α+β type titanium alloy profile and method for manufacturing the same

Abstract

An α+β type titanium alloy shaped material according to an embodiment of the present invention has a 0.2% yield strength of 830 MPa or higher, an elongation percentage of 10% or higher, a fatigue strength of 450 MPa or higher, and an acicular structure. The titanium alloy shaped material has a void area ratio of 1.0 x 10-5% or less, a torsion angle of ±3.0˚ or less from one end to the other end thereof, and a warp height (mm) / total length (m) of ±2.17 or less.

Description

[Technical Field]

[0001] This invention relates to α+β type titanium alloy profiles and methods for producing the same. [Background technology]

[0002] Titanium alloys, with their high specific strength and excellent corrosion resistance, are used in automotive parts such as connecting rods and mufflers, as well as in consumer goods such as golf club heads and building materials. Furthermore, due to their good biocompatibility, titanium alloys are widely used in jewelry such as watches and eyeglass frames, and in medical applications such as implants.

[0003] Titanium alloys are classified into three types: α-type alloys, β-type alloys, and α+β-type alloys. α-type titanium alloys are those whose microstructure at room temperature is mainly composed of the α phase, β-type titanium alloys are those whose microstructure at room temperature is mainly composed of the β phase, and α+β-type titanium alloys are those whose microstructure at room temperature is composed of both the α and β phases. The α phase of a titanium alloy is a phase with a hexagonal close-packed structure (hcp), and the β phase of a titanium alloy is a phase with a body-centered cubic structure (bcc).

[0004] In pure titanium, the microstructure at room temperature is entirely α-phase, and the microstructure above the β-transformation point temperature is entirely β-phase. However, by adding alloying elements that stabilize the β-phase to pure titanium, the temperature at which the β-phase can stably exist decreases, allowing it to remain even at room temperature. The temperature range above the β-transformation point temperature, where the microstructure is solely β-phase, is called the β-single-phase temperature range. Conversely, the temperature range below the β-transformation point temperature, where the microstructure contains both α-phase and β-phase, is called the α+β two-phase temperature range.

[0005] Among titanium alloys, α+β type titanium alloys have a long track record of use due to their excellent balance of strength and ductility, as well as their fatigue properties. One application of α+β type titanium alloys that has recently attracted attention is their use in reducing weight and improving fuel efficiency in the automotive and motorcycle sectors. α+β type titanium alloys weigh approximately 60% of carbon steel and stainless steel, yet possess almost equivalent strength and fatigue properties. Therefore, by replacing carbon steel or stainless steel crankshafts and engine parts with those made of α+β type titanium alloy, the overall weight of the engine can be reduced, leading to improved power output and reduced fuel consumption.

[0006] Generally, engine parts such as crankshafts, valves, and connecting rods are made from hot-worked materials such as round bars, billets, square bars, and wrought materials supplied in shapes with a uniform cross-section. Hereinafter, these materials will be collectively referred to as "profiles." That is, a "profile" refers to a wrought material that has a uniform cross-section along its entire length and is supplied in a straight shape. Profiles have bends and twists at the stage of manufacture by means of extrusion and rolling, so these are usually removed by means of straightening and annealing. Profiles of α+β type titanium alloy can be applied to the wide range of applications mentioned above by machining, or further machining, or machining after further machining.

[0007] α+β type titanium alloys can achieve an excellent balance of strength and ductility, as well as superior fatigue properties, by performing strong processing in the α+β two-phase temperature range below the β transformation point to control the metal structure to an equiaxed grain structure. Conventionally, complex-shaped profiles have been manufactured by machining forged products or thick plates produced by strong processing in the α+β two-phase temperature range.

[0008] However, since machining increases the manufacturing cost of machine parts, it is preferable to minimize it as much as possible. Currently, in order to reduce manufacturing costs, development is underway on manufacturing technologies that can improve production efficiency by producing long profiles with a cross-sectional shape closer to that of the final product. By using profiles with a cross-section close to that of the final product as the material, the amount of machining can be reduced and the yield can be improved. In addition, productivity can be improved by increasing the length of the profiles.

[0009] For example, in the Eugene Séjournet process, one of the extrusion methods, ingots or round billets obtained by hot forging ingots are used as raw materials. As shown in Figure 1, the raw material (billet 5) is inserted into a container 1, and a hydraulic load is applied to the stem 2 to push the billet 5 in the extrusion direction 11 via a dummy block 3, passing it through a die 4 to form various cross-sectional shapes, thereby making it possible to obtain long profiles 6.

[0010] When processing such complex shapes, the hot deformation resistance of α+β type titanium alloys increases sharply in the temperature range below the β transformation point temperature. Therefore, in order to process α+β type titanium alloys in the temperature range below the β transformation point temperature and control the microstructure to an equiaxed structure, large equipment capable of applying high loads is required. This results in high equipment costs. Furthermore, depending on the required cross-sectional shape, it may be impossible to hot work α+β type titanium alloys in the temperature range below the β transformation point temperature. Moreover, even if processing is possible, if the temperature of a part of the cross-section of the profile exceeds the β transformation point temperature due to processing heat, an equiaxed structure and a needle-like structure obtained by processing above the β transformation point temperature will coexist within the cross-section of the profile, resulting in a significant difference in mechanical properties within the cross-section. For this reason, when hot working α+β type titanium alloys into complex shapes, it has generally been done by processing above the β transformation point temperature, which allows for manufacturing with low hot deformation resistance and minimizes surface defects, and by performing forced cooling as needed, thereby controlling the metal structure to a fine needle-like structure and achieving the necessary excellent balance of strength and ductility and fatigue properties.

[0011] The problem here is that hot-worked profiles generally develop bends and twists. These bends and twists are caused by residual processing stress in the profile and by the fact that the cooling rate of the profile is not uniform across different parts of the profile. Therefore, in the manufacturing of profiles, further straightening processing is required after the hot-working described above.

[0012] Straightening methods include roll straightening and tensile straightening. Roll straightening is a method of improving the bending and twisting of a profile by passing the profile between multiple rolls, causing it to bend and then straighten. Tensile straightening is a method of improving the bending and twisting of a profile by fixing both ends of the profile and applying a tensile load to deform it to the plastic deformation range.

[0013] Because α+β type titanium alloys have high cold deformation resistance, straightening is generally performed at a hot temperature. For profiles of α+β type titanium alloys with a needle-like structure, straightening at a hot temperature allows for straightening with a lower load compared to room temperature. Furthermore, α+β type titanium alloys have a narrower elastic range at high temperatures compared to room temperature, allowing for straightening with only a small amount of deformation. In addition, at high temperatures, the soft β phase fraction of α+β type titanium alloys is higher than at room temperature, making strain recovery easier.

[0014] However, when α+β type titanium alloy profiles are straightened in the low-temperature range of the α+β two-phase temperature range to achieve large deformations, numerous coarse voids, which are internal defects, are formed. This poses a problem in that the excellent balance of strength and ductility, as well as the fatigue properties required for the profiles, cannot be achieved. Furthermore, α+β type titanium alloys have a low Young's modulus. Therefore, when α+β type titanium alloys are hot-stretched, significant springback occurs after cooling to room temperature, and the torsion is often not sufficiently corrected.

[0015] For the reasons stated above, α+β type titanium alloy profiles requiring shape correction have not always been able to achieve the necessary excellent balance of strength and ductility, fatigue properties, and superior shape simultaneously.

[0016] Patent Document 1 discloses a method for straightening α+β type titanium alloys, which involves straightening age-hardened billets, square bars, tubes, etc., by applying tensile stress at a temperature of about 0.3 times the melting temperature (about 480°C), and then cooling while applying tensile stress.

[0017] Patent Document 2 discloses a method for straightening α+β type titanium alloys, which involves annealing a rolled titanium alloy rod for microstructure adjustment, followed by straightening the bend by press straightening and then warm straightening at 600°C to the β transformation temperature. [Prior art documents] [Patent Documents]

[0018] [Patent Document 1] Japanese Patent No. 6058535 [Patent Document 2] Japanese Patent Application Publication No. 2011-137204 [Overview of the project] [Problems that the invention aims to solve]

[0019] In typical hot straightening of α+β type titanium alloys, the material is straightened by 2-3% at a high temperature (around 700-740°C), then held at that high temperature for 10-30 minutes while maintaining the straightened length, before being unloaded and cooled to around 400-600°C before being removed from the straightening machine and allowed to cool. If the profile is round, the cooling rate inside the cross-section is uniform, resulting in a small difference in thermal shrinkage and making warping less likely during cooling after straightening.

[0020] However, when the shape of the profile is square or close to the product shape and complex, for example, due to the presence of corners with a high cooling rate or differences in plate thickness, the temperature difference within the cross-section immediately before heating and cooling increases, and the thermal shrinkage difference also increases. As a result, warping occurs in the profile after being removed from the straightening machine and cooled, or even if no warping occurs after cooling, there is a problem that the residual stress inside the profile is large, and warping occurs when the profile is cut after cooling. Fig. 2 shows an example (a) of a conventional simple-shaped profile of an α+β type titanium alloy profile and an example (b) of a profile with a cross-sectional shape closer to the product shape. Fig. 2 is a cross-sectional view, and the profile extends from the front to the back of the paper surface.

[0021] Also, Figs. 3 to 6B show schematic diagrams related to the shape of the profile that causes problems and the shape changes during cutting. Fig. 3 is an example of the cross-sectional shape of a hot-straightened profile. Fig. 4 is an example of warping in the profile. Fig. 5 is an example of twisting in the profile. Figs. 6A and 6B are examples of dimensional changes after cutting. Note that the warping in Fig. 4 and the twisting in Fig. 5 need to be within the ± limit values in the up-down direction of the paper surface, and the dimensional changes in Figs. 6A and 6B need to be within the ± limit values in the left-right direction of the paper surface.

[0022] In Patent Document 1, after performing α+β two-phase region heat treatment on a titanium alloy, a tensile stress is applied during cooling to balance the stress generated during cooling and straighten titanium and titanium alloy products. However, in this technology, the tensile stress applied to the profile for straightening is at least 20% of the yield stress of the profile at the straightening temperature. When such a large tensile stress is applied to a profile having a acicular structure, stress may concentrate at sites with large twisting or bending, and voids may occur inside. Also, since the titanium alloy described in Patent Document 1 is rolled at the α+β two-phase region temperature, it is presumed that the structure is not an acicular structure but mainly consists of equiaxed grains.

[0023] Patent Document 2 describes a method for manufacturing titanium alloy round bars by annealing them for microstructural adjustment, then straightening them by press straightening, removing surface defects by cutting, performing final straightening by warm straightening at 600°C to the β transformation point temperature, and finally polishing the surface. However, with roll straightening, the transport time from the heating device to the roll straightening machine differs in the longitudinal direction, resulting in differences in cross-sectional dimensions in the longitudinal direction, which reduces productivity and overall yield.

[0024] One possible method to reduce warping is to straighten the material at low temperatures where the temperature difference within the cross-section is small. However, at low temperatures, the elastic limit is large, requiring a very large amount of strain to eliminate torsion. Furthermore, in α+β type titanium alloys with an acicular structure, stress concentration occurs at the β grain boundaries under tensile stress, generating voids internally. This becomes the starting point for fatigue fracture, significantly reducing fatigue properties and degrading the balance between strength and ductility. In other words, many problems remain with conventional methods when manufacturing α+β type titanium alloy profiles in complex shapes that are closer to the actual product shape.

[0025] In view of the above circumstances, the object of the present invention is to provide an α+β type titanium alloy profile that exhibits less warping and twisting, and excellent yield strength, fatigue strength, and ductility, as well as a method for manufacturing the same. [Means for solving the problem]

[0026] The gist of this invention is as follows:

[0027] (1) An α+β type titanium alloy profile according to one aspect of the present invention is The material contains, by mass%, Al: 4.4-6.5%, Fe: 0.5-2.9%, Si: 0-0.50%, O: 0-0.25%, C: 0-0.08%, N: 0-0.05%, Ni: 0-0.15%, Cr: 0-0.25%, and Mn: 0-0.25%, with the remainder being Ti and impurities, and the mass% content of Fe, Ni, Cr, and Mn satisfies the condition 0.5% ≤ %Fe + %Ni + %Cr + %Mn ≤ 2.9%. The material has a 0.2% yield strength of 830 MPa or higher, an elongation of 10.0% or higher, and a fatigue strength of 450 MPa or higher, possesses a needle-like structure, and has a void area ratio of 1.0 × 10⁻⁶. -5 The values ​​are less than or equal to %, the twist angle from one end to the other is within ±3.0°, and the curvature height (mm) / total length (m) is within ±2.17. (2) An α+β type titanium alloy profile according to one aspect of the present invention contains, by mass%, Al: 4.4~5.5%, Fe: 1.4~2.3%, Mo: 1.5~5.5%, O: 0~0.20%, C: 0~0.08%, N: 0~0.05%, Si: 0~0.10%, Ni: 0~0.15%, Cr: 0~0.25%, and Mn: 0~0.25%, with the remainder being Ti and impurities. The mass%, content of Fe, Ni, Cr, and Mn satisfies 1.4% ≤ %Fe+%Ni+%Cr+%Mn ≤ 2.3%, the 0.2% yield strength is 830 MPa or more, the elongation is 10% or more, and the fatigue strength is 450 MPa or more. It has an acicular structure and a void area ratio of 1.0 × 10 -5 The values ​​are less than or equal to %, the twist angle from one end to the other is within ±3.0°, and the curvature height (mm) / total length (m) is within ±2.17. ( 3 ) (1) or (2)In the α+β type titanium alloy profile described above, the 0.2% yield strength may be 850 MPa or higher. ( 4 )the above (1) to (3) any one of the above In the α+β type titanium alloy profile described above, the area ratio of the void is 1.0 × 10 -6 It can be less than a percent. ( 5 ) Above (1)~( 4 In the α+β type titanium alloy profiles described in any one of the items of the above, the average prior β particle size may be 500 μm or less. ( 6 ) Above (1)~( 5 In the case of the α+β type titanium alloy profile described in any one of the items of the above, the maximum residual stress within the cross-section may be +400 MPa or less. ( 7 ) Above (1)~( 6 The α+β type titanium alloy profile described in any one of the items in the ) may be an extruded profile. 。 (8 )A method for producing an α+β type titanium alloy profile according to another aspect of the present invention is as described above (1)~( 7 A method for manufacturing an α+β type titanium alloy profile as described in any one of the items of the Act, comprising the steps of: hot working an α+β type titanium alloy to obtain a profile; heating the profile to a straightening temperature of β transformation point temperature -400°C or higher and β transformation point temperature -200°C or lower, applying a strain of 0.1% or more and 8% or less in the longitudinal direction at the straightening temperature, and further applying a torque such that the longitudinal twist of the profile is within ±3.0%; and cooling the profile to 500°C or lower while applying tensile stress and the torque. ( 9 )the above( 8 In the method for manufacturing the α+β type titanium alloy profile described in ( ), the tensile stress applied to the profile during cooling may be 20% or less of the 0.2% yield strength at room temperature. ( 10 )the above (8) or (9) The method for manufacturing the α+β type titanium alloy profile described herein may further include a step of holding the profile at 500 to 650°C during its cooling. ( 11 )the above (8)~(10)In the method for manufacturing an α+β type titanium alloy profile described in any one of the above, the average cooling rate of the profile from the straightening temperature to 500°C may be 10°C / s or less. [Effects of the Invention]

[0028] According to the present invention, it is possible to provide an α+β type titanium alloy profile that exhibits less warping and twisting, and has excellent yield strength, fatigue strength, and ductility, as well as a method for manufacturing the same. [Brief explanation of the drawing]

[0029] [Figure 1] This is a diagram explaining the Eugene Séjournet procedure. [Figure 2] These are examples of cross-sectional shapes of α+β type titanium alloy profiles, where (a) is an example of a simple profile and (b) is an example of a profile that is closer to the actual product shape. [Figure 3] This is an example of the cross-sectional shape of a hot-formed profile. [Figure 4] This is an example of warping in profiled materials. [Figure 5] This is an example of twisting in profiled material. [Figure 6A] This is an example of dimensional variation after machining in a profile with a T-shaped cross-section. [Figure 6B] This is an example of dimensional variation after machining in a profile with a U-shaped cross-section. [Figure 7A] This is a micrograph of an example of needle-shaped tissue. [Figure 7B] This is a micrograph of an example of isaxial tissue. [Figure 8] This is an example of voids contained in α+β type titanium alloy profiles. [Figure 9A] This is an example of a cross-sectional shape of an α+β type titanium alloy profile. [Figure 9B] This is an example of a cross-sectional shape of an α+β type titanium alloy profile. [Figure 9C] This is an example of a cross-sectional shape of an α+β type titanium alloy profile. [Modes for carrying out the invention]

[0030] The inventors diligently investigated methods for obtaining α+β type titanium alloy profiles that exhibit low warping and torsion, as well as excellent yield strength, fatigue strength, and ductility. As a result, they revealed that the above-mentioned problems can be solved by straightening the α+β type titanium alloy profile under predetermined heating and stress conditions, and then cooling it to a low temperature while applying a tensile load. Previously, if warping or twisting could not be eliminated by straightening processes, the affected parts were discarded, which reduced the manufacturing yield of machine parts. Therefore, eliminating warping and twisting can improve the manufacturing yield of machine parts and reduce costs.

[0031] The following describes in detail an α+β type titanium alloy profile according to one aspect of the present invention.

[0032] The α+β type titanium alloy profile according to this embodiment has an excellent balance of strength and ductility, as well as fatigue strength. Furthermore, it has excellent shape even in the case of complex shapes close to the final product. Specifically, the 0.2% yield strength is 830 MPa or higher, preferably 840 MPa or higher, or 850 MPa or higher; the total elongation is 10.0% or higher, preferably 10.3% or higher, or preferably 12.0% or higher; and the fatigue strength is 450 MPa or higher, preferably 480 MPa or higher. The upper limit of the 0.2% yield strength is not particularly limited, but for example, it may be 1400 MPa or lower, 1300 MPa or lower, or 1200 MPa or lower. The upper limit of the total elongation is also not particularly limited, but for example, it may be 30% or lower, 28% or lower, or 25% or lower. The upper limit of the fatigue strength is also not particularly limited, but for example, it may be 800 MPa or lower, 750 MPa or lower, or 700 MPa or lower. Such mechanical properties can be achieved, for example, by the manufacturing method described later. In the following description, "total growth" may sometimes be simply referred to as "growth."

[0033] The profile according to this embodiment is a so-called α+β type titanium alloy. Below the β transformation point temperature, the α+β type titanium alloy is composed of an α phase having an HCP structure and a β phase having a BCC structure. Above the β transformation point temperature, the α phase transforms into the β phase, resulting in the α+β type titanium alloy consisting only of the β phase.

[0034] Furthermore, the α+β type titanium alloy profile according to this embodiment has a needle-like structure. The needle-like structure is a microstructure that occurs when an α+β type titanium alloy is cooled from a temperature above the β transformation point temperature. As illustrated in Figure 7A, in the needle-like structure, a grain boundary α phase is formed at the prior β grain boundaries, and a structure is formed in which the α phase and β phase are arranged in layers within the prior β grains. Here, the prior β grain boundary is the trace of the grain boundary of the β phase that existed at a temperature above the β transformation point temperature. When an α+β type titanium alloy in the temperature range above the β transformation point temperature is cooled, the α phase preferentially precipitates at the grain boundaries of the prior β phase, and this becomes the grain boundary α phase.

[0035] The α+β type titanium alloy profile according to this embodiment may consist only of a needle-like structure, or it may have both a needle-like structure and an equiaxed structure. An equiaxed structure is a structure that does not have the characteristics of the needle-like structure described above, and is composed of α grains and a transformed β phase, as illustrated in Figure 7B, for example. Here, the transformed β phase is a layered structure of β phase and α phase, which is a structure that is in the β phase during hot working and undergoes a phase transformation between the β phase and the α phase during cooling after hot working.

[0036] α+β type titanium alloys with a needle-like structure are prone to void formation at prior β grain boundaries under tensile stress. Generally, the vicinity of voids within a metal has a high stress concentration factor under tensile stress, making it a likely starting point for fracture during tensile and fatigue tests. This is thought to reduce elongation, as well as the 0.2% proof stress and fatigue strength.

[0037] Conventional straightening techniques involve performing tensile or roll straightening in a high-temperature range within the narrow elastic α+β two-phase temperature range to prevent void formation, followed by air cooling, or fixing the material with a mold to suppress deformation due to thermal stress during cooling. This allows for the acquisition of both material properties and dimensional accuracy. However, if the profile has ends with a fast cooling rate, or if the profile has a complex shape, differences in cooling rate or plate thickness can lead to large temperature differences within the cross-section during heating, resulting in large differences in thermal shrinkage during cooling. This causes significant stress to be generated inside the profile during cooling after straightening, causing plastic deformation such as warping and twisting. Furthermore, even if plastic deformation does not occur during cooling after straightening, or if plastic deformation is suppressed by a mold, residual stress is generated inside the profile, causing shape defects such as warping during machining.

[0038] In the α+β type titanium alloy profile according to this embodiment, the average particle size of the region surrounded by prior β grain boundaries, i.e., the average prior β particle size, is not particularly limited, but is preferably 500 μm or less. This further enhances the mechanical properties of the profile. However, when the profile is manufactured by hot rolling and then straightened, the average prior β particle size is usually greater than 500 μm. Hot rolling is usually performed in multiple passes using a hot rolling mill, and β grains grow as the titanium alloy moves between these passes. Also, since the amount of strain introduced per pass during hot rolling is small, the strain recovers as the titanium alloy moves between these passes, making recrystallization less likely. For these reasons, the average prior β particle size of the titanium alloy profile obtained by hot rolling becomes coarser. On the other hand, for example, when the profile is formed by hot extrusion and then straightened, the average prior β particle size is 500 μm or less. The average prior β particle size may be 450 μm or less, 400 μm or less, or 350 μm or less. The lower limit of the average prior β particle size is not particularly limited, but for example, the average prior β particle size may be 50 μm or more, or 80 μm or more. Since hot rolling does not allow for complex cross-sectional shapes of the profiles, it is preferable that the α+β type titanium alloy profiles according to this embodiment be extruded profiles manufactured by hot extrusion.

[0039] The method for measuring the average old β particle size is as follows. First, a test piece for microstructure observation is taken from inside the shaped material. Ensure that this test piece does not include the region from the surface of the shaped material to a depth of 0.5 mm. Then, wet-polish and buff-polish this test piece to make the polished surface mirror-like. Corrode this polished surface using hydrofluoric nitric acid to reveal the metal microstructure. When observing the corroded surface with an optical microscope, the grain boundary α-phase as shown in Fig. 7A can be identified. This grain boundary α-phase can be regarded as the old β grain boundary. By measuring the equivalent circle diameter of the portion surrounded by the old β grain boundary using the cutting method, the average value of the old β particle size can be obtained. When measuring the particle size by the cutting method, the size of the measurement region is 7 mm square, the measurement magnification is 50 - 200 times, and 15 straight lines are described at equal intervals vertically and horizontally in the measurement region.

[0040] The α + β type titanium alloy shaped material according to this embodiment also has a void area ratio of 1.0×10 -5 % or less and a microstructure of acicular structure. As a result, a high strength-ductility balance and fatigue characteristics can be obtained. The microstructure being an acicular structure means that 50% or more of the microstructure by area ratio is an acicular structure, preferably 60% or more, more preferably 70% or more, and most preferably 100% is an acicular structure. The microstructure other than the acicular structure is an equiaxed structure. Also, the void area ratio is preferably 5.0×10 -6 % or less, 1.0×10 -6 % or less, or 1.0×10 -7 % or less. As described above, voids become the starting points of fatigue fracture and significantly reduce fatigue characteristics. Furthermore, voids also have an adverse effect on strength and ductility. Therefore, the smaller the void area ratio, the better. The lower limit value of the void area ratio is not particularly limited and may be, for example, 0%. The void area ratio may be defined as 1.0×10 -9 % or more, 5.0×10 -9 % or more, or 1.0×10 -8 % or more as well.

[0041] The following describes the needle-like and equiaxed structures in detail. A needle-like structure is a structure composed of layers of α and β phases, surrounded by thin grain boundary α phases, as shown in Figure 7A, for example. An equiaxed structure is a structure composed of granular proecution α phases or transformed β phases not surrounded by grain boundary α phases, as shown in Figure 7B, for example. When an α+β type titanium alloy is heated to the β temperature range and then allowed to cool naturally without machining, the structure mainly becomes needle-like. On the other hand, when an α+β type titanium alloy is hot-worked in the α+β two-phase temperature range and then cooled, an equiaxed structure is formed. The amount of needle-like and equiaxed structures varies depending on the temperature of the hot-working and the amount of strain applied. The method for measuring the amount of needle-like structure is as follows: First, a test piece for microstructure observation is taken from inside the profile. This test piece does not include the area from the surface of the profile to a depth of 0.5 mm. Then, this test piece is polished, and the polished surface is corroded with hydrofluoric acid to reveal the metallic structure. When the corroded surface is observed with an optical microscope, the grain boundary α phase as shown in Figure 7A can be identified. The area ratio of the region enclosed by grain boundary α and the grain boundary α phase included in the measurement field of view is considered to be the area ratio of the needle-like structure of the profile, relative to the total area of ​​the measurement field of view.

[0042] Voids are identified by mirror-polishing samples taken from the cross-section and longitudinal section of the profile and then observing them using an optical microscope. The sample should not include any area from the surface of the profile down to a depth of 0.5 mm. Voids are defined as the sum of the areas of voids with an equivalent circular diameter of 1 μm or more. When observed with an optical microscope, voids can be seen as circular, dark-colored regions, as shown in Figure 8, for example.

[0043] The residual stress is measured using the 2D method with X-ray diffraction (Bruker AXS D8 Discover). The X-ray tube is made of Cu, the collimator diameter is 2.0 mm, and the diffraction line used is the (302) line of the Tiα phase. The cross-section at the midpoint of the profile is mirror-polished, and the residual stress distribution in the direction of the cross-section is measured to determine the residual stress value that has the maximum absolute value. The cross-section is the section perpendicular to the length direction of the profile. A specific example of measuring the maximum value of residual stress is described below. For example, in the case of a T-shaped profile shown in Figure 6A, which consists of a horizontal plate and a vertical plate abutting the center of the horizontal plate, the residual stress is maximum at one point along the central axis of the vertical plate. Therefore, by measuring the residual stress along the central axis of the vertical plate, the maximum residual stress within the cross-section of the profile can be determined.

[0044] In this embodiment, it is preferable that the α+β type titanium alloy profile, after hot straightening, exhibits minimal dimensional change when subjected to machining. Regarding dimensional change during machining, it was determined that a change within ±2.0 mm within the cross-section would not pose a problem when processing into parts or products. Preferably, the dimensional change during machining is within ±1.5 mm, and even more preferably within ±1.0 mm; naturally, smaller dimensional changes are preferable. This dimensional change is caused by residual stress within the profile. For the dimensional change within the cross-section during machining to be within ±2.0 mm, it is preferable that the maximum residual stress within the cross-section be +400 MPa or less, and more preferably +300 MPa or less. It is estimated that the dimensional change during machining will also decrease as the absolute value of the residual stress, which is the main cause of dimensional change during machining, decreases.

[0045] The chemical composition of the titanium alloy according to this embodiment is not particularly limited as long as it has the components of an α+β type titanium alloy as illustrated above, but for example, it can have the components shown in [1] and [2] below. Hereinafter, "%" represents "mass%".

[0046] [1] Al: 4.4-6.5%, Fe: 0.5-2.9%, Si: 0-0.50%, O: 0-0.25%, C: 0-0.080%, N: 0-0.050%, Ni: 0-0.15%, Cr: 0-0.25%, Mn: 0-0.25%, remainder: Ti and impurities totaling 0.4% or less, where the mass percentages of Fe, Ni, Cr, and Mn satisfy the condition 1.4% ≤ %Fe + %Ni + %Cr + %Mn ≤ 2.9%.

[0047] Al: 4.4~6.5% Al is an α-stabilizing element and may be included to increase the fraction of the α phase. Considering the balance of 0.2% yield strength, ductility, and toughness, the Al content is preferably 4.4 to 6.5%.

[0048] Fe: 0.5~2.9% Fe is a β-stabilizing element, and its inclusion lowers the β-transformation temperature. Furthermore, Fe improves the 0.2% proof stress. Considering the balance between 0.2% proof stress, segregation during solidification, and elongation, the Fe content is preferably 0.5-2.9%.

[0049] Si:0~0.50%, O:0~0.25%, C:0~0.080%, N:0~0.050% Si, O, C, and N are not essential elements, and the lower limit of their content is 0. These elements also have similar effects as α-stabilizing elements, increasing the fraction of the α phase and improving the 0.2% yield strength. Considering the balance with ductility, the preferred content is Si: 0~0.50%, O: 0~0.25%, C: 0~0.080%, and N: 0~0.050%. However, if the content of these elements is 0%, the refining cost will increase. Therefore, the lower limits for O, C, and N may be set to 0.010% or 0.050% respectively. For Si, the lower limit may be the lower limit of the measurement accuracy of chemical analysis, which is 0.001%.

[0050] Ni:0~0.15%, Cr:0~0.25%, Mn:0~0.25% Ni, Cr, and Mn are not essential elements, and the lower limit of their content is 0. These elements function similarly to Fe, so their inclusion is acceptable. When the content of Ni, Cr, and Mn increases, intermetallic compounds (Ti2Ni, TiCr2, TiMn), which are equilibrium phases, are formed, degrading fatigue strength and room-temperature ductility. Therefore, it is preferable that the content be Ni: 0-0.15%, Cr: 0-0.25%, and Mn: 0-0.25%. The lower limits for Ni, Cr, and Mn may be the lower limits of the measurement accuracy of chemical analysis, which are 0.001%, 0.001%, or 0.001%, respectively.

[0051] The total amount of Ni, Cr, Mn, and Fe is preferably between 0.50% and 2.90%, taking into consideration the balance between room temperature tensile strength and room temperature ductility.

[0052] Remainder: Ti and impurities totaling 0.4% or less The remainder, if the chemical composition is [1], consists of Ti and impurities. Examples of impurities include Cl, Na, Mg, which are introduced during the titanium refining process, and Zr, Sn, Cu, Mo, Nb, and Ta, which are introduced from scrap. As the content of any of these impurities increases, they form compounds with Ti, reducing toughness and consequently decreasing workability. Furthermore, if the total content of impurities becomes excessive, ductility decreases, leading to a deterioration in workability. Therefore, it is preferable to control the total amount of impurity elements to 0.4% or less so as not to hinder the effect of the α+β type titanium alloy profile according to this embodiment. It is also preferable that the content of each individual impurity element be 0.1% or less.

[0053] [2] Al: 4.4-5.5%, Fe: 1.4-2.3%, Mo: 1.5-5.5%, O: 0-0.20%, C: 0-0.080%, N: 0-0.05%, Si: 0-0.10%, Ni: 0-0.15%, Cr: 0-0.25%, Mn: 0-0.25%, remainder: Ti and impurities totaling 0.4% or less, where the mass percentages of Fe, Ni, Cr, and Mn satisfy the condition %Fe, %Ni, %Cr, and %Mn = 1.4% ≤ %Fe + %Ni + %Cr + %Mn ≤ 2.3%.

[0054] Al: 4.4-5.5% Al is an α-stabilizing element and is included to increase the fraction of the α phase. Considering the balance of 0.2% yield strength, ductility, and toughness, the Al content is preferably 4.4 to 5.5%.

[0055] Fe: 1.4~2.3% Fe is a β-stabilizing element, and its inclusion lowers the β-transformation temperature. Furthermore, Fe improves the 0.2% proof stress. Considering the balance between 0.2% proof stress, segregation during solidification, and elongation, the Fe content is preferably 1.4-2.3%.

[0056] Mo: 1.5~5.5% Mo is a β-stabilizing element and, like Fe, can lower the β-transformation temperature of titanium alloys. Furthermore, Mo improves the 0.2% yield strength, ductility, and fatigue strength, as well as hot workability. Considering the balance with solidification segregation, the Mo content is preferably between 1.5% and 5.5%.

[0057] O:0~0.20%, C:0~0.080%, N:0~0.050%, Si:0~0.10% The presence of O, C, N, and Si is not mandatory, and the lower limit of their content is 0. These elements also have similar effects as α-stabilizing elements, and their inclusion increases the fraction of the α phase and improves the 0.2% yield strength. Considering the balance with ductility, the preferred content is O: 0~0.20%, C: 0~0.080%, N: 0~0.050%, and Si: 0~0.10%. However, if the content of O, C, and N is 0%, the refining cost will increase. Therefore, the lower limits for O, C, and N may be set to 0.010% or 0.050% respectively. For Si, the lower limit may be the lower limit of the measurement accuracy of chemical analysis, which is 0.001%.

[0058] Ni:0~0.15%, Cr:0~0.25%, Mn:0~0.25% Ni, Cr, and Mn are not essential elements, and the lower limit of their content is 0. These elements function similarly to Fe, so their inclusion is acceptable. When the content of Ni, Cr, and Mn increases, intermetallic compounds (Ti2Ni, TiCr2, TiMn), which are equilibrium phases, are formed, degrading fatigue strength and room-temperature ductility. Therefore, it is preferable that the content be Ni: 0-0.15%, Cr: 0-0.25%, and Mn: 0-0.25%. The lower limits for Ni, Cr, and Mn may be the lower limits of the measurement accuracy of chemical analysis, which are 0.001%, 0.001%, or 0.001%, respectively.

[0059] The total amount of Ni, Cr, Mn, and Fe is preferably between 1.40% and 2.30%, taking into consideration the balance between room temperature tensile strength and room temperature ductility.

[0060] Remainder: Ti and impurities totaling 0.4% or less The remainder, even when the chemical composition is [2], consists of Ti and impurities. Examples of impurity elements include Cl, Na, Mg, which are introduced during the titanium refining process, and Zr, Sn, Cu, Nb, and Ta, which are introduced from scrap. As the content of any of these impurities increases, they form compounds with Ti, reducing toughness and consequently decreasing workability. Furthermore, if the total content of impurities becomes excessive, ductility decreases, leading to a deterioration in workability. Therefore, it is preferable to control the total amount of impurity elements to 0.4% or less so as not to hinder the effect of the α+β type titanium alloy profile according to this embodiment. It is also preferable that the content of each individual impurity element be 0.1% or less.

[0061] The cross-sectional shape of the α+β type titanium alloy profile according to this embodiment is not particularly limited, as long as it is a shape that allows for the measurement of the twist angle from one end to the other of the profile. In the case of a round bar, the twist angle cannot be measured, so a circular shape is naturally excluded from the cross-sectional shape, but any other cross-sectional shape may be applied to the α+β type titanium alloy profile according to this embodiment. Examples of cross-sectional shapes include L-shaped, T-shaped, H-shaped, U-shaped, π-shaped, + (plus sign) shaped, and - (minus sign) shaped. It may also have shapes as shown in Figures 9A to 9C.

[0062] Figure 9A is a cross-sectional view of a profile having a shape similar to a so-called H-shaped profile. An H-shaped profile has a shape in which first and second parallel-extending plates and a third plate whose end face abuts against the surfaces of the first and second plates are integrally formed. On the other hand, the profile in Figure 9A has a shape in which first, second, and third parallel-extending plates, a third plate whose end face abuts against the surfaces of the first and second plates, and a fourth plate whose end face abuts against the surfaces of the second and third plates are integrally formed. In other words, the profile in Figure 9A has a shape like two H-shaped profiles placed side by side. A shape like three or more H-shaped profiles placed side by side is also acceptable.

[0063] Figure 9B is a cross-sectional view of a profile having a shape similar to a so-called T-shaped profile. A T-shaped profile has a shape in which a first flat plate and a second flat plate with one end face abutting against one side of the first flat plate are integrated together. On the other hand, the profile in Figure 9B has a shape in which these first and second flat plates and a third flat plate with one end face abutting against one side of the second flat plate are integrated together.

[0064] Figure 9C is also a cross-sectional view of a profile having a shape similar to a so-called T-shaped profile. The profile in Figure 9C has a shape in which the first and second flat plates, a third flat plate with its end face abutting against one surface of the second flat plate, and a fourth flat plate with its end face abutting against the other surface of the second flat plate are integrated together.

[0065] The shape described above is merely one example of the cross-sectional shape of the α+β type titanium alloy profile according to this embodiment. For example, various variations can be adopted for the cross-sectional shape of the α+β type titanium alloy profile according to this embodiment, such as a cross-sectional shape that includes a knob-like protrusion as part of the cross-sectional shape exemplified above.

[0066] Next, we will explain the manufacturing method for α+β type titanium alloy profiles that exhibit the aforementioned excellent shape. Hereafter, "profile temperature" refers to the surface temperature of the profile as measured by a radiation thermometer.

[0067] First, an α+β type titanium alloy is hot-worked using manufacturing methods such as extrusion, die forging, or rolling to obtain a profile of the desired shape. The manufacturing method of the profile is not particularly limited, but hot extrusion is preferred considering manufacturing efficiency, etc. However, warping and twisting inevitably occur in profiled materials. With conventional straightening processes, it is difficult to correct warping and twisting while ensuring the mechanical properties of the profiled material. In particular, if the profiled material is manufactured by hot extrusion, the magnitude of warping and twisting becomes significant. In the manufacturing method of α+β type titanium alloy profiled material according to this embodiment, this problem is solved by the straightening process described later.

[0068] Next, the profile is heated and held in the α+β two-phase region. Then, at a straightening temperature between -400°C (β transformation point temperature) and -200°C (β transformation point temperature), a strain of at least 0.1% to 8% in the longitudinal direction is applied to the profile. Furthermore, a torque is applied to the profile to keep the longitudinal twist within ±3.0°. This corrects the shape of the profile. Here, the β transformation point temperature is the temperature at which the material becomes a single β phase during heating. If the straightening temperature is too high, the warping after cooling will be large. On the other hand, if the straightening temperature is too low, a large strain is required to correct the twist, which may increase the void area ratio and reduce the tensile and fatigue properties.

[0069] Generally, twisting and warping behave differently with respect to straightening temperature, exhibiting a trade-off relationship. Specifically, warping is more likely to occur at higher straightening temperatures, while twisting is more likely to persist at lower straightening temperatures. Warping occurs because the temperature difference within the profile during high-temperature straightening causes a difference in the amount of thermal shrinkage during cooling. The higher the straightening temperature, the greater the temperature difference within the profile, and therefore the greater the warping. On the other hand, twisting is caused by the strain introduced during extrusion. The lower the straightening temperature, the greater the elastic limit of the profile. Therefore, even with the same amount of straightening, the lower the straightening temperature, the smaller the strain (plastic strain) introduced into the profile. For these reasons, lowering the straightening temperature to eliminate warping usually results in a larger amount of twisting because it is harder to correct, and raising the straightening temperature to eliminate twisting usually results in a larger amount of warping. To obtain a superior shape from α+β type titanium alloy profiles, it is necessary to simultaneously eliminate twisting and warping. To achieve this, the profile is heated and held in the α+β two-phase region from below the β transformation point temperature of 200°C to above the β transformation point temperature of -400°C. Then, a strain of at least 0.1% to 8% in the longitudinal direction is applied to the profile, and a torque is also applied to reduce the longitudinal twist to within ±3.0° to correct the shape. This is preferable as it prevents the generation of voids and achieves both the elimination of twisting and warping. This is also to avoid requiring extremely high rigidity in the straightening device, which would make it very expensive. In the straightening process, it is preferable to apply both longitudinal strain and longitudinal twist to the profile, but if the profile has a cross-sectional shape that is less prone to twisting, the application of twist may be omitted in the straightening process.

[0070] Next, the α+β type titanium alloy profile, after shape correction, is cooled to below 500°C while a predetermined stress is applied. From the viewpoint of suppressing void formation, compression is preferable. However, considering the occurrence of buckling and the practicalities of the straightening device, it is considered far easier industrially to apply tensile stress. Therefore, it is preferable to apply a tensile stress to the profile that is 20% or less of the 0.2% yield strength at room temperature. Note that the profile shrinks as the temperature decreases, so it is difficult to keep the tensile stress applied to the profile constant during cooling, but fluctuations in tensile stress are not a problem if it is 20% or less of the 0.2% yield strength. Furthermore, it is preferable to cool the profile to 500°C or below while applying a predetermined torque in addition to tensile stress. The torque is preferably a value that maintains the longitudinal torsion within ±3.0°. By satisfying these conditions, changes (deterioration) in shape due to cooling and residual stress can be reduced. Preferably, the profile is cooled to 450°C, and more preferably to 400°C.

[0071] The method for cooling the profile to below 500°C is not particularly limited as long as it can uniformly cool the titanium alloy. The cooling method may be natural cooling or accelerated cooling. Specifically, examples of accelerated cooling include cooling in a gas atmosphere such as Ar, N, H, He, and cooling in a liquid such as water or oil. To uniformly cool titanium alloys, it is preferable that the average cooling rate from the straightening temperature to 500°C be 10°C / s or less, 1°C / s or less, 0.5°C / s or less, or 0.1°C / s or less. Furthermore, as mentioned above, when the applied stress is tensile stress, it is thought that compressive stress will remain more on the surface of the profile, where the temperature drops faster than the interior. Residual stress greatly affects the processing of parts and products, but from the perspective of improving fatigue properties, compressive residual stress is generally desirable. In other words, it is desirable to adjust the residual stress to some extent. Note that the "average cooling rate from the straightening temperature to 500°C" is the value obtained by dividing the difference between the straightening temperature and 500°C by the time it took for the profile temperature to drop from the straightening temperature to 500°C.

[0072] After straightening, cooling while applying stress and / or torque concentrates stress on areas with poor shape, such as areas with significant warping and twisting. This causes localized plastic deformation in the areas where large stresses are applied. As a result, the shape is improved effectively and efficiently, residual stress is relieved, the excellent shape achieved by hot straightening at high temperatures is maintained, and the deterioration of shape such as warping and twisting that occurs during cooling is significantly suppressed, and residual stress in the titanium alloy can be reduced. Because the residual stress of the profile cooled while applying stress and / or torque is effectively and efficiently reduced, the shape change when machined after cooling is also reduced.

[0073] The torque applied during cooling can be appropriately set according to the cross-sectional shape of the profile. For example, as described above, it is preferable to use a torque that maintains the longitudinal torsion within ±3.0°. More preferably, a torque that maintains the longitudinal torsion within ±2.0°, and even more preferably within ±1.0°, is preferable. The torque during cooling may be constant, but it may also be varied within a predetermined range. For example, the torque can be increased as the temperature decreases.

[0074] As mentioned above, the stress applied during cooling is preferably a tensile stress of 20% or less of the 0.2% yield strength of the profile at room temperature. More preferably, the stress applied during cooling is 15% or less of the 0.2% yield strength of the profile at room temperature, and even more preferably 10% or less. Although there is no lower limit for the tensile stress, from the standpoint of shape and residual compressive stress, it is preferable to have a stress of 1% or more of the 0.2% yield strength of the profile at room temperature. By satisfying these conditions, plastic deformation, changes (deterioration) in shape due to cooling, and residual stress can be reduced. The tensile stress during cooling may be constant, but it may also be varied within a predetermined range. For example, a tensile stress corresponding to the elastic limit at each temperature can be applied.

[0075] Furthermore, the atmosphere during the process from the start of straightening the profile shape to cooling the profile to below 500°C is not particularly limited and may, for example, be carried out in the atmosphere. On the other hand, if necessary, it is permissible to use different atmospheres around part or all of the profile in a predetermined temperature range or the entire temperature range, as long as it does not significantly deteriorate the surface properties. For example, it is possible in this manufacturing method to apply gases such as argon and nitrogen to the atmosphere during the profile straightening process. Specifically, this could involve, for example, covering the entire apparatus with a chamber to prevent oxidation, or blowing a sealing gas onto a part of it to prevent oxidation.

[0076] Furthermore, after correcting the shape of the profile, the profile temperature may be maintained at the temperature at which the correction was performed. Alternatively, the profile temperature may be maintained in a predetermined temperature range between the β-transformation point temperature -400°C and the β-transformation point temperature -200°C. In addition, it is possible to maintain the profile in a predetermined temperature range, for example, 500 to 650°C, during the cooling of the profile. Such temperature maintenance is performed to eliminate strain generated during correction and thermal strain generated during cooling, further enhancing the properties of the profile according to this embodiment.

[0077] Standards for round bars of Ti-6Al-4V, a general-purpose α+β type titanium alloy, namely JIS H4650 and ASTM B348, specify 0.2% yield strength and elongation requirements. Specifically, these standards stipulate that a 0.2% yield strength of 828 MPa or higher and an elongation of 10% or higher are preferable. However, these standards do not specify the microstructure. Therefore, even if the microstructure of the profiled material is needle-like, the same lower limits as for equiaxed microstructure, which has a superior balance of strength and ductility, are set.

[0078] According to the method for manufacturing α+β type titanium alloy profiles of this embodiment, if the amount of straightening is 8% or less, the yield strength is improved by 0.2% without a decrease in elongation, and the strength-ductility balance satisfies ASTM B348, while the fatigue strength does not decrease. This is thought to be because when stress is applied to the titanium alloy and straightened, dislocations accumulate at the grain boundaries, strengthening the grain boundaries without forming voids.

[0079] Standard AMS2245B includes specifications for torsion and warp. Specifically, it specifies that torsion should be 1° or less per foot and within ±3.0° over the entire length, and warp should preferably be within ±0.65mm per 300mm (warp height (mm) / total length (m) = ±2.17).

[0080] However, for more complex and precise shapes, it is necessary to achieve even stricter specifications for warp and twist. Preferably, the twist should be 0.5° or less per foot and within ±2.0° over the entire length, and the warp should be within ±0.45 mm per 300 mm (warp height (mm) / total length (m) = ±1.50). The titanium alloy profiles according to this embodiment can achieve these values. The size of the α+β type titanium alloy profile according to this embodiment is not particularly limited. With currently used manufacturing equipment, the profile typically has a diameter of 30 to 300 mm and a total length of 5 to 20 m. Here, the diameter of the profile refers to the diameter of the smallest circle that can enclose the cross-section of the profile. For example, if the cross-section of the profile is triangular, the diameter of the profile is the diameter of the circumscribed circle that passes through all three vertices of the cross-section of the profile. However, in principle, even larger profiles can be manufactured by using appropriate manufacturing equipment. Alternatively, profiles with a total length of less than 5 m may be manufactured by hot extrusion. Furthermore, after a profile of 5 m or more in length is manufactured using manufacturing equipment, it may be cut to produce an α+β type titanium alloy profile with a total length of less than 5 m. In any case, the performance of the titanium alloy profile according to this embodiment is maintained. [Examples]

[0081] The following describes the effects of the present invention on improving warping and twisting after hot straightening, using examples. However, the present invention is not limited to the following examples.

[0082] (Examples) To confirm the effect of the amount of straightening on the tensile and fatigue properties of the profiles, T-shaped or U-shaped (Figure 3) test specimens, which are prone to temperature differences within the cross-section, were manufactured using titanium alloys with the components and β-transformation point temperatures shown in Table 1. The components listed in Table 1 are the values ​​measured after hot straightening, as described later. However, hot extrusion and hot straightening do not affect the compositional changes of the titanium alloy. Titanium alloys having the components shown in Table 1 have a two-phase structure consisting of α and β phases between room temperature and the β transformation point temperature. The test material was obtained by hot forging a real ingot in the α+β temperature range, and then heating the resulting φ200 billet to the β single-phase temperature range and extruding it to produce an α+β type titanium alloy profile. The β-transformation temperature of the titanium alloy was determined by the following procedure. First, the estimated β transformation point temperature Tβ of the titanium alloy was calculated by substituting the content of the element corresponding to each symbol into the elemental symbols listed in the formula below. Tβ=882+21.1×[Al]-13.9×[V]-13.9×[Fe]-9.5×[Mo]-12.1×[Cr]+23.3×[Si]+183.3×[O]+580×[C]+1040×[N] However, the above estimate is not accurate and may differ from the actual β transformation point temperature by about 100°C. Therefore, the true β transformation point temperature was determined using the following procedure. (A) The titanium alloy is heated to various temperatures within a temperature range of approximately ±50°C of the β transformation point, and then rapidly cooled. The heating temperature is varied by 5°C increments within the aforementioned range. (B) Observe the various structures obtained by (A) above and confirm whether or not needle-like structures are formed in them. If needle-like structures are confirmed in the titanium alloy, it means that the heating temperature was above the β transformation point temperature. (C) The smallest heating temperature at which a needle-like structure was obtained is considered to be the β-transformation point temperature of the titanium alloy. If needle-like structures are obtained at all heating temperatures, repeat steps (A) to (C) above in a temperature range lower than the temperature range mentioned above. If needle-like structures are not obtained at all heating temperatures, repeat steps (A) to (C) above in a temperature range higher than the temperature range mentioned above.

[0083] [Table 1]

[0084] The obtained profiles were subjected to hot straightening under various straightening conditions. Subsequently, the profiles were allowed to cool in the air to a predetermined temperature. Table 2 shows the straightening temperature (see "Straightening Temperature" column), the amount of longitudinal strain applied to the profile at the straightening temperature (see "Longitudinal Strain" column), the longitudinal torsion applied to the profile at the straightening temperature (see "Twist Angle Before Cooling" column), and the cooling stop temperature (see "Removal Temperature" column). Furthermore, Table 2 also shows the tensile stress applied to the profile during cooling (see "Tensile Stress During Cooling" column), the holding temperature of the profile during cooling (see "Holding Temperature During Cooling" column), and the average cooling rate of the profile from the straightening temperature to 500°C (see "Cooling Rate" column). For example, in No. 1, a tensile stress of 75 MPa and a torque that produced a torsion of -0.6° were applied to the test material at a straightening temperature of 720°C. The test material was then allowed to cool to 400°C while the tensile stress and torque were still applied, and then removed from the tensioning device. The average cooling rate of the profile from the straightening temperature to 500°C was 0.5°C / second. On the other hand, in No. 17, no torque was applied before cooling. Therefore, in No. 17, "None" was written in the "Presence or Absence of Torque During Cooling" column. Also, in No. 17, a "-" was written in the "Twist Angle During Cooling" column to indicate that no torque was applied. In Example 18, cooling during the straightening process is interrupted at 600°C, and the temperature is held there. In this Example 18, the average cooling rate when the surface temperature of the profile decreases from 720°C to 600°C is set to 0.5°C / second, the surface temperature of the profile is held at 600°C for 10 minutes, and the average cooling rate when the surface temperature of the profile decreases from 600°C to 500°C is set to 0.5°C / second.

[0085] <Fatigue Test> For fatigue strength measurement, a round bar test specimen with a diameter of 5.08 mm and a parallel section length of 15.24 mm was prepared from the area marked with a circle in Figure 3, ensuring that the region less than 0.5 mm deep from the surface of the profile was not included. Using this round bar test specimen, a fatigue test was performed under conditions of stress ratio R = σmin / σmax = 0.1 (both σmin and σmax are tensile stresses) and room temperature, with a number of cycles of 1.0 × 10⁻⁶. 7The fatigue strength was defined as the maximum value of σmax, the strength at which the object did not break in a given number of cycles.

[0086] Furthermore, for each profile, the 0.2% yield strength, elongation, void area ratio, maximum residual stress, maximum dimensional change after machining, torsion after straightening, and warp after straightening were measured. The results are shown in Table 2. 0.2% yield strength and elongation were measured in accordance with ASTM E8, using test specimens with a parallel section length of 28 mm and a diameter of 6.25 mm taken from the profile, with a yield strength of 0.005 / min (8.3 × 10) at a strain of 2% or less. -5 ( / s), 0.1 / min (1.67 × 10) when strain is 2% or more. -3 The evaluation was performed by conducting tensile tests at a speed of ( / s). The sampling locations for the test specimens were indicated by circles in Figure 3. Furthermore, similar results for fatigue strength, 0.2% yield strength, and elongation can be obtained by preparing the test specimens in such a way that they do not include any region less than 0.5 mm deep from the surface. The void area ratio and the maximum residual stress were evaluated using the method described above. The maximum dimensional change after machining was determined by measuring the dimensional change caused by the tilting of a portion of the cross-section during machining using calipers, as shown in Figures 6A and 6B, and evaluating the maximum value. Figure 6A is a cross-sectional view of the example labeled "T" in the "Shape" column, and Figure 6B is a cross-sectional view of the example labeled "U" in the "Shape" column. The curvature and twist after correction were measured according to the AMS2245B standard. In Table 2, items that are underlined are outside the scope of the present invention.

[0087] Furthermore, the microstructure was also evaluated. The microstructure evaluation was carried out using the following procedure. First, a specimen for microstructure observation was taken from a cross-section half the length of the obtained profile. The dimensions of the specimen for microstructure observation are not particularly limited, and for profiles with a small cross-sectional area, the cross-section itself may be used as the specimen for microstructure observation. Here, a 10 mm square specimen was taken. The sampling location was the circled position in Figure 3. The observation surface of this specimen was mirror-polished, and then etched using hydrofluoric acid to reveal the microstructure. The microstructure was then observed using an optical microscope. As a result of the observation, it was confirmed that the metal structure in all examples was needle-like. At the circled position in Figure 3, located at a depth of 0.5 mm or more from the surface of the profile, the metal structure was confirmed to be needle-like. On the other hand, in the region less than 0.5 mm from the surface of the profile, i.e., the surface region, areas where needle-like and equiaxed structures were mixed were observed. This is presumed to be because, during hot extrusion, the temperature tends to decrease in the surface region, and processing strain was introduced in a temperature range below the β transformation point temperature. Furthermore, in the region at a depth of 0.5 mm or more from the surface of the profile, it is preferable that the metal structure is mainly composed of needle-like structures, and particularly preferable that the proportion of needle-like structures in the metal structure is 100%. For the corroded specimens described above, an optical microscope with a magnification of 200x was used to measure a 7mm square measurement area. Fifteen straight lines were drawn vertically and horizontally at equal intervals within the measurement area, and the original β particle size was measured using the sectioning method. In all examples, the average prior β particle size was 500 μm or less. In one example, the average prior β particle size was 60 μm.

[0088] [Table 2]

[0089] [Table 3]

[0090] Examples No. 1-13 and No. 18-20 are examples of straightening and then cooling while applying torque to maintain the angle and tensile stress, using various alloy types and temperatures. In all cases, the elongation during straightening, i.e., the longitudinal strain, was 8% or less, resulting in a 0.2% yield strength of 830 MPa or more, elongation of 10% or more, and fatigue strength of 450 MPa or more, demonstrating an excellent balance of strength and ductility and fatigue characteristics. Furthermore, because torque and tensile stress were applied during cooling after hot straightening, the maximum residual stress exceeded +400 MPa, the maximum dimensional change after machining was 2.5 mm or less, and profiles with excellent shape characteristics, such as minimal warping and twisting, were obtained even in the hot-straightened state. Note that in No. 13, the longitudinal strain during straightening was at the upper limit of 8%, resulting in slight void formation, but good mechanical properties were ensured. In No. 24, a tensile stress exceeding 20% ​​of the 0.2% yield strength at room temperature was applied to the profile, resulting in slight void formation, but good mechanical properties were ensured. In No. 25, the profile was forcibly cooled using mist while held in a straightening machine, resulting in a high cooling rate and consequently an increase in the maximum residual stress within the cross-section. However, even in the profile of No. 25, warping and torsion were minimal, and the yield strength, fatigue strength, and ductility were within acceptable limits.

[0091] Numbers 14 and 15 are profiles that were cooled without applying tensile stress after hot straightening, maintaining their straightened length until cooled to a low temperature, and then removed from the straightening machine. As a result, No. 14, which was hot straightened at a temperature of 720°C, had a post-straightening warp (mm) / total length (m) outside the range of ±2.17, a maximum residual stress exceeding +400 MPa, and a maximum dimensional change of more than 2.0 mm before and after cutting. On the other hand, No. 15, which was hot straightened at a temperature of 600°C, lower than the β transformation point temperature of -400°C, did not have its torsion corrected with a 3% elongation, and even after cooling, the torsion was outside the range of ±3.0°.

[0092] In No. 16, the removal temperature after hot straightening exceeded 500°C, resulting in thermal stress during cooling. The maximum residual stress exceeded +400 MPa, and the maximum values ​​of warping after cooling and dimensional changes before and after cutting exceeded 2.0 mm.

[0093] In No. 17, no torque was applied to maintain the torsion during cooling. As a result, there was significant springback during cooling, and the torsion of the profile after cooling was outside the ±3.0° range.

[0094] Because No. 18 was held at 500-650°C during cooling, and because No. 19 had a cooling rate of less than 0.1°C / s, the yield strength was 0.2% higher than other alloys with the same alloy composition, exceeding 850 MPa.

[0095] For samples No. 21-22, the longitudinal strain during straightening exceeded 8%, resulting in numerous voids within the profile, and depending on the alloy type, the ductility or fatigue strength fell below the lower limit.

[0096] In No. 23, straightening was not performed in the two-phase region, and twisting was applied during cooling, resulting in significant twisting after the straightening process.

[0097] According to the present invention, it has been confirmed that the shape of α+β type titanium alloy, which has excellent strength, elongation, and fatigue properties, can be significantly improved by shape correction, and that stable manufacturing is possible. [Explanation of Symbols]

[0098] 1 container 2 stems 3 Dummy Blocks 4 dice 5 Billets 6 Shaped materials 11 Extrusion direction

Claims

1. In mass%, Al: 4.4-6.5%, Fe: 0.5-2.9%, Si: 0 to 0.50%, O: 0-0.25%, C: 0-0.08%, N: 0-0.05%, Ni: 0 to 0.15%, Cr: 0-0.25%, and Mn: 0-0.25% It contains Ti and impurities, The content of Fe, Ni, Cr, and Mn, expressed in mass percent, satisfies the condition 0.5% ≤ %Fe + %Ni + %Cr + %Mn ≤ 2.9%. The 0.2% proof strength is 830 MPa or higher, the elongation is 10% or higher, and the fatigue strength is 450 MPa or higher. It has needle-like tissue, The area ratio of the void is 1.0 × 10⁻⁶. －５ It is less than %. The twist angle from one end to the other is within ±3.0°. An α+β type titanium alloy profile with a curvature height (mm) / total length (m) within ±2.

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2. In mass%, Al: 4.4-5.5%, Fe: 1.4-2.3%, Mo: 1.5-5.5%, O: 0-0.20%, C: 0-0.08%, N: 0-0.05%, Si: 0 to 0.10%, Ni: 0 to 0.15%, Cr: 0-0.25%, and Mn: 0-0.25% It contains Ti and impurities, The content of Fe, Ni, Cr, and Mn, expressed in mass percent, satisfies the condition 1.4% ≤ %Fe + %Ni + %Cr + %Mn ≤ 2.3%. The 0.2% proof strength is 830 MPa or higher, the elongation is 10% or higher, and the fatigue strength is 450 MPa or higher. It has needle-like tissue, The area ratio of the void is 1.0 × 10⁻⁶. －５ It is less than %. The twist angle from one end to the other is within ±3.0°. An α+β type titanium alloy profile with a curvature height (mm) / total length (m) within ±2.

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3. The α+β type titanium alloy profile according to claim 1 or 2, characterized in that the aforementioned 0.2% yield strength is 850 MPa or more.

4. The area ratio of the void is 1.0 × 10 －６ An α+β type titanium alloy profile according to any one of claims 1 to 3, characterized in that it is less than or equal to %.

5. An α+β type titanium alloy profile according to any one of claims 1 to 4, characterized in that the average prior β particle size is 500 μm or less.

6. An α+β type titanium alloy profile according to any one of claims 1 to 5, characterized in that the maximum residual stress within the cross-section is +400 MPa or less.

7. An α+β type titanium alloy profile according to any one of claims 1 to 6, characterized in that it is an extruded profile.

8. A method for manufacturing an α+β type titanium alloy profile according to any one of claims 1 to 7, The process involves hot working an α+β type titanium alloy to obtain a profiled material, The process involves heating the profile to a straightening temperature between the β-transformation point temperature of -400°C or higher and the β-transformation point temperature of -200°C or lower, applying a strain of 0.1% to 8% in the longitudinal direction at the straightening temperature, and further applying a torque that reduces the longitudinal twist of the profile to within ±3.0%. A step of cooling the profile to 500°C or below while applying tensile stress and torque to the profile, A method for manufacturing an α+β type titanium alloy profile, characterized by comprising the following features.

9. The method for manufacturing an α+β type titanium alloy profile according to claim 8, characterized in that the tensile stress applied to the profile during cooling is 20% or less of the 0.2% yield strength at room temperature.

10. Furthermore, the method for manufacturing an α+β type titanium alloy profile according to claim 8 or 9 is characterized by comprising a step of holding the profile at 500 to 650°C during cooling.

11. A method for producing an α+β type titanium alloy profile according to any one of claims 8 to 10, characterized in that the average cooling rate of the profile from the straightening temperature to 500°C is 10°C / s or less.

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