elements described above, the Cu—Al—Mn-based alloy material of the present invention can further contain, optional additionally alloying element(s), at least one selected from the group consisting of 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 (for example, Pr and Nd). These elements exhibit an effect of enhancing the physical strength of the Cu—Al—Mn-based alloy material, while maintaining cold workability. The content in total of these optional additionally elements is preferably 0.001 to 10.000 mass %, and particularly preferably 0.001 to 5.000 mass %. If the content of these optional additionally elements is too large, the martensite transformation temperature is lowered, and the β single phase texture becomes unstable.
Ni, Co, Fe and Sn are elements that are effective for strengthening of the matrix microstructure. Co makes the grains coarse by forming Co—Al intermetallic compound, but Co in an excess amount causes lowering of toughness of the alloy. A content of Co is 0.001 to 2.000 mass %. A content of Ni and Fe is respectively 0.001 to 3.000 mass %. A content of Sn is 0.001 to 1.000 mass %.
Ti is bonded to N and O, which are inhibitory elements, and forms oxynitride. Also, Ti forms boride when added in combination with B, to enhance physical strength. A content of Ti is 0.001 to 2.000 mass %.
V, Nb, Mo and Zr have an effect of enhancing hardness, to enhance abrasion resistance. Further, since these elements are hardly solid-solubilized into the matrix, the elements precipitate as a β phase (bcc crystals), to enhance physical strength. Contents of V, Nb, Mo and Zr are respectively 0.001 to 1.000 mass %.
Cr is an element effective for retaining abrasion resistance and corrosion resistance. A content of Cr is 0.001 to 2.000 mass %. Si has an effect of enhancing corrosion resistance. A content of Si is 0.001 to 2.000 mass %. W is hardly solid-solubilized into the matrix, and thus has an effect of precipitation strengthening. A content of W is 0.001 to 1.000 mass %.
Mg has an effect of eliminating N and O, which are inhibitory elements, fixes S that is an inhibitory element as sulfide, and has an effect of enhancing hot workability or toughness. Addition of a large amount of Mg brings about grain boundary segregation, and causes embrittlement. A content of Mg is 0.001 to 0.500 mass %.
P acts as a de-acidifying agent, and has an effect of enhancing toughness. A content of P is 0.01 to 0.50 mass %. Be, Sb, Cd, and As have an effect of strengthening the matrix microstructure. Contents of Be, Sb, Cd and As are respectively 0.001 to 1.000 mass %.
Zn has an effect of raising the shape memory treatment temperature. A content of Zn is 0.001 to 5.000 mass %. When appropriate amounts of B and C are used, a pinning effect is obtained, and thereby an effect of coarsening the grains is obtained. Particularly, combined addition of B and C together with Ti and Zr is preferred. Contents of B and C are respectively 0.001 to 0.500 mass %.
Ag has an effect of enhancing cold workability. A content of Ag is 0.001 to 2.000 mass %. When an appropriate amount of misch metal is used, a pinning effect is obtained, and thereby an effect of coarsening the grains is obtained. A content of misch metal is 0.001 to 5.000 mass %. Misch metal refers to an alloy of rare earth elements, such as La, Ce, and Nd, for which separation into simple substances is difficult.
The Cu—Al—Mn-based alloy material of the present invention has a recrystallized texture. Further, the Cu—Al—Mn-based alloy material of the present invention has a recrystallized texture that is substantially formed from (composed of) a β single phase. The expression ‘having a recrystallized texture substantially formed from a β single phase’ means that the proportion occupied by a β phase in the recrystallization texture is generally 90% or more, and preferably 95% or more.
In the technical field of the present invention, even if a large number of grains exist randomly without being aligned a uniform crystalline orientation, if this is a so-called bamboo structure (as schematically shown in
Thus, controlling a Cu—Al—Mn-based alloy material to have a predetermined texture and a predetermined grain size constitutes the technical significance of the present invention. That is, according to the present invention, when a predetermined texture is formed, the alloy material stably exhibits superelastic characteristics, and in addition to that, even if predetermined small grains (grains X) are co-present at a certain low existence ratio in the bamboo structure formed by predetermined large grains (grains Y or Z), exhibition of superelasticity capable of enduring a number (for example, 100 times) of repeated deformations has been made possible. As such, a remarkable effect can be obtained, which is unpredictable from the conventional means.
There also has been a demand for a bamboo structure in the conventional technologies, but only large grains could be controlled, and the control of small grains could not be achieved. Thus, alloy materials exhibited satisfactory superelasticity after several repeated cycles, but the quantity of residual strain increased after a large number of cycles. This is because residual strain is accumulated in the grain boundaries. Small grains that caused residual strain after a large number of repeated deformations were controlled to be eliminated up to a certain mixed use ratio, and thereby the residual strain after a large number of repetitions could be made small. As such, a remarkable effect can be obtained, which is unpredictable from the conventional means.
In the Cu—Al—Mn-based copper alloy of the present invention, grains having small grain sizes (grains X defined in the present invention) exist in an amount of existence (existence proportion) as low as 15% or less, but most of the grains are grains having large grain sizes (for example, grains Y and Z defined in the present invention, in which the grain lengths satisfy the relationships of a≧b). For example, in the case of a rod material, regarding a small grain (this is referred to as grain X) in which the grain length (ax for the grain X) in the working direction (RD) with respect to the sample diameter R is R/2 or less, and the grain length (bx for the grain X) in a direction perpendicular to the working direction (RD) is R/4 or less, the amount of existence of the grains X is 15% or less, and preferably 10% or less, of the total amount of the alloy material. Further, in the case of a sheet material, regarding a small grain (this is referred to as grain X) in which the grain length (ax for the grain X) in the working direction with respect to the sample width R (direction perpendicular to the RD, that is, sample length in the TD) is R/2 or less, and the grain length (bx for the grain X) in a direction perpendicular to the working direction (RD) is R/4 or less, the amount of existence of the grains X is 15% or less, and preferably 10% or less, of the total amount of the alloy material. Herein, the amount of existence of the grains X can be determined based on the proportion of the area (area ratio) occupied by the relevant grains at a surface or a cross-section of the Cu—Al—Mn-based copper alloy material. For the measurement, an area of a surface or a cross-section in the longitudinal direction of the alloy material, in which measurement has been arbitrarily made at 4 or more points, can be employed. In regard to the grain X according to the present invention, evaluation shall be performed at the surface of the Cu—Al—Mn-based alloy material, where the working ratio is substantially higher than the working ratio at the central portion due to the influence of additional shear stress in the working process or the friction at a tool surface, and the grains are likely to become fine.
The large grains, grain Y and grain Z (or grains Y′ and Z′ in the final state), are such that the grain lengths thereof (a and b) satisfy the relationships of a≧b. In regard to the grain Y and the grain Z (or grains Y′ and Z′ in the final state), it is particularly preferable that the grain lengths (a and b, or a′ and b′ in the final state) satisfy the relationships of a≧1.5b (or a′≧1.5b′ in the final state). In the Cu—Al—Mn-based alloy material of the present invention, the superelastic characteristics for repeated deformations can be further enhanced by achieving a balance between the state of the grain sizes and preferably the texture that will be explained below.
Regarding this large grain, with regard to the grain Y (or grain Y′ in the final state) in which the grain length a in the working direction and the grain length b in a direction perpendicular to the working direction satisfy the relationships of a≧b, and the angle formed by the normal line of the (111) plane of the crystal and the working direction (RD) is 15° or larger, the amount of existence of the grain Y (or grain Y′ in the final state) is 85% or more of the total amount of the alloy material. It is preferable that the amount of existence of the grain Y is 90% or more.
Further, among grains Y as described above, with regard to the grain Z (or grain Z′ in the final state) in which the angle formed by the normal line of the (101) plane of the crystal and the working direction (RD) is 20° or less, it is preferable that the amount of existence of the grain Z is 50% or more of the total amount of the alloy material. It is more preferable that the amount of existence of the grain Z (or grain Z′ in the final state) is 60% or more.
In the case where the sum total of the amount of existence of grains X and the amount of grains Y (grains Y include grains Z) is less than 100%, this means that grains having a size other than the sizes of the grains X and the grains Y exist, in addition to the grains X and the grains Y. In this case, the size of the grains having a size other than the sizes of the grains X and the grains Y is larger than that of the grains X and smaller than that of the grains Y.
In regard to the Cu—Al—Mn-based alloy material of the present invention, in the case where the crystalline orientation of a sample is analyzed at a plane that faces the stress axis direction (working direction, RD) by electron backscatter diffraction pattern analysis (EBSP) (taking the area of the alloy material in which measurement has been made arbitrarily at three or more points (magnification of 100×)), 85% or more, and preferably 90% or more, of the grains have a texture in which the angle formed by the normal line of the (111) plane and the working direction is 15° or larger (see
More preferably, the Cu—Al—Mn-based alloy material of the present invention has a texture in which, among the grains Y, in addition to the grain lengths and the texture described above, preferably 50% or more of the grains, and more preferably 60% or more of the grains, are such that the angle formed by the normal line of the (101) plane of the crystal and the working direction (RD) is within the range of 20°. In other words, among the grains Y, the proportion of the grains in which the angle formed by the normal line of the (101) plane of the crystal and the working direction (RD) is 20° or less, is preferably 50% or more, and more preferably 60% or more, of all the grains. In the present invention, such a grain is referred to as grain Z.
In the present invention, the degree of integration in directions other than the <111> direction or the degree of integration in the <101> direction are measured by a SEM-EBSD method. The specific measurement method will be explained below.
The Cu—Al—Mn-based alloy material of the present invention is cut such that the plane facing the stress axis direction (working direction, RD) becomes an observation plane, and the alloy material is embedded in an electroconductive resin and is subjected to vibration-type buff finish (polishing). Measurement is made by an EBSD method at four or more sites in a measurement region of about 800 μm×2,000 μm, under the conditions of a scan step of 5 μm. Herein, regarding the specimen for measuring the recrystallized texture, a specimen extracted at the time point of completion of [Step 5-4] is used. This is because if the Cu—Al—Mn-based alloy material of the present invention is subjected to the entire shape memory heat treatment including the final step [Step 5-10], since grains grow coarsely, it becomes difficult to analyze the texture. Thus, when a specimen is extracted at the time point of completion of [Step 5-4], which is a step of the mid course, the distribution of the crystalline orientation before coarsening of the grains can be checked, and thus, the specimen is inspected in the state described above. For the analysis, the crystalline orientation obtained from all of the analysis results using an OIM software program (trade name, manufactured by TSL) is plotted on an inverse pole figure (see, for example,
With the methods for working and heat treatment according to the present invention, the grain size in the final steps of a shape memory heat treatment can be controlled, without destroying the proportions of the controlled crystalline orientations. Thus, the range of the orientation property of the crystalline orientation according to the present invention is equal to that of the orientation property of the final crystalline orientation.
For example, in Example 1 indicated in Table 3-2, the results obtained by analyzing a specimen extracted at the time point of completing [Step 5-4] at four points in an analytic region having a size of about 800 μm×2,000 μm by the SEM-EBSD method are recorded, as the values of the amounts of existence of the grains Y and the grains Z. Thus, it is shown that the amount of grains Y (proportion of area ratio) in which the angle formed by the normal line of the (111) plane and the working direction was 15° or larger, was 88%, and that, among the grains Y, the amount of grains Z in which the angle formed by the normal line of the (101) plane of the crystal and the working direction was 20° or less, was 60%. That is, in those cases, the magnitude of the grain size is not considered.
On the other hand, regarding the working conditions and the like, for a material that had been produced in the same manner as in Example 1 and subjected up to [Step 5-10], an arbitrary grain was measured by the SEM-EBSD method, the orientation property of the crystalline orientation of the grain was clarified, and then the grain lengths of the grain and the area ratios were determined by calculation. As a result, the amount of grains (hereinafter, grains Y′) in which the angle formed by the normal line of the (111) plane and the working direction was 15° or larger was 89%, and the amount of grains (hereinafter, grains Z′) in which the angle formed by the normal line of the (101) plane and the working direction was 20° or less was 65%. For the grains Y′ and the grains Z′, the crystalline orientations were checked by the SEM-EBSD method, then images of the grain size were taken using a digital camera or the like, and thereby the area (area ratio) is calculated.
When the amounts of existence of the crystalline orientations of the grains at the time points of [Step 5-4] and [Step 5-10] were compared, using the analytic method such as described above. In Example 26, while grains Y 91% and grains Z 60% were obtained at the time point of [Step 5-4] (the state in the mid course of the production), grains Y′ 95% and grains Z′ 68% were obtained at the time point of [Step 5-10] (the final state); in Example 27, while grains Y 88% and grains Z 55% were obtained at the time point of [Step 5-4], grains Y′ 88% and grains Z′ 60% were obtained at the time point of [Step 5-10]; and in Example 39, while grains Y 85% and grains Z 54% were obtained at the time point of [Step 5-4], grains Y′ 85% and grains Z′ 55% were obtained at the time point of [Step 5-10]. Thus, it was confirmed that grains grew almost without any change in the orientation property of the crystalline orientation, and grains were coarsened. This indicates that in the heat treatment steps according to the present invention, generation of new nuclei is not induced by a heat treatment, and grains are coarsened. In addition to the fact that there are limitations on the size of the specimen in the SEM-EBSD method, and that the textures in the mid course can be easily checked, consistency with the final crystalline orientation is confirmed as described above. Thus, the amounts of existence of grains Y and grains Z, which are the textures in the mid course, can be regarded and handled as the amounts of existence of grains Y′ and grains Z′ of the final texture. Accordingly, it can be said that the amount of existence (proportion) of the grains according to the present invention that exhibit a predetermined orientation as checked in the mid course of the production process represents an amount of existence equivalent to that of the final texture state.
In the case where the crystalline orientation of each grain after performing the final heat treatment is measured by the SEM-EBSD method, the measurement region includes grains X, and the area ratio is checked by measuring the crystalline orientations of at least 20 or more at the minimum of grains including grains Y and Z (or grains Y′ and Z′) other than the grains X. In regard to the evaluation of the area ratio in the final state, since grains have been coarsened, the EBSD method is not performed, and the area ratio is calculated from a photograph or the like. That is, in Step [5-4], measurement of the crystalline orientation and the area ratio is performed by the EBSD method, but in [Step 5-10], only the crystalline orientation is measured by the EBSD method, and measurement of the area ratio is performed using a photograph or the like. Herein, for the confirmation of the texture after the final heat treatment of [Step 5-10], measurement of the crystalline orientation and the grain size of the same material at a different position in the longitudinal direction was performed, and similar results were acknowledged.
Further, since the grains X of the material after the final heat treatment had a small grain size, the crystalline orientation was not evaluated, and an evaluation of the grain size and the area ratio only was performed. The measurement range for the area ratio of the grain size related to the grains X is defined as a range including 20 or more at the minimum of grains, similarly to the range in which the grains Y′ and the grains Z′ are identified.
The method of measuring the grain size and the method of measuring the crystalline orientation, each according to the present invention, are performed respectively and independently.
In regard to the Cu—Al—Mn-based alloy material of the present invention, regarding the production conditions for obtaining a superelastic alloy material which stably provides satisfactory superelastic characteristics and has excellent resistance to repeated deformations, a production process such as described below may be mentioned. A representative example of the production process is illustrated in
In the following explanation, the treatment temperature and treatment time (retention time) for a heat treatment, and the working ratio (cumulative working ratio) of cold-working, all being described with the terms “(for example,)” are representatively indicated with the values used in Process No. a in Example 1, and the present invention is not intended to be limited to these values.
In the entire production process, particularly, when the heat treatment temperature [3] for intermediate annealing [Step 3] is set to the range of 400° C. to 680° C., and the cold-working ratio or the working ratio for cold wire-drawing [5] for the cold work (specifically, cold rolling or cold wire-drawing) [Step 4-1] is set to the range of 30% or more, a Cu—Al—Mn-based alloy material which stably provide satisfactory superelastic characteristics is obtained. In addition to those, in the shape memory heat treatment [Step 5-1] to [Step 5-10], the speeds of temperature raising [10] and [16] in heating [Step 5-3] and [Step 5-7] from the temperature ranges [8] and [14] for obtaining the (α+β) phase (which may vary depending on the alloy composition, but usually near 300° C. to 700° C., and preferably 400° C. to 650° C.) to the temperature ranges [11] and [17] for obtaining the β single phase (which may vary depending on the alloy composition, but usually 700° C. or higher, preferably 750° C. or higher, and more preferably 900° C. to 950° C.), and the speed of temperature lowering [13] in cooling [Step 5-5] from the temperature range [11] for obtaining the β single phase to the temperature range [14] for obtaining the (α+β) phase, are all controlled to a predetermined slow range such as 0.1° C./min to 20° C./min. Further, after the heating [Step 5-3] from the temperature range [8] for obtaining the (α+β) phase to the temperature range [11] for obtaining the β single phase, a series of steps including: from retention [Step 5-4] in a temperature range [11] for obtaining the β single phase for a predetermined time [12]; cooling [Step 5-5] from the temperature range [11] for obtaining the β single phase to the temperature range [14] for obtaining the (α+β) phase at a speed of temperature lowering [13] of 0.1° C./min to 20° C./min; retention [Step 5-6] in the temperature range [14] for a predetermined time [15]; heating [Step 5-7] from the temperature range [14] for obtaining the (α+β) phase to the temperature range [17] for obtaining the β single phase at a speed of temperature raising [16] of 0.1° C./min to 20° C./min; to retention [Step 5-8] in the temperature range [17] for a predetermined time [18], that is, a series including from [Step 5-4] to [Step 5-8], is repeated at least two times (Step [5-9]). Thereafter, rapid cooling [Step 5-10] is carried out lastly.
Further, before the [Step 5-9] of repeating at least two times from [Step 5-4] to Step [5-8] including these temperature lowering [Step 5-5] and temperature raising [Step 5-7], it is preferable to perform heating [Step 5-1] to the temperature range [8] for obtaining the (α+β) phase at a speed of temperature raising [7], and then retention [Step 5-2] in this temperature range [8] for a certain retention time [9]. As such, once retention [Step 5-2] in the temperature range [8] for obtaining the (α+β) phase is performed and then temperature raising [Step 5-3] to the temperature range [11] for obtaining the β single phase is performed, the amount of precipitation of the α phase or the size is maintained constant and small. Thus, in the case where a grain coarsening treatment is performed by rapid cooling [Step 5-10] at the end of the shape memory heat treatment, an effect of having enlarged grains may be readily obtained.
Thus, first, when temperature raising [Step 5-1] to the temperature range [8] for obtaining the (α+β) phase is performed, and then the alloy material is subjected to retention [Step 5-2] in this temperature range [8] for obtaining the (α+β) phase (for example, 500° C.) for 2 to 120 minutes [9]. In regard to the alloy material is heated by the heat treatment [Step 5-1] described above, it is desirable if the material reaches the temperature range [8] for obtaining the (α+β) phase by temperature raising. Thus, there are no particular limitations on the speed of temperature raising [7] to this [Step 5-1], and it is not necessary to perform slow temperature raising in the present invention. This speed of temperature raising [7] can be set to, for example, 30° C./min, but the speed of temperature raising may be faster, or on the contrary, may be slower. In regard to the retention [Step 5-2], the retention time [9] in the temperature range [8] for obtaining the (α+β) phase is preferably 10 to 120 minutes. Further, fixing of the amount of precipitation of the α phase is implemented by [Step 5-2]. Since the amount of precipitation of the α phase can be controlled by [Step 5-2], there is no problem even if the speed of temperature raising of [Step 5-1] is not defined. For this reason, the speed of temperature raising of [Step 5-1] can be carried out at a faster speed, and the overall time taken for the production can be shortened. This is one of the advantages for the production method of the present invention.
Thereafter, temperature raising [Step 5-3] from the temperature range [8] for obtaining the (α+β) phase (for example, 500° C.) to the temperature range [11] for obtaining the β single phase (for example, 900° C.) at the speed of temperature raising [10] is performed, and the alloy material is retained [Step 5-4] in this temperature range [11] for a predetermined time [12]. Then, temperature lowering [Step 5-5] to the temperature range [14] for obtaining the (α+β) phase at the speed of temperature lowering [13] is performed, the alloy material is retained [Step 5-6] in this temperature range [14] for a predetermined time [15], and temperature raising (in the temperature raising [Step 5-7] after the second temperature raising, speed of temperature raising [16]) is performed again as described above. Steps including from this [Step 5-4] to [Step 5-8] is repeated [Step 5-9] two or more times [20] in total. Thereafter, rapid cooling [Step 5-10] is performed at the end, and the alloy material is subjected to a solution treatment. It is preferable to perform such an overall process.
Herein, along with slowing of the speeds of temperature raising [10] and [16] and the speed of temperature lowering [13] for the shape memory heat treatment (in the present specification, this is referred to as slow temperature raising and slow temperature lowering, respectively), when the temperature lowering [Step 5-5] and temperature raising [Step 5-7] are repeated two or more times, desired satisfactory superelasticity can be obtained even after repeated deformations. The speeds of temperature raising [10] and [16] and the speed of temperature lowering [13] are all 0.1° C./min to 20° C./min, preferably 0.1° C./min to 10° C./min, and more preferably 0.1° C./min to 3.3° C./min. Further, in regard to the shape memory heat treatment, after the last heating treatment in the slow temperature lowering [Step 5-5] and slow temperature raising [Step 5-7] (in the depicted example, [Step 5-7] and [16] on the rightmost side in the diagram) that are repeated at least two or more times, the alloy material is subjected to a solution treatment by rapid cooling [Step 5-10] (so-called quenching). This rapid cooling can be carried out by, for example, water cooling by introducing a Cu—Al—Mn-based alloy material that has been subjected to a shape memory heat treatment up to retention and heating to the β single phase [Step 5-8], into cooling water.
Preferably, a production process such as follows may be mentioned.
In a usual manner, after melting and casting [Step 1] and hot-working [Step 2] of hot rolling or hot forging is carried out, intermediate annealing [Step 3] at 400° C. to 680° C. [3] for 1 to 120 minutes [4], and then cold-working [Step 4-1] of cold rolling or cold wire-drawing at a working ratio of 30% or higher [5] are carried out. Herein, the intermediate annealing [Step 3] and the cold-working [Step 4-1] may be carried out once each in this order, or may be repeated [Step 4-2] in this order at a number of repetitions [6] of two or more times. Thereafter, the shape memory heat treatment [Step 5-1] to [Step 5-10] is carried out.
The shape memory heat treatment [Step 5-1] to [Step 5-10] includes: heating [Step 5-3] from a temperature range [8] for obtaining an (α+β) phase (for example, 500° C.) to a temperature range [11] for obtaining a β single phase (for example, 900° C.) at a speed of temperature raising [10] of 0.1° C./min to 20° C./min, preferably 0.1° C./min to 10° C./min, and more preferably 0.1° C./min to 3.3° C./min; retention [Step 5-4] at that heating temperature [11] for 5 minutes to 480 minutes, and preferably 10 to 360 minutes [12]; cooling [Step 5-5] from a temperature range [11] for obtaining a β single phase (for example, 900° C.) to a temperature range [14] for obtaining an (α+β) phase (for example, 500° C.) [14] at a speed of temperature lowering [13] of 0.1° C./min to 20° C./min, preferably 0.1° C./min to 10° C./min, and more preferably 0.1° C./min to 3.3° C./min; and retention [Step 5-6] at that temperature [14] for 20 to 480 minutes, and preferably 30 to 360 minutes [15]. Thereafter, the alloy material is subjected to: the heating [Step 5-7] again from a temperature range [14] for obtaining an (α+β) phase (for example, 500° C.) to a temperature range [17] for obtaining a β single phase (for example, 900° C.) at the speed of temperature raising [16] of the slow temperature raising; and retention [Step 5-8] at that temperature [17] for 5 minutes to 480 minutes, and preferably 10 to 360 minutes [18]. Repetition [Step 5-9] of such slow temperature lowering [13] [Step 5-5] and slow temperature raising [16] [Step 5-7] is carried out at a number of repetitions [19] of at least two times. Then, the shape memory heat treatment includes: rapid cooling [Step 5-10], for example, water cooling.
The temperature range for obtaining an (α+β) single phase is set to 300° C. to below 700° C., and preferably 400° C. to 650° C.
The temperature range for obtaining a β single phase is set to 700° C. or higher, preferably 750° C. or higher, and more preferably 900° C. to 950° C.
After the shape memory heat treatment [Step 5-1] to [Step 5-10], it is preferable to perform an aging heat treatment [Step 6] at below 300° C. [21] for 5 to 120 minutes [22]. If the aging temperature [21] is too low, the β phase is unstable, and if the alloy material is left to stand at room temperature, the martensite transformation temperature may change. On the contrary, if the aging temperature [21] is too high, precipitation of the α phase occurs, and the shape memory characteristics or superelasticity tends to be decreased conspicuously.
By repeatedly performing [Step 4-2] intermediate annealing [Step 3] and cold-working [Step 4-1], the crystalline orientation can be integrated more preferably. The number of repetitions [6] of intermediate annealing [Step 3] and cold-working [Step 4-1] may be one time, but is preferably two or more times, and more preferably three or more times. This is because, as the number of repetitions [6] of the intermediate annealing [Step 3] and the cold-working [Step 4-1] is larger, the degree of integration facing the <101> direction increases, to enhance the characteristics.
The intermediate annealing [Step 3] is carried out at 400° C. 680° C. [3] for 1 minute to 120 minutes [4]. It is preferable that this intermediate annealing temperature [3] is set to a lower temperature, and preferably to 400° C. to 550° C.
The cold-working [Step 4-1] is carried out at a working ratio [5] of 30% or higher. Herein, the working ratio is a value defined by formula:
Working ratio (%)={(A1-A2)/A1}×100
wherein A1 represents the cross-sectional area of a specimen obtained before cold-working (cold-rolling or cold-wire-drawing); and A2 represents the cross-sectional area of the specimen obtained after cold-working.
The cumulative working ratio ([6]) in the case of repeatedly performing this intermediate annealing [Step 3] and cold-working [Step 4-1] two or more times is preferably set to 30% or higher, and more preferably 45% or higher. There are no particular limitations on the upper limit of the cumulative working ratio, but the cumulative working ratio is usually 95% or lower.
In regard to the shape memory heat treatment [Step 5-1] to [Step 5-10], first, in [Step 5-1], temperature raising is carried out after the cold-working, from room temperature to a temperature range [8] for obtaining an (α+β) phase (for example, 500° C.) at the speed of temperature raising [7] (for example, 30° C./min). Then, retention [Step 5-2] is performed in a temperature range [8] for obtaining an (α+β) phase (for example, 500° C.) for 2 to 120 minutes, and preferably 10 to 120 minutes [9]. Then, when heating [Step 5-3] is performed from a temperature range [8] for obtaining an (α+β) phase (for example, 500° C.) to a temperature range [11] for obtaining a β single phase (for example, 900° C.), the speed of temperature raising [10] is set to 0.1° C./min to 20° C./min, preferably 0.1° C./min to 10° C./min, and more preferably 0.1° C./min to 3.3° C./min, of the slow temperature raising. Then, the alloy material is retained [Step 5-4] in this temperature range [11] for 5 to 480 minutes, and preferably 10 to 360 minutes [12]. Then, cooling [Step 5-5] is performed from a temperature range [11] for obtaining a β single phase (for example, 900° C.) to a temperature range [14] for obtaining an (α+β) phase (for example, 500° C.) at a speed of temperature lowering [13] of 0.1° C./min to 20° C./min, preferably 0.1° C./min to 10° C./min, and more preferably 0.1° C./min to 3.3° C./min, and the alloy material is retained [Step 5-6] in this temperature range [14] for 20 to 480 minutes, and preferably 30 to 360 minutes [15]. Then, heating [Step 5-7] is performed again from a temperature range [14] for obtaining an (α+β) phase (for example, 500° C.) to a temperature range [17] for obtaining a β single phase (for example, 900° C.) at the speed of temperature raising [16] of the slow temperature raising, and the alloy material is retained [Step 5-8] in this temperature range [17] for 5 to 480 minutes, and preferably 10 to 360 minutes [18]. Repetition [Step 5-9] of such a [Step 5-4] to [Step 5-8] (conditions [11] to [18]) is carried out at least two times [19].
The cooling speed [20] at the time of rapid cooling [Step 5-10] is usually set to 30° C./sec or more, preferably 100° C./sec or more, and more preferably 1,000° C./sec or more.
The final optional aging heat treatment [Step 6] is usually carried out at 70° C. to 300° C. [21] for 5 to 120 minutes [22], and preferably at 80° C. to 250° C. [21] for 5 to 120 minutes [22].
The superelastic Cu—Al—Mn-based alloy material of the present invention has the following physical properties (characteristics).
In regard to the Cu—Al—Mn-based alloy material of the present invention, the amount of residual strain in repeated deformations of repeating 100 times loading and unloading of a stress equivalent to an amount of strain of 5% (see, for example,
Further, in the case where the difference between the stress value of the 0.2% proof stress value and the stress value exhibited when a strain of 5% is loaded is defined as the difference of stress (see, for example,
The Cu—Al—Mn-based alloy material of the present invention is a shaped body that is elongated in the working direction (RD). As described previously, the working direction (RD) is the rolling direction for rolling if the alloy material is a sheet material, and is the wire-drawing direction for wire-drawing if the alloy material is a rod material. The alloy material of the present invention is elongated in the working direction (RD), but it is not necessarily essential that the longitudinal direction of the alloy material is consistent with the working direction. In the case where the Cu—Al—Mn-based alloy material of the present invention, which is a lengthy object, has been cut or bent, whether the alloy material is included in the present invention or not is determined, by considering which direction the original working direction of the alloy material is directed to. There are also no particular limitations on the specific shape of the Cu—Al—Mn-based alloy material of the present invention, and, for example, any shape of rod (wire), sheet (strip), and the like may be taken. There are also no particular limitations on the sizes of the Cu—Al—Mn-based alloy material of the present invention. For example, in the case of the rod, the diameter thereof may be employed 0.1 mm to 50 mm; or alternatively, the diameter of the rod may be the size of 8 mm to 16 mm depending on the use thereof. Further, the sheet may also have the thickness of 1 mm or more, for example, 1 mm to 15 mm. Herein, in regard to the production method of the present invention described above, a sheet material (a strip material) can be obtained by performing rolling instead of wire-drawing.
In the present invention, a rod material (or a wire material) may be any shape of a square rod (or a square wire) or a rectangular rod (or a rectangular wire), in addition to a round rod (or a round wire). In order to obtain the square rod (or the square wire), the round rod (or the round wire) obtained as above is subjected to, in a usual manner, for example, cold-working using a working machine, cold-working using a cassette roller die, pressing, drawing, and the like, to carry out a rectangular wire-working. Further, when a cross-section shape obtainable in rectangular wire-drawing is appropriately adjusted, a square rod (or a square wire) having a square cross-sectional shape and a rectangular rod (or a rectangular wire) having a rectangular cross-sectional shape can be produced individually. Further, the rod material (or the wire material) of the present invention may also have a tubular shape, which is a hollow shape having a tube wall, or the like.
The Cu—Al—Mn-based alloy material of the present invention can be preferably used as a vibration damping material or a construction material. This vibration damping material or construction material is constructed from the rod material or sheet material described above. Examples of the vibration damping material or construction material are not particularly limited, but, for example, may include brace, fastener, anchor bolt, and the like.
The vibration damping structure of the present invention is preferably constructed of the Cu—Al—Mn-based alloy material. This vibration damping structure is constructed of the vibration damping material. Examples of the vibration damping structure are not particularly limited, but any kinds of the structures may be used as long as the structures are constructed of using the above-described brace, fastener, anchor bolt, and the like.
The Cu—Al—Mn-based alloy material of the present invention can also be utilized as a civil engineering and construction material enabling prevention of the pollution of noises or vibrations. For example, the alloy material can be used by forming a composite material together with concrete.
The Cu—Al—Mn-based alloy material of the present invention can also be used as a vibration-absorbing member for an aircraft, an automobile, or the like. The alloy material can also be applied to the field of transportation equipment intended for an effect of attenuating (reducing) noises.
The present invention will be described in more detail based on examples given below, but the invention is not meant to be limited by these.
Samples (specimen) of rods (wires) were produced under the following conditions.
As the raw materials that give the compositions of the Cu—Al—Mn-based alloy as indicated in Table 1-1 and Table 1-2, pure copper, pure Mn, pure Al, and if necessary materials of other optionally adding alloying elements were subjected respectively to melting in a high-frequency induction furnace. The Cu—Al—Mn-based alloys thus melted were cooled, to obtain ingots with diameter of 80 mm×length of 300 mm. The ingot thus obtained was subjected to hot extrusion at 800° C., and then rod materials having a diameter of 10 mm were produced, in Example 1, according to the working process illustrated in Process No. a shown in Table 2 (a flow chart thereof is presented in
The various processes for the steps described in Table 2 as well as Table 3-1, Tables 4-1 and 4-2 described below, correspond to the number indicated in parentheses ([Process #]) shown in
For the melting and casting conditions of [1], as described above, the material was melted in the air and then was cooled and cast in a mold having a predetermined size.
The hot-working temperature of [2] was set to 800° C.
The intermediate annealing temperature of [3] was set to 550° C.
The intermediate annealing time of [4] was set to 100 minutes.
The cold-working ratio of [5] was set to 30%.
The number of repetitions of [3] to [5] in [6] was set to three times, and the cumulative cold-working ratio was set to 65%.
The speed of temperature raising from room temperature to a temperature range for obtaining an (α+β) phase in [7] was set to 30° C./min.
The retention temperature at a temperature range for obtaining an (α+β) phase in [8] was set to 500° C.
The retention time at the temperature range for obtaining an (α+β) phase in [9] was set to 60 minutes.
The retention temperature at a temperature range for obtaining a β single phase in [11] was set to 900° C.
The retention time at the temperature range for obtaining a β single phase in [12] was set to 120 minutes.
The retention temperature at a temperature range for obtaining an (α+β) phase in [14] was set to 500° C.
The retention time at the temperature range for obtaining an (α+β) phase in [15] was set to 60 minutes.
The retention temperature at a temperature range for obtaining a β single phase in [17] was set to 900° C.
The retention time at the temperature range for obtaining a β single phase in [18] was set to 120 minutes.
The rapid-cooling speed from the temperature range for obtaining a 1 single phase in [20] was set to 50° C./sec.
The aging temperature in [21] was set to 150° C.
The aging time in [22] was set to 20 minutes.
Texture observation was performed using an optical microscope or with the naked eye, and the analysis of crystalline orientation was performed using an EBSD method. For the evaluation of superelastic characteristics, loading and unloading of stress by a tensile test was repeated 100 times, a stress-strain curve (S-S curve) was determined, and the residual strain was determined, to evaluate the superelastic characteristics. The tensile test was carried out by cutting five specimens (N=5) from one sample material. In the following test results, the residual strain is the average value of five results.
The results of tests and evaluations of Examples according to the present invention and Comparative Examples are summarized in Tables 3-1 to 3-2 and Tables 4-1 to 4-2, together with the kind of the alloy material (see Tables 1-1 and 1-2) and the working process conditions (see Table 2, Table 3-1, and Tables 4-1 to 4-2).
The methods for tests and evaluations are described in detail below.
a. Recrystallized Texture Orientation
Before the evaluation of resistance to repeated deformations of superelasticity that will be described below, each of specimens was cut such that the plane facing the stress axis direction (working direction, RD) would be an observation plane, followed by embedding in an electrically conductive resin and subjected to vibration-type buffer finish (polishing). Measurement was carried out at four (4) points or more, by an EBSD method, in a measurement region having a size of about 800 μm×2,000 μm, under the conditions of a scan step of 5 μm. Herein, for the sample from which the recrystallized texture was analyzed, a specimen extracted at the time point of completion of [Step 5-4] was used. This is because, if the Cu—Al—Mn-based alloy of the present invention material is subjected up to [Step 5-10], which is the final step of the shape memory heat treatment, since grains have grown coarsely, it becomes difficult to perform a texture analysis. Thus, when a specimen is extracted at a time point of completion of [Step 5-4], which is the step of the mid course, the distribution of the crystalline orientation before coarsening of the grains can be checked. Thus, the samples were checked in the state described above. The crystalline orientations obtained from all of the measurement results, using an OIM software (trade name, manufactured by TSL), were plotted on an inverse pole figure (for example, please see
According to the definitions of the present invention, a grain, which has predetermined grain sizes (a≧b), and in which the angle formed by the normal line of the (111) plane and the working direction (RD) was 15° or larger, is defined as grain Y, and the amount of existence (area ratio) of the grains Y is indicated as “amount of existence (%) of grains Y” in the tables. Further, among the grains Y, a grain, in which the angle formed by the normal line of the (101) plane and the working direction (RD) was 20° or less, is defined as grain Z, and the amount of existence of the grains Z is indicated as “amount of existence (%) of grains Z”.
In regard to the amount of existence (%) of grains Y, an amount of existence of 90% or more was judged excellent and was rated as “A”; an amount of existence of 85% or more and less than 90% was judged satisfactory and was rated as “B”; and an amount of existence of less than 85% was judged unacceptable and was rated as “C”. These grades are indicated in the tables.
Further, in regard to the amount of existence (%) of grains Z, an amount of existence of 60% or more was judged excellent and was rated as “A”; an amount of existence of 50% or more and less than 60% was judged satisfactory and was rated as “B”; and an amount of existence of less than 50% was judged unacceptable and was rated as “C”. These grades are indicated in the tables.
An inverse pole figure produced from the results obtained by measuring the crystalline orientation observed in a plane facing the working direction (RD) of Example 1 by EBSD is presented in
Apart from those, in regard to the samples of Examples and Comparative Examples, the amount of existence of grains Y in which the angle formed by the normal line of the (111) plane and the working direction (RD) is 15° or larger, and the amount of existence of grains Z in which the angle formed by the normal line of the (101) plane and the working direction (RD) is 20° or less, were measured in the same manner by the EBSD method.
b. Gain Size of a Recrystallized Texture
Before a tensile test for an evaluation of the resistance to repeated deformations of superelasticity described below, a specimen in a rod form was etched on the surface with an aqueous solution of ferric chloride, and the grain size was checked. The entire length of the specimen to be checked was not particularly set up, but it was considered that a length equal to or longer than the gauge length of the tensile test that will be described below would be needed. Thus, in the present invention, a length of 100 mm or more was used. The respective samples of Example 1 and Comparative Example 1 were etched with an aqueous solution of ferric chloride, and then texture photographs were taken. The photographs are shown in
According to the definition in the present invention, a grain which satisfies the predetermined relationships of grain sizes (ax and bx) is designated as grain X, and the amount of existence (area ratio) of the grains X is designated as “amount of existence (%) of grains X” in the tables. When the grain sizes of Example 1 and Comparative Example 1 were compared, in Example 1, the amount of grains X was 15% or less, and the relationships of a≧b was satisfied in all of the grains Y (and grains Z). On the other hand, in Comparative Example 1, the grain X existed at an area proportion of more than 15%, and thus the definition in the present invention was not satisfied.
Among the grains of the rod material for which the grain size was measured by the method described above, a sample in which the existence proportion of grains X was 10% or less of the total area of measurement was judged excellent and was rated as “A”; a sample in which the existence proportion of grains X was more than 10% and 15% or less was judged satisfactory and was rated as “B”; and a sample in which the existence proportion was more than 15% was judged poor and was rated as “C”. The ratings are indicated in the tables.
Further, in regard to the grain size in the grains Y (and grains Z), since it is required that the relationships of a≧b be satisfied, the grain size was judged based on the average value of the value of a/b. The value of a/b of a grain Y is indicated as “a/b size of grain Y” in the tables. A sample in which the value of a/b was 1.5 or more was judged excellent and was rated as “A”; a sample in which the value of a/b was less than 1.5 and 1.0 or more was judged satisfactory and was rated as “B”; and a sample in which the value of a/b was less than 1.0 was judged poor and was rated as “C”. The ratings are presented in the tables.
In the case where the sum of the amount of existence of grains X and the amount of existence of grains Y (the grains Y include grains Z) was less than 100%, other grains existed, which had a size other than the sizes of the grains X and the grains Y. In this case, the size of the other grains having a size other than the grains X and the grains Y, was larger than the grains X and smaller than the grains Y.
c. Resistance to Repeated Deformations [Residual Strain after Repeating 100 Cycles—Loading and Unloading of 5% Strain]
Loading and unloading of a stress that resulted in a strain of 5% were repeated, and a stress-strain curve (a S-S curve) was determined. The residual strain after one cycle and the residual strain after 100 cycles were determined (see
Twenty specimens having a length of 170 mm were cut out from each sample and were subjected to the test. The residual strain after 100 cycles of loading and unloading of a strain of 5% was determined from the stress-strain curve (the S-S curve). In the tables, the residual strain after 100 cycles is indicated as “residual strain after cycles”.
Regarding the test conditions, a tensile test of alternately repeating loading and unloading of a stress that gives a strain amount of 5% at a gauge length of 100 mm was carried out 100 times at a test speed of 5%/min. An evaluation was carried out according to the following criteria.
The case where the residual strain was 1.5% or less, was judged to have excellent superelastic characteristics and was rated as “A”; the case where the residual strain was 2.0% or less but more than 1.5%, was judged to have satisfactory superelastic characteristics and was rated as “B”; and the case where the residual strain was large such as more than 2.0%, was judged to have unacceptable superelastic characteristics and was rated as “C”. The results are shown in the tables.
In regard to the representative residual strain, a stress-strain curve (a S-S curve) is presented in
d. Difference of Stress in 5% Strain and 0.2% Strain
Loading and unloading of a stress that gives 5% strain is performed, and the difference between the stress value for 0.2% proof stress and the stress value exhibited when a strain of 5% is loaded is determined, as the “difference of stress” from the stress-strain curve (the S-S curve) (see
As is apparent from the results described above, in Examples 1 to 49, since the grain size and the texture orientation defined in the present invention are satisfied, the resistance to repeated deformations of superelasticity, and the difference of stress between a 5% strain and a 0.2% strain, are excellent. Further, as described above, it was also confirmed that the orientation of the grains (Y and Z) immediately after [Step 5-4] was consistent with the orientation of coarse grains (Y′ and Z′) after the final heat treatment ([Step 5-10]).
Contrary to the above, each of the Comparative Examples resulted in the results in which any of the characteristics was poor.
Among these, in Comparative Examples 1 to 10 shown in Table 3-1 to Table 3-2, and Comparative Examples 32 to 34 shown in Table 4-2, production itself was impossible (Comparative Example 8); at least one condition of the grain size or the texture orientation as defined in the present invention was not satisfied (Comparative Examples other than Comparative Example 8); and the resistance to repeated deformations of superelasticity was poor. In Comparative Examples 9 and 10, the difference of stress was also poor. These are all Comparative Examples with respect to the production method of the present invention. In Comparative Example 8, the intermediate annealing temperature was too low, and disconnection occurred. On the other hand, in Comparative Example 9, the intermediate annealing temperature was too high, and the texture orientation could not be controlled as desired.
Further, since all of Comparative Examples 11 to 31 shown in Table 4-2 did not satisfy the predetermined alloy composition defined in the present invention, the production itself of the materials was impossible (Comparative Examples 11 to 15, 17 to 20, 22, 26, and 30), or although the conditions for the grain size or the texture orientation defined in the present invention were satisfied, the resistance to repeated deformations of superelasticity was poor (Comparative Examples other than Comparative Examples 11 to 15, 17 to 20, 22, 26, and 30).
It can be seen from the results described above that, even if a desired texture could be formed, if the alloy material is not produced under conditions in which the retention at temperature ranges [8] and [14] for obtaining an (α+β) phase for predetermined times [9] and [15] in [Step 5-2] or [Step 5-6], the speeds of temperature raising [10] and [16] in [Step 5-3] and [Step 5-7], the speed of temperature lowering [13] in [Step 5-5], and the number of repetitions [19] of temperature lowering and temperature raising in [Step 5-9] are appropriately satisfied, it is difficult to cause coarsening of grains Y (including grains Z) while the texture is maintained, and to control the amount of existence of grains X to be at a low level. Thus, the grain size or the texture defined in the present invention cannot be satisfied, the difference of stress becomes small (low in characteristics of vibration damping), and the resistance to repeated deformations of superelasticity becomes poor.
Further, the test results were omitted but not shown. However, for the cases of the Cu—Al—Mn-based alloy materials of the present invention, which had the preferred alloy compositions within the ranges defined in the present invention other than those described in Tables 1-1 and 1-2, and for the cases of the sheets (strips) but not the rods (wires), the similar results as those of Examples can be obtained.
Having described our invention as related to the present embodiments, it is our intention that the invention not be limited by any of the details of the description, unless otherwise specified, but rather be construed broadly within its spirit and scope as set out in the accompanying claims.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2014-052462 | Mar 2014 | JP | national |
This application is a Continuation of PCT International Application No. PCT/JP2015/056856 filed on Mar. 9, 2015, which claims priority under 35 U.S.C. 10 §119 (a) to Japanese Patent Application No. 2014-052462 filed in Japan on Mar. 14, 2014. Each of the above applications is hereby expressly incorporated by reference, in its entirety, into the present application.
| Number | Date | Country | |
|---|---|---|---|
| Parent | PCT/JP2015/056856 | Mar 2015 | US |
| Child | 15264113 | US |
