CN118434893A - Alpha+beta titanium alloy section bar and manufacturing method thereof - Google Patents

Abstract

An aspect of the present invention provides an α+β titanium alloy material, wherein the α+β titanium alloy material has a 0.2% yield strength of 830MPa or more, an elongation of 10% or more, and a fatigue strength of 450MPa or more, has a needle-like structure, has an area ratio of a void of 1.0x10 ^‑5% or less, has a torsion angle from one end to the other end of within ±3.0°, and has a warp height (mm)/total length (m) of ±2.17.

Description

Alpha+beta titanium alloy section bar and manufacturing method thereof

Technical Field

The invention relates to an alpha+beta titanium alloy section and a manufacturing method thereof.

Background

Titanium alloys are used for various applications such as connecting rods for automobiles, muffler parts, golf club heads, and construction materials, because of their high specific strength and excellent corrosion resistance. Titanium alloys have also been widely used for medical applications such as ornaments for clocks and spectacle frames, and implants, because of their good compatibility with living bodies.

Titanium alloys are classified into three kinds of alloys of alpha type, beta type and alpha + beta type. The α -type titanium alloy is a titanium alloy whose metallographic structure at normal temperature is mainly composed of an α -phase, the β -type titanium alloy is a titanium alloy whose metallographic structure at normal temperature is mainly composed of a β -phase, and the α+β -type titanium alloy is a titanium alloy whose metallographic structure at normal temperature is composed of both an α -phase and a β -phase. The alpha phase of the titanium alloy refers to a phase having a hexagonal closest packing structure (hcp), and the beta phase of the titanium alloy refers to a phase having a body centered cubic structure (bcc).

In pure titanium, the metallographic structure at normal temperature is all alpha phase, and the metallographic structure at a temperature exceeding the beta transformation temperature is all beta phase. However, by adding an alloy element for stabilizing the β phase to pure titanium, the temperature at which the β phase can exist stably is reduced, and the β phase remains even at normal temperature. The temperature region exceeding the β transformation temperature and having a metallographic structure of β single phase is referred to as β single phase temperature region. In addition, a temperature region below the β transformation temperature and the metallographic structure including both the α phase and the β phase is referred to as an α+β two-phase temperature region.

Among titanium alloys, α+β type titanium alloys also have excellent balance of strength/ductility and fatigue characteristics, and thus have long experience in use. As one of the uses of α+β titanium alloys, recently, attention has been paid to reduction in fuel consumption due to weight reduction in the fields of automobiles and motorcycles. The α+β titanium alloy is about 60% by weight of carbon steel or stainless steel, and has strength and fatigue characteristics substantially equivalent to those of the same. Therefore, by replacing carbon steel or stainless steel crankshafts and engine components with components made of an α+β titanium alloy, the weight of the entire engine can be reduced, and the output can be improved and the fuel consumption can be reduced.

In general, materials used as engine parts such as crankshafts, valves, and connecting rods are hot-worked products such as round bars, billets, square bars, and stretched materials supplied in a shape having a uniform cross section. Hereinafter, these materials are collectively referred to as "profile". That is, the "profile" means a stretched material having a uniform cross section over the entire length and supplied in a straight line shape. The profile has bending, torsion in the stage of manufacture by extrusion, rolling, etc., and is therefore usually removed by straightening, annealing, etc. The α+β titanium alloy profile can be also used in the above-mentioned wide variety of applications by cutting, further processing, cutting after further processing, and the like.

The α+β titanium alloy can realize excellent balance of strength and ductility and fatigue characteristics by controlling the metallographic structure to an equiaxed grain structure by performing strong working in an α+β two-phase temperature range of β transformation temperature or lower. Conventionally, a complicated shape profile is manufactured by cutting a forged material or a thick plate manufactured by performing a strong working in an α+β two-phase temperature range.

However, since the cutting process increases the manufacturing cost of the machine component, it is preferable to reduce the number of the machine components as much as possible. Now, in order to reduce the manufacturing cost, the following manufacturing technique is being developed: it is possible to manufacture a long bar-shaped profile having a sectional shape closer to that of the final product, so as to improve production efficiency. By using a profile having a cross section close to the final product as a material, the cutting amount can be reduced and the yield can be improved. In addition, by increasing the length of the profile, productivity can be improved.

For example, in a high-speed extrusion method of a glass lubricant, which is one of extrusion processes, an ingot or a round billet obtained by hot forging the ingot is used as a raw material. As shown in fig. 1, a raw material (billet 5) is inserted into a container 1, a load by hydraulic pressure is applied to a rod 2, and the billet 5 is pushed in an extrusion direction 11 via an extrusion pad 3 and is molded into various cross-sectional shapes through a die 4, whereby a long profile 6 can be obtained.

When processed into such a complex shape, the resistance to thermal deformation of the α+β type titanium alloy in a temperature region lower than the β transformation temperature is abruptly increased. Thus, in order to process an α+β type titanium alloy in a temperature region lower than the β transformation temperature and control the structure to an equiaxed structure, a large-sized apparatus capable of applying a high load is required. Thereby, the equipment cost becomes high. In addition, depending on the desired cross-sectional shape, the α+β type titanium alloy may not be hot worked in a temperature region below the β transformation temperature. Even when the processing is possible, if the local temperature in the section of the profile exceeds the β -transformation temperature due to the processing heat, the equiaxed structure and the needle-like structure obtained by the processing at or above the β -transformation temperature are mixed in the section of the profile, and a significant difference in mechanical properties occurs in the section. Therefore, in general, when an α+β titanium alloy is to be hot-worked into a complex shape, it can be produced with a low heat distortion resistance, and the metallographic structure can be controlled to a fine needle structure by working at a β transformation temperature or higher at which surface defects are less likely to occur and forcibly cooling the alloy as needed, thereby achieving the desired excellent balance of strength and ductility and fatigue characteristics.

In general, bending and torsion are generated in the profile after the hot working. Bending and torsion are caused by residual processing strain in the profile and by the cooling rate of the profile being different at each location. Therefore, in the production of the profile, further correction processing is required after the above-mentioned hot working.

The straightening processing method includes roll straightening and stretching straightening. The roll straightening is a method of passing a profile through a plurality of rolls to bend and return the profile to bending, thereby improving bending and torsion of the profile. The stretching correction is a method of fixing both ends of a profile and applying a stretching load to deform the profile into a plastic deformation region, thereby improving bending and torsion of the profile.

Since the α+β titanium alloy has a high cold deformation resistance, the straightening process is also generally performed as a hot process. In the shape of an α+β titanium alloy having a needle-like structure, the heat straightening can be performed at a lower load than at room temperature. In addition, since the α+β titanium alloy has a narrower elastic region at a high temperature than at room temperature, it can be corrected with a small deformation amount. In addition, in the α+β titanium alloy, the soft β phase fraction is higher than room temperature at high temperature, and the strain is easily recovered.

However, when the α+β titanium alloy material is corrected so that the deformation amount is large in the low temperature region of the α+β two-phase temperature region, a plurality of internal defects, that is, coarse voids, are formed. This causes a problem that excellent balance of strength and ductility and fatigue characteristics required for the profile cannot be achieved. And, the Young's modulus of the alpha+beta type titanium alloy is low. Therefore, when the α+β titanium alloy is subjected to hot straightening, a large spring back occurs after cooling to room temperature, and torsion cannot be sufficiently improved, which is often the case.

For the above reasons, it is now that the α+β titanium alloy shape requiring shape correction does not necessarily satisfy the required excellent balance of strength and ductility, fatigue characteristics, and excellent shape.

Patent document 1 discloses a correction method for an α+β titanium alloy as follows: the blank, square bar, tube, etc. after the time hardening is subjected to a tensile stress at a temperature about 0.3 times the melting temperature (about 480 ℃) to be corrected, and further cooled while applying a tensile stress.

Patent document 2 discloses a correction method for an α+β titanium alloy as follows: after annealing for adjusting the structure is performed on the rolled round titanium alloy rod, bending correction by press correction and warm correction within 600-beta transformation temperature are performed.

Prior art literature

Patent literature

Patent document 1: japanese patent No. 6058535

Patent document 2: japanese patent application laid-open No. 2011-137204

Disclosure of Invention

Problems to be solved by the invention

In the thermal correction of a general α+β titanium alloy, correction is performed as follows: straightening 2-3% in a high temperature region (about 700-740 ℃), maintaining the length of the titanium alloy in the high temperature region for about 10-30 minutes while simultaneously serving as annealing, removing the load, cooling to about 400-600 ℃, and taking the titanium alloy from the straightening machine and naturally cooling. If the profile is circular, the cooling rate inside the cross section is uniform, so that the difference in thermal shrinkage is small, and warpage is not easy to occur during natural cooling after correction.

However, in the case where the shape of the profile is a quadrangle, or in the case where the profile is a shape close to the product shape and is a complex shape, for example, the difference in plate thickness due to the presence of corners having a relatively high cooling rate increases the temperature difference in the cross section immediately before heating and natural cooling, and the difference in heat shrinkage increases. As a result, the following problems occur: the warpage occurs in the profile after being removed from the leveler and cooled, or even in the case where the warpage does not occur after cooling, the residual stress in the interior of the profile is large, and the warpage occurs when the profile is cut after cooling. Fig. 2 (a) shows an example of a conventional simple-shaped profile of an α+β titanium alloy profile, and fig. 2 (b) shows an example of a profile having a cross-sectional shape closer to the shape of a product. Fig. 2 is a cross-sectional view of a profile extending from the surface of the paper to the back.

Fig. 3 to 6B are schematic views showing the shape of the profile, which is a problem, and the shape change during cutting. Fig. 3 is an example of the cross-sectional shape of the profile after heat correction. Fig. 4 shows an example of warpage in a profile. Fig. 5 shows an example of torsion in the profile. Fig. 6A and 6B illustrate an example of dimensional change after cutting. The warp and twist in fig. 4 and 5 are required to be within a limit of ±s above and below the paper surface, and the dimensional change in fig. 6A and 6B is required to be within a limit of ±s around the paper surface.

In patent document 1, after heat treatment of a titanium alloy in an α+β two-phase region, a tensile stress is applied during cooling, and thereby the stress generated during cooling is balanced, and titanium alloy products are corrected. In this technique, however, the tensile stress applied to the profile for straightening is set to at least 20% of the yield stress of the profile at the straightening temperature. If such a large tensile stress is applied to a profile having a needle-like structure, the stress concentrates on a portion where torsion or bending is large, and a void may be generated inside. Further, since the titanium alloy described in patent document 1 is rolled at a temperature in the α+β two-phase region, it is assumed that the structure is not a needle-like structure but mainly composed of equiaxed grains.

In patent document 2, after annealing for adjusting the structure of the round titanium alloy rod, final straightening is performed by bending straightening by press straightening, cutting removal of surface flaws, and warm straightening at a 600 to β transformation temperature, and then surface polishing is performed, thereby producing a round titanium alloy rod product. However, in the roll straightening, the conveyance time from the heating device to the roll straightening machine is different in the longitudinal direction, and therefore, a difference in cross-sectional dimension occurs in the longitudinal direction, and productivity and overall yield are reduced.

As means for reducing warpage, a method of straightening at a low temperature in which a temperature difference in a cross section is small is considered. But the elastic limit is large at low temperatures, so that the elimination of torsion requires a very large amount of strain. In addition, since the α+β titanium alloy having a needle structure causes stress concentration at β grain boundaries under tensile stress, voids are generated inside, which becomes a starting point of fatigue fracture, fatigue characteristics are significantly reduced, and balance of strength and ductility is also poor. That is, there are many problems in the conventional method in terms of manufacturing the α+β titanium alloy material into a complex shape that is closer to the shape of the product.

In view of the above, an object of the present invention is to provide an α+β titanium alloy material which is less in warpage and torsion and is excellent in conditional yield strength, fatigue strength and ductility, and a method for producing the same.

Solution for solving the problem

The gist of the present invention is as follows.

(1) An aspect of the present invention provides an α+β titanium alloy material, wherein the α+β titanium alloy material has a 0.2% yield strength of 830MPa or more, an elongation of 10.0% or more, and a fatigue strength of 450MPa or more, has a needle-like structure, has a void area ratio of 1.0×10 ^-5% or less, has a torsion angle from one end to the other end of within ±3.0°, and has a warp height (mm)/total length (m) of ±2.17.

(2) The α+β titanium alloy material according to (1) above, wherein the 0.2% yield strength is 850MPa or more.

(3) The α+β titanium alloy material according to (1) or (2), wherein the area ratio of the voids is 1.0X10 ^-6% or less.

(4) The α+β -type titanium alloy material according to any one of (1) to (3) above, wherein the average primary β -grain diameter is 500 μm or less.

(5) The α+β titanium alloy material according to any one of (1) to (4) above, wherein the maximum residual stress in the cross section is +400MPa or less.

(6) The α+β titanium alloy material according to any one of the above (1) to (5) may be an extrusion material.

(7) The α+β titanium alloy material according to any one of (1) to (6) above, wherein the material contains, in mass%, al:4.4 to 6.5 percent of Fe:0.5 to 2.9 percent of Si:0 to 0.50 percent, O:0 to 0.25 percent, C:0 to 0.08 percent, N:0 to 0.05 percent of Ni:0 to 0.15 percent, cr:0 to 0.25%, and Mn:0 to 0.25 percent, and the balance of Ti and impurities, wherein the content of Fe, ni, cr, mn percent of Fe,% of Ni,% of Cr and% of Mn in mass percent satisfies 0.5 percent to less than or equal to percent of Fe+% of Ni+% of Cr+% of Mn to less than or equal to 2.9 percent.

(8) The α+β titanium alloy material according to any one of (1) to (6) above, wherein the material contains, in mass%, al:4.4 to 5.5 percent of Fe:1.4 to 2.3 percent of Mo:1.5 to 5.5, O:0 to 0.20 percent, C:0 to 0.08 percent, N:0 to 0.05 percent of Si:0 to 0.10 percent of Ni:0 to 0.15 percent, cr:0 to 0.25%, and Mn:0 to 0.25 percent, and the balance of Ti and impurities, wherein the content of Fe, ni, cr, mn percent of Fe,% Ni,% Cr and% Mn in mass percent satisfies 1.4 percent to less than or equal to 1.4 percent of Fe+% Ni+% Cr+% Mn to less than or equal to 2.3 percent.

(9) Another aspect of the present invention provides a method for producing an α+β titanium alloy material according to any one of (1) to (8), comprising: a step of hot working an alpha+beta titanium alloy to obtain a profile; heating the profile to a correction temperature of a beta transus temperature of-400 ℃ or higher and a beta transus temperature of-200 ℃ or lower, applying a strain of 0.1% or higher and 8% or lower in the longitudinal direction at the correction temperature, and further applying a torque for making the torsion of the profile in the longitudinal direction within + -3.0%; and cooling to 500 ℃ or lower while applying a tensile stress and the torque to the profile.

(10) The method for producing an α+β titanium alloy material according to (9) above, wherein the tensile stress applied during cooling of the material is 20% or less of the 0.2% yield strength at room temperature.

(11) The method for producing an α+β titanium alloy material according to (9) or (10), further comprising the step of maintaining the material at 500 to 650 ℃ during cooling of the material.

(12) The method for producing an α+β titanium alloy material according to any one of (9) to (11), wherein an average cooling rate of the material from the straightening temperature to 500 ℃ is 10 ℃/s or less.

ADVANTAGEOUS EFFECTS OF INVENTION

According to the present invention, an α+β titanium alloy material having less warpage and torsion and excellent conditional yield strength, fatigue strength and ductility, and a method for producing the same can be provided.

Drawings

Fig. 1 is a view illustrating a high-speed extrusion method of a glass lubricant.

Fig. 2 shows an example of the cross-sectional shape of an α+β type titanium alloy material, (a) shows an example of a simple shape material, and (b) shows an example of a material closer to the shape of a product.

Fig. 3 is an example of the cross-sectional shape of the profile after heat correction.

Fig. 4 shows an example of warpage of a profile.

Fig. 5 shows an example of the torsion of the profile.

Fig. 6A shows an example of dimensional change after cutting of a profile having a T-shaped cross section.

Fig. 6B shows an example of dimensional change after cutting of a profile having a U-shaped cross section.

Fig. 7A is a photomicrograph of an example of a needle-like structure.

Fig. 7B is a photomicrograph of an example of an equiaxed structure.

Fig. 8 shows an example of voids included in an α+β titanium alloy profile.

Fig. 9A is an example of the cross-sectional shape of an α+β titanium alloy profile.

Fig. 9B is an example of the cross-sectional shape of an α+β titanium alloy profile.

Fig. 9C is an example of the cross-sectional shape of an α+β titanium alloy profile.

Detailed Description

The present inventors have conducted intensive studies on a method for obtaining an α+β titanium alloy material having less warpage and torsion and excellent yield strength, fatigue strength and ductility. As a result, it was found that the above problems can be solved by straightening an α+β titanium alloy material under predetermined heating conditions and stress conditions, and then cooling the material to a low temperature while applying a tensile load.

In addition, in the past, when warpage and torsion cannot be eliminated by correction processing, the portions including warpage and torsion are discarded, which reduces the manufacturing yield of the mechanical component. Therefore, when warpage and torsion are eliminated, effects such as improvement in the manufacturing yield of the mechanical component and reduction in cost can be obtained.

The following describes in detail an α+β titanium alloy material according to an embodiment of the present invention.

The α+β titanium alloy material of the present embodiment has excellent balance of strength and ductility and fatigue strength. Further, even in the case of a complex shape close to a product, it has an excellent shape. Specifically, the 0.2% yield strength is 830MPa or more, preferably 840MPa or more, 850MPa or more, the total elongation is 10.0% or more, preferably 10.3% or more, preferably 12.0% or more, and the fatigue strength is 450MPa or more, preferably 480MPa or more. The upper limit of the 0.2 yield strength is not particularly limited, and for example, the 0.2% yield strength may be 1400MPa or less, 1300MPa or less, or 1200MPa or less. The upper limit of the total elongation is not particularly limited, and for example, the total elongation may be 30% or less, 28% or less, or 25% or less. The upper limit of the fatigue strength is not particularly limited, and for example, the fatigue strength may be 800MPa or less, 750MPa or less, or 700MPa or less. Such mechanical properties can be achieved by a manufacturing method described later, for example. In the following description, the "total elongation" may be simply referred to as "elongation".

The profile of the present embodiment is a so-called α+β titanium alloy. The α+β titanium alloy is composed of an α phase having a HCP structure and a β phase having a BCC structure at a temperature lower than the β transformation temperature. When the α+β titanium alloy is equal to or higher than the β transformation temperature, the α phase changes to the β phase, and the α+β titanium alloy is composed of only the β phase.

The α+β titanium alloy material of the present embodiment has a needle structure. The needle structure is a structure form generated when the α+β titanium alloy is cooled from a temperature equal to or higher than the β transformation temperature. As illustrated in fig. 7A, in the needle-like structure, a grain boundary α phase is formed at the original β grain boundary, and a structure in which α phase and β phase are arranged in layers is formed in the original β grain. Here, the original β grain boundary refers to a trace of a grain boundary of a β phase existing at a temperature equal to or higher than the β transformation temperature. When an α+β titanium alloy in a temperature range equal to or higher than the β transformation temperature is cooled, an α phase is preferentially precipitated at the grain boundary of the original β phase, and becomes a grain boundary α phase.

The α+β titanium alloy material according to the present embodiment may be composed of only a needle structure, or may have both a needle structure and an equiaxed structure. The equiaxed structure is a structure that does not have the features of the needle-like structure, and is composed of alpha grains and a phase change beta phase, for example, as illustrated in fig. 7B. The phase transition β phase (Transformed β) is a layered structure of β phase and α phase, and is a structure of β phase in hot working and β phase and α phase in cooling after hot working.

Alpha + beta titanium alloys with needle-like structure tend to void at the original beta grain boundaries under tensile stress. It is considered that the stress concentration coefficient in the vicinity of the voids in the metal is high under tensile stress, and the stress concentration coefficient tends to become a starting point of fracture in a tensile test or a fatigue test, and the elongation is lowered, and the 0.2% yield strength and the fatigue strength are also lowered.

In the conventional straightening technique, stretching straightening and roll straightening are performed in a high-temperature region of an α+β two-phase temperature region having a narrow elastic region so as not to generate a void, and then natural cooling is performed, or fixing is performed with a mold so as to suppress deformation due to thermal stress during cooling, and cooling is performed. Thus, both material properties and dimensional accuracy are obtained. However, when the profile has an end portion with a relatively high cooling rate or when the profile has a complicated shape, for example, a difference in cooling rate or plate thickness increases in temperature in a cross section during heating, and a difference in heat shrinkage during cooling increases. Therefore, when cooling in the straightening process, a significant stress is generated in the profile, and plastic deformation such as warpage and torsion is caused to the profile. In addition, even when plastic deformation is not generated during cooling in the straightening process or when plastic deformation is suppressed by the mold, residual stress is generated inside the profile, and thus, shape defects such as warpage are caused during cutting processing of the profile.

In the α+β type titanium alloy material according to the present embodiment, the average value of the crystal grain diameters of the regions surrounded by the original β grain boundaries, that is, the average original β crystal grain diameter is not particularly limited, and is preferably 500 μm or less, for example. This can further improve the mechanical properties of the profile. In addition, in the case of producing a profile by hot rolling followed by corrective working, the average raw β grain diameter is generally greater than 500 μm. Hot rolling is typically performed in multiple passes using a hot rolling apparatus, with beta grains growing as the titanium alloy moves between the passes. Further, since the amount of strain introduced in each pass during hot rolling is small, when the titanium alloy moves between the passes, strain is recovered, and recrystallization is difficult to occur. For these reasons, the average original β grain diameter of the titanium alloy shape obtained by hot rolling may coarsen. On the other hand, for example, in the case of forming a profile by hot extrusion and then performing correction processing, the average original β crystal grain diameter is 500 μm or less. The average raw β crystal grain diameter may be 450 μm or less, 400 μm or less, or 350 μm or less. The lower limit of the average raw β crystal grain diameter is not particularly limited, and for example, the average raw β crystal grain diameter may be 50 μm or more or 80 μm or more. Since the cross section of the profile cannot be made into a complicated shape by hot rolling, the α+β titanium alloy profile of the present embodiment is preferably an extrusion profile produced by hot extrusion.

The method for measuring the average raw β grain diameter is as follows. First, a test piece for tissue observation was collected from the inside of the profile. The test piece was set to have no depth region of 0.5mm from the surface of the profile. Then, the test piece was subjected to wet polishing and finish polishing to set the polished surface to a mirror surface state. The polished surface was etched with fluoronitric acid to develop a metallographic structure. When the etched surface is observed by an optical microscope, the grain boundary α phase shown in fig. 7A can be determined. The grain boundary alpha phase can be considered as the original beta grain boundary. The average value of the original β crystal grain diameter can be obtained by measuring the equivalent circle diameter of the portion surrounded by the original β crystal grain boundary by the cutting method. When the grain diameter is measured by the cutting method, the size of the measurement region is 7mm square, the measurement magnification is 50 to 200 times, and 15 lines each in the longitudinal and transverse directions are recorded at equal intervals in the measurement region.

The area ratio of the voids in the α+β titanium alloy material of the present embodiment is 1.0x10 ^-5% or less, and the structure is a needle-like structure. As a result, a high balance of strength and ductility and fatigue properties can be obtained. The term "needle-like tissue" means that 50% or more of the tissue is needle-like tissue, preferably 60% or more, more preferably 70% or more, and most preferably 100% of the tissue is needle-like tissue. The tissue other than needle-like tissue is equiaxed. The area ratio of the voids is preferably 5.0X10 ^-6% or less, 1.0X10 ^-6% or less, or 1.0X10 ^-7% or less. As described above, the void becomes a starting point of fatigue fracture, and the fatigue characteristics are significantly reduced. In addition, voids adversely affect strength and ductility. Therefore, the smaller the area ratio of the voids is, the more preferable. The lower limit of the area ratio of the void is not particularly limited, and may be, for example, 0%. The area ratio of the voids may be set to 1.0X10 ^-9% or more, 5.0X110 ^-9% or more, or 1.0X10 ^-8% or more.

The needle-like structure and the equiaxed structure are described in detail below. The needle-like structure is a structure composed of a lamellar arrangement of an α phase and a β phase as shown in fig. 7A, for example, and a relatively thin grain boundary α phase surrounding the structure. The equiaxed structure is a structure composed of, for example, granular pre-eutectoid α phase or phase transition β not surrounded by grain boundary α phase as shown in fig. 7B. If the α+β titanium alloy is naturally cooled without machining after being heated to the β temperature region, the structure thereof is mainly a needle-like structure. On the other hand, when an α+β titanium alloy is hot worked in an α+β two-phase temperature range and then cooled, an equiaxed structure is generated. The amounts of needle-like structures and equiaxed structures vary depending on the temperature of the hot working, the amount of strain applied, and the like.

The amount of needle-like tissue was measured as follows. First, a test piece for tissue observation was collected from the inside of the profile. The test piece was set to have no depth region of 0.5mm from the surface of the profile. Then, the test piece was polished, and the polished surface was etched with fluoronitric acid to develop a metallographic structure. When the etched surface is observed by an optical microscope, the grain boundary α phase shown in fig. 7A can be determined. The ratio of the area ratio of the region surrounded by the grain boundary α and the grain boundary α phase included in the measurement field to the total area of the measurement field is regarded as the area ratio of the needle-like structure of the profile.

After mirror polishing the samples collected from the cross section and the longitudinal section of the profile, the voids were observed using an optical microscope. The sample was set to contain no depth region from the surface of the profile to a depth of 0.5 mm. The voids are the sum of the areas of voids having an equivalent circle diameter of 1 μm or more in terms of area. When observed by an optical microscope, for example, as shown in fig. 8, the void can be visually recognized as a circular dark region.

In the method of measuring residual stress, the measurement was performed by a 2D method using an X-ray diffraction method (Bruker AXSD Discover). The tube ball was Cu, the diameter of the collimator was 2.0mm, the diffraction line was measured as Ti alpha phase (302), the cross section of the length center portion of the profile was mirror polished, the residual stress distribution in the cross section was measured, and the residual stress value with the largest absolute value was obtained. The cross section refers to a section perpendicular to the longitudinal direction of the profile.

A specific example for measuring the maximum value of the residual stress is described below. For example, in the case of a T-section profile shown in fig. 6A, which is composed of a cross plate and a vertical plate that is in contact with the center of the cross plate, the residual stress is maximized at any position of 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 in the section of the profile can be obtained.

Preferably, the heat-corrected α+β titanium alloy material of the present embodiment has a small dimensional change when subjected to cutting processing. When the dimensional change during cutting is within ±2.0mm in cross section, it is determined that there is no problem in processing into a member or a product. The dimensional change during cutting is preferably within.+ -. 1.5mm, more preferably within.+ -. 1.0mm, and of course, the smaller the dimensional change, the better.

The dimensional change is due to residual stresses inside the profile. In order to change the dimension of the cross section in the cutting process within ±2.0mm, the maximum residual stress in the cross section is preferably +400MPa or less, more preferably +300MPa or less. It is also estimated that the residual stress, which is a factor of dimensional change during cutting, has a small absolute value, and therefore the dimensional change during cutting is also small.

The chemical composition of the titanium alloy of the present embodiment is not particularly limited as long as it is an α+β type titanium alloy as exemplified above, and may be, for example, the following compositions [1] and [2 ]. Hereinafter "%" means "% by mass".

[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％、 The balance: the content of the element (Fe, ni, cr, mn) in mass percent of Fe,% Ni,% Cr,% Mn is less than or equal to 1.4% and less than or equal to 2.9% of Fe+Ni+Cr+Mn.

Al：4.4～6.5％

Al is an alpha stabilizing element, and may be contained in order to increase the fraction of the alpha phase. Considering a balance of 0.2% yield strength, ductility, toughness, the Al content is preferably 4.4 to 6.5%.

Fe：0.5～2.9％

Fe is a β -stabilizing element, and when contained, has an effect of lowering the β -phase transition temperature. In addition, fe has the effect of increasing the yield strength by 0.2%. Considering the balance between 0.2% yield strength and segregation at solidification and elongation, the content of Fe is preferably 0.5 to 2.9%.

Si：0～0.50％、O：0～0.25％、C：0～0.080％、N：0～0.050％

Si, O, C, N is not necessarily contained, and the lower limit of the content is 0. Also, elements having the same effect as the α -stabilizing element have an effect of increasing the fraction of the α phase and increasing the 0.2% yield strength by containing these elements. Considering the balance with ductility, the content is preferably Si:0 to 0.50 percent, O:0 to 0.25 percent, C:0 to 0.080 percent, N:0 to 0.050 percent. However, when the content of these elements is set to 0%, refining costs increase. Therefore, the lower limit value of O, C, N may be set to 0.010% or 0.050%. The lower limit value of 0.001% of the measurement accuracy in the chemical analysis may be set as the lower limit value of Si.

Ni：0～0.15％、Cr：0～0.25％、Mn：0～0.25％

Ni, cr, mn are not necessarily contained, and the lower limit of the content is 0. These elements may also be contained because they function as Fe. When the content of Ni, cr, and Mn becomes high, intermetallic compounds (Ti ₂Ni、TiCr₂, tiMn) are generated as equilibrium phases, and fatigue strength and room temperature ductility become poor. Therefore, the content is preferably Ni:0 to 0.15 percent, cr:0 to 0.25 percent of Mn:0 to 0.25 percent. The lower limit values of Ni, cr, and Mn may be set to 0.001%, and 0.001% of the lower limit value of the measurement accuracy in chemical analysis, respectively.

Considering the balance of room temperature tensile strength and room temperature ductility, the total amount of Ni, cr, mn, fe is preferably 0.50% or more and 2.90% or less.

The balance: ti and less than 0.4% of total impurity

In the case of the chemical composition [1], the balance is Ti and impurities. Examples of the impurity elements include Cl, na, mg, zr, sn, cu, mo, nb, ta mixed in the titanium refining step and impurities mixed from scrap. Any impurities will, when the content increases, cause a decrease in toughness with the Ti forming compounds, with the result that the workability decreases. Further, if the total content of impurities is too large, ductility decreases, and thus workability deteriorates. Therefore, in order not to hinder the effect of the α+β type titanium alloy material of the present embodiment, the total amount of impurity elements is preferably controlled to 0.4% or less. The content of each impurity element is preferably 0.1% or less.

[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％, The balance: the content of the element (Fe, ni, cr, mn) in mass percent of Fe,% Ni,% Cr,% Mn is less than or equal to 1.4% and less than or equal to 1.4% of Fe+Ni+Cr+Mn is less than or equal to 2.3%.

Al：4.4～5.5％

Al is an alpha stabilizing element, and is an element contained to increase the fraction of the alpha phase. Considering a balance of 0.2% yield strength, ductility, toughness, the Al content is preferably 4.4 to 5.5%.

Fe：1.4～2.3％

Fe is a β -stabilizing element, and when contained, has an effect of lowering the β -phase transition temperature. In addition, fe has the effect of increasing the yield strength by 0.2%. Considering the balance between 0.2% yield strength and segregation at solidification and elongation, the content of Fe is preferably 1.4 to 2.3%.

Mo：1.5～5.5％

Mo is a β -stabilizing element, and can lower the β -transformation temperature of the titanium alloy in the same manner as Fe. Mo also improves 0.2% yield strength, ductility, and fatigue strength, and improves hot workability. In view of balance with solidification segregation, the Mo content is preferably 1.5 to 5.5%.

O：0～0.20％、C：0～0.080％、N：0～0.050％、Si：0～0.10％

O, C, N, si is not necessarily contained, and the lower limit of the content is 0. These elements also have the same effect as the α -stabilizing element, and by containing these elements, the effect of increasing the fraction of the α phase and increasing the 0.2% yield strength is exhibited. Considering the balance with ductility, the content is preferably O:0 to 0.20 percent, C:0 to 0.080 percent, N:0 to 0.050 percent, si:0 to 0.10 percent. However, when the O, C, N content is set to 0%, the refining cost increases. Therefore, the lower limit value of O, C, N may be set to 0.010% or 0.050%. The lower limit value of 0.001% of the measurement accuracy in the chemical analysis may be set as the lower limit value of Si.

Ni：0～0.15％、Cr：0～0.25％、Mn：0～0.25％

Ni, cr, mn are not necessarily contained, and the lower limit of the content is 0. These elements may also be contained because they function as Fe. When the content of Ni, cr, and Mn becomes high, intermetallic compounds (Ti ₂Ni、TiCr₂, tiMn) are generated as equilibrium phases, and fatigue strength and room temperature ductility become poor. Therefore, the content is preferably Ni:0 to 0.15 percent, cr:0 to 0.25 percent of Mn:0 to 0.25 percent. The lower limit values of Ni, cr, and Mn may be set to 0.001%, or 0.001% of the lower limit value of the measurement accuracy in chemical analysis, respectively.

Considering the balance of room temperature tensile strength and room temperature ductility, the total amount of Ni, cr, mn, fe is preferably 1.40% or more and 2.30% or less.

The balance: ti and less than 0.4% of total impurity

In the case of the chemical composition [2], the balance is Ti and impurities. Examples of the impurity elements include Cl, na, mg, zr, sn, cu, nb, ta mixed in the titanium refining step and impurities mixed from scrap. Any impurity forms a compound with Ti when the content increases, resulting in a decrease in toughness, and as a result, workability decreases. Further, if the total content of impurities is too large, ductility decreases, and thus workability deteriorates. Therefore, in order not to hinder the effect of the α+β type titanium alloy material of the present embodiment, the total amount of impurity elements is preferably controlled to 0.4% or less. The content of each impurity element is preferably 0.1% or less.

The cross-sectional shape of the α+β titanium alloy material according to the present embodiment is not particularly limited as long as it is a shape capable of measuring the torsion angle from one end to the other end of the material. In the case of the round bar, since the torsion angle cannot be measured, the round 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 of the present embodiment. Examples of the cross-sectional shape include L-shape, T-shape, H-shape, U-shape, pi-shape, plus-shape, minus-shape, and the like. The shape shown in fig. 9A to 9C may be used.

Fig. 9A is a cross-sectional view of a profile having a shape similar to a so-called H-profile. The H-shaped profile has a shape formed by integrating a 1 st flat plate, a 2 nd flat plate and a 3 rd flat plate, wherein the 1 st flat plate and the 2 nd flat plate extend in parallel, and the end face of the 3 rd flat plate is abutted against the surfaces of the 1 st flat plate and the 2 nd flat plate. On the other hand, the profile of fig. 9A has a shape in which a 1 st flat plate, a 2 nd flat plate, a 3 rd flat plate, and a 4 th flat plate are integrally formed, the 1 st flat plate, the 2 nd flat plate, and the 3 rd flat plate extend in parallel, the end face of the 3 rd flat plate abuts against the surfaces of the 1 st flat plate and the 2 nd flat plate, and the end face of the 4 th flat plate abuts against the surfaces of the 2 nd flat plate and the 3 rd flat plate. That is, the profile of fig. 9A has a shape in which two H-shaped profiles are laterally arranged. It also allows a shape in which 3 or more H-shaped profiles are arranged.

Fig. 9B is a cross-sectional view of a profile having a shape similar to a so-called T-profile. The T-shaped profile has a shape in which a1 st flat plate and a2 nd flat plate having an end face abutting against one surface of the flat plate are integrally formed. On the other hand, the profile of fig. 9B has a shape in which the 1 st flat plate, the 2 nd flat plate, and the 3 rd flat plate with the end face abutting against one surface of the 2 nd flat plate are integrally formed.

Fig. 9C is also a cross-sectional view of a profile having a shape similar to a so-called T-profile. The profile of fig. 9C has a shape in which the 1 st flat plate, the 2 nd flat plate, the 3 rd flat plate with the end face abutting against one surface of the 2 nd flat plate, and the 4 th flat plate with the end face abutting against the other surface of the 2 nd flat plate are integrally formed.

The above-described shape is merely an example of the cross-sectional shape of the α+β titanium alloy material according to the present embodiment. For example, the cross-sectional shape of the α+β titanium alloy material according to the present embodiment can be modified in various ways such as a cross-sectional shape including a protrusion having a tumor shape in a part of the cross-sectional shape exemplified above.

Next, a method for producing an α+β titanium alloy material having the above-described excellent shape will be described. Hereinafter, the temperature of the profile refers to the surface temperature of the profile measured by a radiation thermometer.

First, an α+β titanium alloy is hot worked by a manufacturing method such as extrusion, die forging, rolling, etc., to obtain a profile of a desired shape. The method for producing the profile is not particularly limited, and is preferably hot extrusion in view of production efficiency and the like.

But the profile must be warped and twisted. According to the usual straightening process, it is difficult to straighten warpage and torsion while ensuring mechanical properties of the profile. In particular, in the case where the method of producing the profile is hot extrusion, the magnitude of warpage and torsion becomes remarkable. In the method for producing an α+β titanium alloy material according to the present embodiment, this problem is solved by the correction processing described later.

The profile is then heated to maintain the two-phase region of α+β. Then, a strain of 0.1% to 8% is applied to the profile at least in the longitudinal direction at a correction temperature of not less than-400 ℃ but not more than-200 ℃. Further, a torque is applied to the profile so that the longitudinal torsion is within + -3.0 DEG. Thereby, the shape of the profile is corrected. Here, the β phase transition temperature is a temperature that becomes β single phase when heated. When the correction temperature is too high, the warp after cooling becomes large. On the other hand, when the correction temperature is too low, a large strain is required to correct torsion, and the void area ratio increases, and the tensile characteristics and fatigue characteristics may decrease.

In general, torsion and warpage exhibit different behavior with respect to correction temperature, showing a trade-off relationship. Specifically, the higher the correction temperature, the more likely warpage occurs, and the lower the correction temperature, the more likely torsion remains. The temperature difference in the profile, which occurs at the time of straightening at high temperature, causes a difference in the amount of heat shrinkage in cooling of the profile, thereby generating warpage. The higher the correction temperature, the greater the temperature difference in the profile and therefore the greater the warpage. On the other hand, torsion is caused by strain introduced during extrusion. The lower the correction temperature, the greater the elastic limit of the profile. Therefore, even with the same correction amount, the lower the correction temperature, the smaller the strain (plastic strain) of the introduced profile. For the above reasons, in general, when the correction temperature is lowered to eliminate warpage, it is difficult to correct torsion, and therefore the torsion amount becomes large, and when the correction temperature is raised to eliminate torsion, the warpage amount becomes large.

In order to obtain an excellent shape using an α+β type titanium alloy profile, it is necessary to achieve both the elimination of torsion and the elimination of warpage. Thereby, the shape is corrected by heating the profile to a two-phase region of α+β from 200 ℃ or lower of β phase transition temperature to-400 ℃ or higher, and then applying a strain of 0.1% or more and 8% or less to the profile at least in the longitudinal direction, and further applying a torque to make the longitudinal torsion within ±3.0°. Thus, it is preferable to prevent the occurrence of voids and to combine the elimination of torsion and the elimination of warpage. This is to avoid the need for extremely high stiffness of the orthotic device, which becomes very expensive. In the straightening step, both strain in the longitudinal direction and torsion in the longitudinal direction are preferably applied to the profile, but in the case where the profile has a cross-sectional shape in which torsion is less likely to occur, torsion may be omitted from the straightening step.

Then, the shape-corrected α+β titanium alloy material is cooled to 500 ℃ or lower while applying a predetermined stress thereto. From the viewpoint of suppressing void formation, compressive stress is preferable. However, it is considered that it is very easy to industrially apply a tensile stress in the actual situation of the correction device in consideration of the occurrence of buckling. Therefore, a tensile stress of 20% or less of the 0.2% yield strength at room temperature is preferably applied to the profile. Further, since the profile shrinks with a decrease in temperature, it is difficult to make the tensile stress applied to the profile constant during cooling, but if the tensile stress is 20% or less of the 0.2% yield strength, the variation in tensile stress does not become a problem.

In addition to the tensile stress, the mold is preferably cooled to 500 ℃ or less while applying a predetermined torque to the mold. The torque is preferably a value that maintains the longitudinal torsion within ±3.0°. By satisfying these conditions, the shape change (deterioration) and residual stress caused by cooling can be reduced. The profile is preferably cooled to 450 ℃, more preferably to 400 ℃.

The method for cooling the profile to 500 ℃ or lower is not particularly limited as long as the titanium alloy can be uniformly cooled. The cooling means may be natural cooling or accelerated cooling. Specifically, examples of the accelerated cooling include cooling in a gas atmosphere such as Ar, N, H, he, cooling in a liquid such as water or oil, and the like.

In order to uniformly cool the titanium alloy, the average cooling rate from the corrected temperature to 500 ℃ is preferably set to 10 ℃/s or less, 1 ℃/s or less, 0.5 ℃/s or less, or 0.1 ℃/s or less. In addition, when the stress applied as described above is a tensile stress, it is considered that compressive stress remains on the surface of the profile whose temperature is lower than that of the inside. The residual stress has a great influence on the processing of the component and the product, but is also generally preferably compressive residual stress in view of improving fatigue characteristics. That is, it is desirable to adjust the residual stress to some extent. Further, "average cooling rate from the corrected temperature to 500 ℃ means a value obtained by dividing the difference between the corrected temperature and 500 ℃ by the time required for the profile temperature to decrease from the corrected temperature to 500 ℃.

After the correction, the cooling is performed while applying stress and/or torque, whereby stress is concentrated on a portion having a poor shape such as a portion having a large warp and a portion having a large torsion. Further, local plastic deformation occurs in a portion to which a large stress acts. This can effectively and efficiently improve the shape, alleviate the residual stress, maintain the excellent shape due to thermal correction at high temperature, significantly suppress the deterioration of the shape such as warpage and torsion generated during cooling, and reduce the residual stress in the titanium alloy. Since the residual stress is effectively and efficiently reduced in the profile cooled while applying the stress and/or torque, the shape change in the case of cutting after cooling is also reduced.

The value of the torque applied during cooling can be set appropriately according to the sectional shape of the profile. For example, as described above, it is preferable to maintain the longitudinal torsion within ±3.0°. More preferably, the torque in the longitudinal direction is maintained within ±2.0°, and still more preferably, within ±1.0°. The torque during cooling may be constant, but may also vary within a prescribed range. For example, the torque may be increased as the temperature decreases.

As described above, the stress applied in cooling is preferably a tensile stress of 20% or less of the 0.2% yield strength of the profile at room temperature. The stress applied during cooling is more preferably 15% or less, and still more preferably 10% or less of the 0.2% yield strength of the profile at room temperature. The lower limit of the tensile stress is not specified, but is preferably 1% or more of the 0.2% yield strength of the profile at room temperature in terms of the shape and the residual compressive stress. By satisfying these conditions, the shape change (deterioration) and residual stress due to plastic deformation and cooling can be reduced. The tensile stress during cooling may be constant or may be varied within a predetermined range. For example, a tensile stress corresponding to the elastic limit at each temperature can be applied.

The atmosphere during the process from the start of the correction of the shape of the profile to the time when the profile is cooled to 500 ℃ or lower is not particularly limited, and may be, for example, atmospheric air. On the other hand, if necessary, it is permissible to set the partial environment or the whole environment of the profile to a different atmosphere in a predetermined temperature range or in a different temperature range without significantly deteriorating the surface properties. As an example, in the present manufacturing method, a gas such as argon or nitrogen can be applied to the atmosphere in the straightening step of the profile. Specifically, for example, the chamber is used to cover the entire apparatus to prevent oxidation, a sealing gas is blown to a part to prevent oxidation, and the like.

In addition, after the shape of the profile is corrected, the temperature of the profile may be maintained at the temperature at which the correction is performed. The profile temperature may be maintained in a predetermined temperature range of from the β -phase transition temperature to 400 ℃ or higher to the β -phase transition temperature to 200 ℃ or lower. In addition, the profile can be maintained at a predetermined temperature range, for example, 500 to 650 ℃ during cooling of the profile. Such temperature maintenance is performed to remove strain generated during correction and thermal strain generated during cooling, thereby further improving the characteristics of the profile of the present embodiment.

The standard of the round bar of Ti-6Al-4V, which is a general-purpose alpha+beta titanium alloy, that is, JIS H4650 and ASTM B348, includes a standard of 0.2% yield strength and elongation. Specifically, 0.2% yield strength is specified in the standard: 828MPa or more, elongation: more than 10% is preferred. But no provision is made for organization in these standards. Therefore, even if the structure of the profile is a needle-like structure, the lower limit value is set to be the same as that of an equiaxed structure excellent in balance of strength and ductility.

According to the method for producing an α+β titanium alloy material of the present embodiment, if the correction amount is 8% or less, the 0.2% yield strength is improved without decreasing the elongation, the strength/ductility balance of ASTM B348 is satisfied, and the fatigue strength is not decreased. This is considered to be because, when the titanium alloy is subjected to stress and is subjected to the straightening, dislocations are accumulated in the grain boundaries, no void is formed, and the grain boundaries are reinforced.

In the standard AMS2245B, there is a standard regarding torsion and warpage. Specifically, the following is specified: the torsion is preferably 1 ° or less in 1 foot and within ±3.0° in the whole length, and the warpage is preferably within ±0.65mm (warpage height (mm)/total length (m) = ±2.17) per 300 mm.

However, in a more complex and highly precise shape, a more strict predetermined value of warpage and torsion needs to be achieved. The torsion is preferably 0.5 ° or less in 1 foot, and within ±2.0° in the whole length, and the warpage is preferably within ±0.45mm (warpage height (mm)/total length (m) = ±1.50) per 300 mm. These values can be achieved by the titanium alloy profile of the present embodiment.

The size of the α+β titanium alloy material according to the present embodiment is not particularly limited. According to the manufacturing equipment commonly used at present, the thickness of the profile is generally 30-300 mm, and the total length is 5-20 m. The thickness of the profile here refers to the diameter of the smallest circle that can contain the cross section of the profile. For example, in the case where the cross section of the profile is triangular, the thickness of the profile refers to the diameter of a circle passing through all three vertices of the cross section of the profile, i.e., the circumscribed circle of the cross section of the profile. In principle, however, even larger profiles can be produced by using corresponding production facilities. On the other hand, profiles with a total length of less than 5m can also be produced by hot extrusion. Alternatively, an α+β type titanium alloy material having a total length of less than 5m may be produced by producing a material having a length of 5m or more by a production apparatus and then cutting the material. In any case, the performance of the titanium alloy material according to the present embodiment can be maintained.

Examples

The effect of improving warp and torsion after heat correction according to the present invention will be described below with reference to examples. The present invention is not limited to the following examples.

Example (example)

In order to confirm the influence of the tensile properties and the correction amount of the profile on the fatigue properties, a titanium alloy having the composition shown in table 1 and the β -phase transition temperature at the time of heating was used to produce a test material having a T-shape or a U-shape (fig. 3) in which a temperature difference was likely to occur in the cross section. The components shown in table 1 are values measured after the heat correction described later. But hot extrusion and hot straightening do not affect the composition change of the titanium alloy.

Titanium alloys having the composition of table 1 have a two-phase structure consisting of an alpha phase and a beta phase between room temperature and beta phase transition temperature. The test material is prepared by carrying out final hot forging on a solid ingot in an alpha+beta temperature region, heating the obtained phi 200 blank to a beta single-phase temperature region and extruding the blank to prepare the alpha+beta titanium alloy section.

The beta transus temperature of the titanium alloy was obtained by the following procedure.

First, the estimated value tβ of the β transformation temperature of the titanium alloy is calculated by substituting the content of the element corresponding to each symbol into the symbol of the element described in the following formula.

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 estimated value is not accurate, and may be different from the actual β -phase transition temperature by about 100 ℃. Therefore, the true β -phase transition temperature was obtained according to the following procedure.

Step (a) is performed by quenching the titanium alloy after heating to various temperatures within a temperature range of β transus temperature ± about 50 ℃. The heating temperature was varied by 5℃each time within the above range.

In the step (B), various tissues obtained in the above (A) were observed to confirm whether or not needle-like tissues were formed in them. The heating temperature of the titanium alloy confirmed to be needle-like structure exceeds the beta transus temperature.

Step (C) regards the minimum value of the heating temperatures at which the needle-like structure is obtained as the beta transus temperature of the titanium alloy.

If needle-like structures are obtained at all heating temperatures, the above steps (a) to (C) are repeated in a temperature range lower than the above temperature range. If needle-like structures cannot be obtained at all heating temperatures, the above steps (a) to (C) are repeated in a temperature range higher than the above temperature range.

TABLE 1

In the table, "no element is positively added.

For the obtained profile, thermal correction was performed under various correction conditions. Then, the profile was naturally cooled to a predetermined temperature in the atmosphere. Table 2 shows the correction temperature (see "correction temperature" row), the strain amount in the longitudinal direction applied to the profile at the correction temperature (see "strain in the longitudinal direction" row), the torsion in the longitudinal direction applied to the profile at the correction temperature (see "torsion angle before cooling" row), and the cooling stop temperature (see "removal temperature" row). The tensile stress applied to the profile during cooling of the profile (see "tensile stress during cooling" column), the holding temperature during cooling of the profile (see "holding temperature during cooling" column), and the average cooling rate of the profile from the corrected temperature to 500 ℃ (see "cooling rate" column) are also shown in table 2.

For example, in No.1, a tensile stress of 75MPa and a torque generating torsion of-0.6 DEG are applied to a test material at a correction temperature of 720 ℃. Then, the test material was naturally cooled to 400 ℃ in a state where tensile stress and torque were applied, and removed from the tension applying device. The average cooling rate of the profile from the correction temperature to 500 ℃ was set at 0.5 ℃/sec.

On the other hand, in No.17, no torque is applied before cooling. Therefore, in No.17, the column "torque with or without cooling" is described as "no". In addition, in No.17, "-" is described as a symbol indicating that no torque is applied in the column of "torsion angle at the time of cooling".

In example 18, cooling during the leveling process was discontinued at 600 ℃ and the temperature was maintained. In example 18, the average cooling rate when the surface temperature of the profile was lowered from 720 to 600 ℃ was set to 0.5 ℃/sec, the surface temperature of the profile was maintained at 600 ℃ for 10 minutes, and then the average cooling rate when the surface temperature of the profile was lowered from 600 to 500 ℃ was set to 0.5 ℃/sec.

< Fatigue test >)

In the measurement of fatigue strength, a round bar test piece having a diameter of 5.08mm and a parallel portion length of 15.24mm was produced from the circled portion in fig. 3 so as not to include a region having a depth of less than 0.5mm from the surface of the profile. Using the round bar test piece, a fatigue test was performed under conditions of a stress ratio r=σmin/σmax=0.1 (σmin, σmax are both tensile stresses) and room temperature, and the maximum value of the strength σmax, which was not broken for 1.0x10 ⁷ times, was set as the fatigue strength.

For each profile, 0.2% yield strength, elongation, void area ratio, maximum value of residual stress, maximum value of dimensional change after cutting, torsion after correction, and warp after correction were measured. The results are shown in Table 2.

Test pieces having a parallel portion length of 28mm and a diameter of 6.25mm were collected from a profile according to ASTM E8, and were subjected to a tensile test at a rate of 0.005/min (8.3X10 ^-5/s) when the strain was 2% or less and at a rate of 0.1/min (1.67X 10 ^-3/s) when the strain was 2% or more, whereby the test pieces were evaluated. The collection position of the test piece is the part marked with a circle in fig. 3.

The same results can be obtained when the test piece is produced so as not to include a region having a depth of less than 0.5mm from the surface, with respect to the fatigue strength, 0.2% yield strength and elongation.

The maximum values of void area ratio and residual stress were evaluated by the above-described methods.

The maximum value of the dimensional change after the cutting process was evaluated by measuring the dimensional change due to the local inclination of the cross section as shown in fig. 6A and 6B accompanying the cutting with a vernier caliper. Fig. 6A is a cross-sectional view of an example of "T" in the "shape" column, and fig. 6B is a cross-sectional view of an example of "U" in the "shape" column.

Corrected warpage and torsion were measured according to AMS 2245B.

In table 2, the underlined sections indicate that they are outside the scope of the present invention.

The tissue was also evaluated. The evaluation of the tissue was performed as follows. First, a specimen for tissue observation was collected from a cross section of 1/2 of the length of the obtained profile. The size of the specimen for tissue observation is not particularly limited, and a section itself may be used as the specimen for tissue observation. Here, a test piece having a square of 10mm was used. The acquisition position is set to the position of fig. 3. The observation surface was mirror polished and etched with fluoronitric acid to develop a structure. The tissue was then observed using an optical microscope. As a result of observation, in any of the examples, it was confirmed that the metallographic structure was a needle-like structure.

In the good position of FIG. 3 located in the region having a depth of 0.5mm or more from the surface of the profile, it was confirmed that the metallographic structure was needle-like. On the other hand, in the region having a depth of less than 0.5mm from the surface of the profile, i.e., in the surface layer region, a region where needle-like structures and equiaxed structures were mixed was observed. It is assumed that this is because the temperature is likely to be lowered in the surface layer region during the hot extrusion, and the processing strain is introduced in the temperature region of β -phase transition temperature or lower. In the region having a depth of 0.5mm or more from the surface of the profile, the metallographic structure is preferably mainly composed of needle-like structures, and the proportion of the needle-like structures in the metallographic structure is particularly preferably 100%.

For the above-mentioned corroded test piece, the original β crystal grain diameter was measured by a cutting method using an optical microscope with a magnification of 200 times, the size of the measurement region was 7mm square, 15 lines each in the longitudinal and transverse directions were recorded at equal intervals in the measurement region.

In any of the embodiments, the average primary β grain diameter is 500 μm or less. In the examples, there are also examples in which the average original β grain diameter is 60 μm.

TABLE 2

The symbol underline indicates that the present invention is not limited to the above-described examples.

TABLE 3

No.1 to 13 and No.18 to 20 are examples in which natural cooling is performed while applying torque and tensile stress for maintaining an angle during cooling after leveling at various alloy types and temperatures. Since the elongation at the time of correction, that is, the strain in the longitudinal direction is 8% or less, the alloy has a 0.2% yield strength of 830MPa or more, an elongation of 10% or more, and a fatigue strength of 450MPa or more, and has excellent balance of strength/ductility and fatigue characteristics. Further, since the profile is cooled while applying torque and tensile stress at the time of cooling after heat straightening, a profile excellent in shape can be obtained in which the maximum value of dimensional change after cutting is 2.5mm or less and warp and torsion are small after heat straightening even if the maximum residual stress exceeds +400 MPa. In addition, in No.13, since the strain in the longitudinal direction at the time of correction is 8% as the upper limit, although a void is slightly generated, good mechanical properties are ensured. In No.24, since a tensile stress exceeding 20% of the 0.2% yield strength at room temperature is applied to the profile, good mechanical properties are ensured although voids are slightly generated. In No.25, since forced cooling is performed by using mist in a state where the profile is held by the leveler, the cooling rate is high, and as a result, the maximum residual stress in the cross section increases. However, in the profile of No.25, warpage and torsion were also small, and the conditional yield strength, fatigue strength and ductility were in acceptable ranges.

No.14 to 15 are profiles taken from the leveler while maintaining the leveled length and cooling to a low temperature without applying tensile stress during cooling after heat leveling. Therefore, in No.14 having a heat straightening temperature of 720 ℃, the warpage amount (mm)/total length (m) after straightening was outside the range of.+ -. 2.17, and the maximum residual stress exceeded +400MPa, and the maximum value of dimensional change before and after cutting exceeded 2.0mm. On the other hand, in No.15 having a heat straightening temperature of 600℃and below the beta transus temperature of-400℃the torsion was not straightened at an elongation of 3%, and after cooling the torsion was also outside the range of.+ -. 3.0 ℃.

In No.16, since the removal temperature after heat straightening exceeds 500 ℃, thermal stress is generated during cooling, the maximum residual stress exceeds +400MPa, and the maximum value of warpage after cooling and dimensional change before and after cutting exceeds 2.0mm.

In No.17, torque to maintain torsion is not applied at the time of cooling. Therefore, the rebound in cooling is large, and the torsion of the cooled profile is out of the range of ±3.0°.

No.18 was maintained at 500 to 650℃during cooling, and No.19 was cooled at a rate of 0.1 ℃/s or less, so that the 0.2% yield strength was higher than that of other alloys having the same alloy composition, and exceeded 850MPa.

Since the strain in the longitudinal direction at the time of correction of No.21 to No.22 exceeds 8%, a plurality of voids are found in the interior of the profile, and the ductility and fatigue strength are lower than the lower limit depending on the type of alloy.

In No.23, correction in the two-phase region is not performed, and torsion is applied at the time of cooling, as a result, large torsion is generated after the correction processing.

According to the present invention, it was confirmed that for an α+β titanium alloy excellent in strength and elongation and fatigue characteristics: by correcting the shape, the shape is significantly improved, and stable production is possible.

Description of the reference numerals

1. A container; 2. a rod; 3. a squeeze pad; 4. a mold; 5. blank material; 6. a section bar; 11. extrusion direction.

Claims (12)

1. An alpha+beta titanium alloy section bar, wherein,

The alpha+beta titanium alloy section has a 0.2% yield strength of 830MPa or more, an elongation of 10% or more, and a fatigue strength of 450MPa or more,

The alpha + beta titanium alloy section has needle-shaped structures,

The area ratio of the voids is 1.0X10 ^-5% or less,

The torsion angle from one end to the other end is within + -3.0 DEG,

The warp height (mm)/total length (m) is within + -2.17.

2. The α+β titanium alloy section bar according to claim 1, wherein,

The 0.2% yield strength is above 850 MPa.

3. An alpha + beta titanium alloy profile according to claim 1 or 2, characterized in that,

The area ratio of the voids is 1.0X10 ^-6% or less.

4. The alpha + beta titanium alloy profile according to any one of claim 1 to 3, wherein,

The average original beta grain diameter is 500 μm or less.

5. The α+β titanium alloy section bar according to any one of claim 1 to 4, wherein,

The maximum residual stress in the cross section is +400MPa or less.

6. The α+β titanium alloy section bar according to any one of claim 1 to 5, wherein,

The alpha+beta titanium alloy section is an extrusion section.

7. The α+β titanium alloy profile according to any one of claims 1 to 6, comprising, in 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 to 0.25%, and

Mn：0～0.25％，

The balance of Ti and impurities,

The content of Fe, ni, cr, mn% Fe,% Ni,% Cr,% Mn expressed in mass% satisfies 0.5% to less than or equal to% Fe+% Ni+% Cr+% Mn less than or equal to 2.9%.

8. The α+β titanium alloy profile according to any one of claims 1 to 6, comprising, 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：0～0.10％、

Ni：0～0.15％、

Cr:0 to 0.25%, and

Mn：0～0.25％，

The balance of Ti and impurities,

The content of Fe, ni, cr, mn% Fe,% Ni,% Cr,% Mn expressed in mass% satisfies 1.4% to less than or equal to% Fe+% Ni+% Cr+% Mn less than or equal to 2.3%.

9. A method for producing the α+β titanium alloy material according to any one of claims 1 to 8, comprising:

a step of hot working an alpha+beta titanium alloy to obtain a profile;

Heating the profile to a correction temperature of a beta transus temperature of-400 ℃ or higher and a beta transus temperature of-200 ℃ or lower, applying a strain of 0.1% or higher and 8% or lower in the longitudinal direction at the correction temperature, and further applying a torque for making the torsion of the profile in the longitudinal direction within + -3.0%; and

And cooling to 500 ℃ or lower while applying a tensile stress and the torque to the profile.

10. The method for producing an α+β titanium alloy section according to claim 9, wherein,

The tensile stress applied in cooling of the profile is 20% or less of a 0.2% yield strength at room temperature.

11. The method for producing an α+β titanium alloy section according to claim 9 or 10, characterized in that,

The manufacturing method further comprises a step of maintaining the shape at 500-650 ℃ during cooling.

12. The method for producing an α+β titanium alloy material according to any one of claims 9 to 11, wherein,

The profile has an average cooling rate from the corrected temperature to 500 ℃ of 10 ℃/s or less.