JP7307315B2 - α+β type titanium alloy bar and its manufacturing method - Google Patents

Description

TECHNICAL FIELD The present invention relates to an α+β type titanium alloy bar and a method for producing the same.

Titanium alloys are used as lightweight and high-strength materials in fields such as aircraft, automobiles, and consumer products such as golf clubs. Titanium alloys that are commonly used are mainly composed of α phase and β phase, and Ti-6Al-4V, Ti-6Al-2Sn-4Zr-2Mo, Ti-5Al-1Fe alloys, etc. are known. there is

It is known that the α-phase of titanium, which has a close-packed hexagonal structure, undergoes creep deformation even at low temperatures such as room temperature when high stress is applied, and even titanium alloys containing the α-phase undergo creep deformation at room temperature. Furthermore, it is known that the tendency of creep deformation in titanium alloys containing the α phase leads to a decrease in life due to fatigue (dwell fatigue) in which a high load condition continues for a certain period of time, as typified by a trapezoidal wave type load cycle. there is (Non-Patent Documents 1-3)

Dwell fatigue is particularly problematic when used as aircraft jet engine components, as it leads to fracture in fewer cycles than triangular or sinusoidal duty cycles, which do not sustain high load conditions. can be

Patent Document 1 (Japanese Patent Application Laid-Open No. 2016-199796) discloses a titanium alloy bar having excellent fatigue properties and a method for producing the same. In Patent Document 1, the angle (inclination of the c-axis) between the c-axis direction of the hexagonal close-packed structure and the length direction of the titanium alloy bar in the proeutectoid α grains is 25° or more and 55° or less, and is equivalent to a circle. It is stated that the area ratio of proeutectoid α-grains having a diameter of 20 μm or more in the metal structure is 2.0% or less. This is described in paragraph 0020 of Patent Document 1, "The basal slip of a close-packed hexagonal crystal is more likely to occur as the crystal orientation (indicated by the symbol "θ" in FIG. 2) is closer to 45°, and the crystal orientation is 25°. It becomes active when it is more than 55 degrees or less. In addition, the larger the size of the equiaxed proeutectoid α-grains contained in the metal structure, the more easily the stress applied to the test piece concentrates. . Therefore, proeutectoid α-grains with a c-axis inclination of 25° or more and 55° or less and an equivalent circle diameter of 20 μm or more are likely to cause basal slip of a dense hexagonal crystal, and stress is likely to concentrate, so fatigue life is shortened. presumably shortened. It is based on the technical idea of "", and is appropriate as a normal fatigue fracture mechanism.

On the other hand, as described in Non-Patent Documents 1 to 3, different fracture mechanisms are known for Dwell fatigue. When (S) and an α-grain (H) with an inclination of about 0° are adjacent to each other, the stress is concentrated on the H-grain and a facet-like fracture surface perpendicular to the stress axis is generated. It is known from another study that it is almost parallel to the hexagonal basal plane.

Patent Document 1 does not mention Dwell fatigue.

Patent Document 2 (Japanese Patent Publication No. 2009-531546) discloses a technique for improving resistance to Dwell fatigue. Here, in the TA6Zr4DE (Ti-6Al-2Sn-4Zr-2Mo) alloy, the rupture life is increased from 5500 times to 10000 times by performing heat treatment for 4 to 8 hours at a temperature of -20 to -15 ° C. at the β transformation point. Improved. However, the process before heat treatment is only stamping in the β region, and the heat treatment process before that is unclear and cannot form a sufficiently fine microstructure, and the effect on the anisotropy of Dwell fatigue life is uncertain.

Patent Document 3 (Japanese Unexamined Patent Application Publication No. 2012-224935) discloses a titanium alloy billet in which the degree of accumulation of the α phase with respect to a specific direction of the c-axis is specified. However, the grain size of the α-phase, which is the starting point of fatigue fracture, is not mentioned, and fatigue characteristics cannot be improved simply by increasing the degree of accumulation.

Patent Document 4 (Japanese Patent Application Laid-Open No. 2014-65967) discloses a titanium alloy billet in which the degree of accumulation of the α phase in a specific direction of the c-axis is specified. However, since this patent document does not intend to improve the fatigue strength, the effect cannot be obtained, and the c-axis accumulation direction is different from that of the present invention.

JP 2016-199796 A Japanese translation of PCT publication No. 2009-531546 JP 2012-224935 A JP 2014-65967 A

M.R.Bache, “A review of dwell sensitive fatigue in titanium alloys: the role of microstructure, texture and operating conditions”, International Journal of Fatigue 25 (2003) 1079-1087 V.Sinha, M.J.Mills, J.C.Williams, “Determination of crystallographic orientation of dwell-fatigue fracture facets in Ti-6242 alloy”, J Mater Sci (2007) 42:8334-8341 Adam L. Pilchak, “Progress in Understanding the Fatigue Behavior of Ti Alloys”, Materials Science Forum Vol.710, pp85-92

The present invention has been made in view of the above circumstances, and an object of the present invention is to provide an α+β type titanium alloy bar having good Dwell fatigue characteristics and a method for producing the same.

Means for solving the above problems are as follows. In the present invention, the good Dwell fatigue property means that stress is applied in two directions, the longitudinal direction of the bar and the direction perpendicular to the diameter direction in the plane perpendicular to the longitudinal direction, that is, the circumferential direction. It means that the difference in Dwell fatigue life is small.

[1] The chemical components are Al: 5.50 to 6.75% by mass, V: 3.5 to 4.5% by mass, Fe: 0.05 to 0.40% by mass, O: 0.05 to 0 α+β type titanium alloy bar material containing .25% by mass and the balance being Ti and impurities, or the chemical components are Al: 5.50 to 6.50% by mass, Sn: 1.75 to 2.25% by mass %, Zr: 3.5 to 4.5% by mass, Mo: 1.8 to 2.2% by mass, Fe: 0.02 to 0.25% by mass, O: 0.02 to 0.15% by mass An α + β type titanium alloy bar material containing and the balance consisting of Ti and impurities ,
In a cross section parallel to the major axis direction of the α+β type titanium alloy bar, the area ratio of α crystal grain aggregates having an equivalent circle diameter of less than 100 μm and an aspect ratio of less than 3.0 is 60% or more,
The normal direction of the (0001) plane of the close-packed hexagonal crystal that constitutes the α crystal grains, the longitudinal direction of the α+β-type titanium alloy rod, and the circumferential direction of the α+β-type titanium alloy rod in the cross section perpendicular to the major axis. The α + β type titanium alloy bar, characterized in that the area ratio of α crystal grains having an equivalent circle diameter of more than 20 μm in the range of 0 ° to 55 ° and ω formed with each is 15% or less .
[ 2 ] A first step of heating a titanium alloy billet obtained by hot working an ingot to a temperature in the β single phase region and then quenching it;
a second step of heating the titanium alloy billet to a temperature in the α+β two-phase region, forging the titanium alloy billet over its entire length in a direction intersecting the longitudinal direction of the titanium alloy billet, and then cooling;
The titanium alloy billet is heated to a temperature in the α + β two-phase region and not higher than the heating temperature in the second step, and the titanium alloy billet is extended from a direction intersecting the major axis direction of the titanium alloy billet to the entire length of the billet. and a third step of forging over and at least finally cooling, in that order,
In the second step, in the cross section perpendicular to the longitudinal direction of the titanium alloy billet, the maximum width passing through the center of gravity of the perpendicular cross section is W1, the minimum width passing through the center of gravity of the perpendicular cross section is W2, and W1/W2 after forging. is 1.3 or less, and the ratio ΔW (ΔW= Wafter/Wini ) between the width Wini of the titanium alloy billet before forging and the width Wafter after forging is 1.05 or less; a step of performing a first rotation step of rotating the titanium alloy billet around the longitudinal axis and a first feeding step of moving the titanium alloy billet in the longitudinal direction, wherein the first forging step is performed at least twice; A step of forging the titanium alloy billet over its entire length by performing the above and performing the first feeding step after performing the first rotating step between the first forging steps,
The third step includes a second forging step of performing a reduction such that W1/W2 after forging is 1.5 or less, a second rotating step of rotating the titanium alloy billet around the long axis, and the titanium a step of performing a second feeding step of moving the alloy billet in the longitudinal direction, wherein the second forging step is performed at least twice, and after performing the second rotating step between the second forging steps a step of forging the titanium alloy billet over its entire length by performing the second feeding step;
The forging ratio in the second step is 1.6 or less, and the forging ratio in the third step is 1.5 or more and 2.5 or less,
The method for producing an α+β type titanium alloy bar according to [1], characterized in that:
[ 3 ] The production of the α+β type titanium alloy bar according to Claim 2 , wherein the first step is a step of heating the titanium alloy billet to a temperature in the β single phase region, working it, and then quenching it. Method.
[ 4 ] The chemical components of the titanium alloy billet are Al: 5.50 to 6.75% by mass, V: 3.5 to 4.5% by mass, Fe: 0.05 to 0.40% by mass, O: The method for producing an α+β type titanium alloy bar according to [ 2 ] or [ 3 ] containing 0.05 to 0.25% by mass, the balance being Ti and impurities.
[ 5 ] The chemical components of the titanium alloy billet are Al: 5.50 to 6.50% by mass, Sn: 1.75 to 2.25% by mass, Zr: 3.5 to 4.5% by mass, Mo: 1.8 to 2.2% by mass, Fe: 0.02 to 0.25% by mass, O: 0.02 to 0.15% by mass, the balance being Ti and impurities [ 2 ] or [ 3 ] .

According to the present invention, it is possible to provide an α+β type titanium alloy bar having good Dwell fatigue properties and a method for producing the same.

FIG. 4 is a diagram for explaining the crystal structure of the titanium alloy bar of the present embodiment, showing the orientation of the major axis direction of the titanium alloy bar and the normal direction of the (0001) plane of the close-packed hexagonal crystal that constitutes the α crystal grains. The figure explaining a difference. It is a diagram for explaining the crystal structure of the titanium alloy bar of the present embodiment, in which the circumferential direction T of the titanium alloy bar and α crystal grains are formed in a cross section perpendicular to the major axis direction of the titanium alloy bar. FIG. 4 is a diagram for explaining the orientation difference from the normal direction of the (0001) plane of the dense hexagonal crystal. (a) is a schematic diagram showing α crystal grains and α crystal grain aggregates in the titanium alloy bar material of the present embodiment, and (b) is a schematic diagram for explaining the aspect ratio of the α crystal grain aggregates. FIG. 4 is a schematic diagram for explaining the method for manufacturing the titanium alloy bar according to the present embodiment, and is a diagram for explaining the positional relationship between the titanium alloy billet and the anvil. FIG. 2 is a schematic diagram for explaining the method for manufacturing the titanium alloy bar according to the present embodiment, and is a diagram for explaining the cross-sectional shape of the titanium alloy billet.

It is known that the tensile properties of titanium alloys have anisotropy depending on the texture. When the normal direction of the (0001) plane is accumulated in the stress direction, the proof stress and tensile strength increase by 0.2%, but when the normal direction of the (10-10) plane is accumulated in the stress direction, it is 0.2%. % yield strength and tensile strength are lowered. The same is true for fatigue characteristics with a normal triangular wave or sine wave. For example, a graph (normalized SN ), if the microstructure is the same, it is represented on substantially the same line regardless of the direction of accumulation of the texture.
Note that "-1" in the description of "(10-10) plane" means that a line is drawn on "1".

However, it was found that the Dwell fatigue characteristics differed in behavior represented by a normalized SN diagram. That is, in the case of a texture in which the normal direction of the basal plane of a dense hexagonal crystal (hereinafter sometimes referred to as (0001) plane) is parallel to the stress axis, the life is significantly longer than that in a texture in which the direction is different. to In addition, when the direction of the stress axis is changed between the major axis direction of the titanium alloy bar and the cross section orthogonal to the major axis direction, the ratio of the (0001) plane orientation of the dense hexagonal crystal changes, so the Dwell fatigue life changes significantly.

Titanium alloy bars used as materials for aircraft engine parts should preferably have a homogeneous metallographic structure and material properties, and should have a small difference in Dwell fatigue life between the long axis direction and the direction perpendicular to the long axis. is preferred.

The anisotropy of Dwell fatigue life is considered to be due to the following mechanism. In Dwell fatigue, crack initiation is accelerated by strain accumulation, and crack propagation is accelerated by the formation of faceted fracture surfaces substantially parallel to the (0001) plane of the dense hexagonal crystal, resulting in a reduction in life. The Dwell fatigue life decreases because crack initiation is accelerated as the area ratio of the coarse α phase having a specific crystal orientation with respect to the stress axis direction increases. In particular, when coarse α-grain aggregates (α-colony) with a large aspect ratio are present, they act as substantially coarse α-grains when stress or strain acts in the long side direction, and cracks occur remarkably. Because it is premature, Dwell fatigue life is significantly reduced.

Therefore, it is advantageous for improving the Dwell fatigue life to have few coarse α crystal grain aggregates (α colonies) with a large aspect ratio and few coarse α crystal grains having a specific crystal orientation with respect to the stress axis direction. be.

In addition, in order to reduce the anisotropy of the Dwell fatigue properties in the long axis direction and the direction perpendicular to the long axis, the ratio of orientation of the (0001) plane of coarse α crystal grains (close-packed hexagonal crystal) in each direction should be small. is advantageous. Here, the orientation of the (0001) plane of the close-packed hexagonal crystal can be represented by the orientation of the bottom surface of the close-packed hexagonal crystal in the long axis direction and in the cross section orthogonal to the long axis. Therefore, the angles θ and ω formed between the normal direction of the (0001) plane of the dense hexagonal crystal and the major axis direction of the titanium alloy bar and the circumferential direction T in the cross section perpendicular to the major axis direction are within 55°. It is preferable that the area ratio of α crystal grains having a certain equivalent circle diameter exceeding 20 μm is small.

In order to control the aspect ratio and existence ratio of the α crystal grain aggregates and the crystal orientation of the α crystal grains as described above, it is important to understand the metal structure change behavior during hot working of the titanium alloy. Generally, in the forging process of a titanium alloy, a step of heating to the β single-phase region to reduce the crystal orientation deviation of the α-phase that previously existed and randomize it is incorporated. However, by processing in the α+β region after that, a new texture of the α phase is formed. In particular, it is difficult to eliminate the α-phase texture formed by the initial processing in the α+β region after cooling from the β single-phase region by subsequent processing in the α+β region. Therefore, it is necessary to control the processing method in the first α+β region after cooling from the β single phase region.

In the present invention, in the processing in the α + β region subsequent to the step of cooling from the β single phase region, when the maximum width passing through the center of gravity of the orthogonal cross section of the material is W1 and the minimum width passing through the center of gravity of the orthogonal cross section is W2, forging A titanium alloy billet is reduced such that the post-forging W1/W2 is 1.3 or less and the ratio ΔW (ΔW=Wa/Wi) between the width Wi of the titanium alloy billet before forging and the width Wa after forging is 1.05 or less. is rotated around the long axis, and then the same forging is performed over the entire length of the titanium alloy billet. can reduce the anisotropy of Dwell fatigue life. In addition, the step of reducing the titanium alloy billet so that W1/W2 after forging is 1.5 or less while rotating the titanium alloy billet around the long axis, and then performing the same forging over the entire length of the titanium alloy billet is performed at a forging ratio of 1. 0.5 or more and 2.5 or less. This makes the titanium alloy bar material suitable for use in aircraft engine parts.

The titanium alloy bar of this embodiment will be described below.
The titanium alloy bar material of the present embodiment preferably has a metallographic structure in which the main phase is the α phase and the β phase is the secondary phase at 25°C. That is, it is preferably a component of an α+β two-phase titanium alloy, and preferably contains 5.50 to 6.75% by mass of Al. In addition to Ti and Al, 3.5 to 4.5% by mass of V, 0.05 to 0.40% by mass of Fe, and 0.05 to 0.25% by mass of O may be contained. When the Al content is 5.50% by mass or more, a titanium alloy bar having high strength and excellent fatigue properties can be obtained. Moreover, when the Al content is 6.75% by mass or less, it is possible to prevent the titanium alloy bar from becoming brittle due to the formation of intermetallic compounds such as Ti ₃ Al.

The titanium alloy bar material of the present embodiment may be formed of, for example, components specified by AMS4928. That is, Al: 5.50 to 6.75% by mass, V: 3.5 to 4.5% by mass, Fe: 0.05 to 0.40% by mass, O: 0.05 to 0.25% by mass may be contained, and the balance may be Ti and impurities. As impurities, for example, N: 0.08% by mass or less, C: 0.08% by mass or less, and H: 0.015% by mass or less may be contained.

Further, the titanium alloy bar material of the present embodiment may be formed of, for example, components specified by AMS4975. That is, Al: 5.5 to 6.5% by mass, Sn: 1.75 to 2.25% by mass, Zr: 3.5 to 4.5% by mass, Mo: 1.8 to 2.2% by mass, It may contain 0.02 to 0.25% by mass of Fe, 0.02 to 0.15% by mass of O, and the balance may be Ti and impurities. As impurities, for example, N: 0.08% by mass or less, C: 0.08% by mass or less, and H: 0.015% by mass or less may be contained.

The titanium alloy bar of the present embodiment is a round bar having a circular cross-sectional shape perpendicular to the major axis direction. It is sufficient if the shape is approximately circular. Although the diameter is not particularly limited, it preferably ranges from 100 mm to 350 mm.

On the other hand, regarding the shape of the intermediate shape from the ingot to the rod material, the cross-sectional shape orthogonal to the long axis direction is not limited to a circular shape, and may be a polygon such as a quadrangle or octagon, or a polygon with rounded corners. It may be rectangular.

Next, the crystal structure of the titanium alloy bar material of this embodiment will be described with reference to FIGS. 1 to 3. FIG.
FIG. 1 shows the orientation of the (0001) plane of the hexagonal close-packed crystal in the long axis direction of the titanium alloy bar and the α crystal grains. Let θ be the angle between the major axis direction of the titanium alloy rod and the normal direction of the (0001) plane of the hexagonal close-packed crystal. In the titanium alloy bar of the present embodiment, in a cross section parallel to the major axis direction of the titanium alloy bar, the area ratio of α crystal grains having an equivalent circle diameter of more than 20 μm with θ in the range of 0° or more and 55° or less is It is preferably 15% or less. If the area ratio exceeds 15%, the Dwell fatigue life is lowered, which is not preferable.

FIG. 2 shows the orientation state of close-packed hexagonal crystals in α crystal grains in a region of minute length dL in the major axis direction of a titanium alloy bar. FIG. 2(a) is a perspective view of the titanium alloy bar seen from the peripheral surface side, and FIG. 2(b) is a perspective view of the titanium alloy bar seen from the cross section perpendicular to the longitudinal direction. T is the circumferential direction in the orthogonal section. The circumferential direction is a direction orthogonal to the radial direction in the orthogonal cross section. Let ω be the angle between the normal direction of the (0001) plane of the hexagonal close-packed crystal and the circumferential direction T. In the titanium alloy bar of the present embodiment, the area ratio of α crystal grains with an equivalent circle diameter of more than 20 μm in the range of ω in the range of 0° to 55° is 15 in the cross section perpendicular to the major axis direction of the titanium alloy bar. % or less. If the area ratio exceeds 15%, the Dwell fatigue life is lowered, which is not preferable.

FIG. 3(a) shows α crystal grains and α crystal grain aggregates (α colonies). Crystal grains are schematically indicated by circles. The arrows in the circled α crystal grains schematically indicate the crystal orientation. An aggregate of α crystal grains having similar crystal orientations is called an α crystal grain aggregate (α colony). In FIG. 3(a), α colonies are schematically indicated by ellipses. FIG. 3(b) explains the aspect ratio of the α crystal grain aggregate (α colony), and the aspect ratio is defined as the ratio of major axis (a1)/minor axis (a2).

In the titanium alloy bar of the present embodiment, in a cross section parallel to the major axis direction, the area ratio of α crystal grain aggregates having an equivalent circle diameter of less than 100 μm and an aspect ratio of less than 3.0 is preferably 60% or more. . By satisfying this condition, the Dwell fatigue life can be improved.

The crystal structure of the titanium alloy bar material of this embodiment can be measured using EBSD (Electron Backscatter Diffraction).

First, a test piece having a cross section in the longitudinal direction as an observation surface is taken from the central part in the longitudinal direction of the titanium alloy bar. The measurement point on the observation plane is a position at a depth of r/2 from the surface of a circular bar whose cross section has a radius of r. Next, a rectangular area of 3 mm in length and 3 mm in width at the measurement point on the observation surface of the test piece is set as a field of view, and measurement is performed using EBSD at a measurement interval of 2.0 μm and an acceleration voltage of 15 kV.

The obtained measurement results are analyzed using OIM (crystal orientation analysis software manufactured by TSL Solutions Co., Ltd.). First, a Partiton for only the α phase is created and analyzed.
Next, α crystal grains are determined by setting the angle difference (misorientation angle) between the orientations (c-axis direction) of adjacent EBSD measurement points to 5° or less, and the area of each α crystal grain is calculated from the number of measurement points of the α crystal grains. , and the circle-equivalent diameter of each α crystal grain is calculated.

In addition, the average value of the c-axis direction at the EBSD measurement points in each α crystal grain is calculated, and using it, the normal direction of the (0001) plane of the α crystal grain and the titanium alloy bar Calculate the angle θ formed with the major axis direction of .

Then, among the α crystal grains, the area ratio of the α crystal grains with an angle θ of 0° to 55° is obtained. Also, the area ratio of α-crystal grains having an equivalent circle diameter of more than 20 μm is determined.

Alternatively, after setting the grain size to more than 20 μm in Partition Properties, use the Crystal Direction Map to determine that the angle θ formed by the normal direction of the (0001) plane of the α crystal grain and the major axis direction of the titanium alloy bar is 0°. Determine the area ratio (Total Fraction) of α crystal grains from 55°.

Furthermore, after setting the grain size to more than 20 μm in the partition properties, a crystal direction map is created and the angle between the normal direction of the (0001) plane of the α crystal grain and the circumferential direction of the titanium alloy bar is 0° or more. Determine the area ratio (Total Fraction) of α crystal grains in the range of 55° or less.

Regarding the α crystal grain aggregate (α colony), the α crystal grain aggregate (α colony) is determined by setting the angle difference (misorientation angle) between the crystal orientations of adjacent EBSD measurement points to 15° or less, and the α crystal grain The area of each α crystal grain aggregate (α colony) is obtained from the number of measurement points of the aggregate (α colony), and the equivalent circle diameter of the α crystal grain aggregate (α colony) is calculated. The aspect ratio is determined by approximating the region of each α crystal grain aggregate (α colony) to an elliptical shape. Specifically, after setting the Grain Size to less than 100 μm and the Aspect Ratio to 0.334 or more in the Partition Properties, a Phase Map is created to determine the area ratio (Total Fraction) of α crystal grain aggregates (α colonies). Here, the Aspect Ratio of 0.334 is the reciprocal of 2.994, which corresponds to an aspect ratio of less than 3.0.

Next, a method for manufacturing a titanium alloy bar according to this embodiment will be described.
The titanium alloy bar material of the present embodiment is produced by melting a raw material adjusted to a predetermined chemical composition to obtain an ingot, and then heating and processing the obtained ingot to a β single phase region. It is obtained by subjecting a titanium alloy billet obtained through α+β forging, which is heated to a two-phase region and processed, to the following steps.

The titanium alloy bar material of the present embodiment is produced by a first step of heating the above titanium alloy billet having a predetermined chemical composition to a temperature in the β single phase region and then quenching it, and then heating the titanium alloy billet to a temperature in the α + β two phase region. A second step of heating to, forging and then cooling, and a third step of heating and forging the titanium alloy billet to a temperature in the α + β two-phase region that is equal to or lower than the heating temperature in the second step. , in this order.
Each step will be described below.

In the first step, the titanium alloy billet is heated to the β single-phase temperature range in a heating furnace and then quenched to homogenize the metal structure and suppress coarsening of crystal grains. Heating in the β single-phase temperature region is performed by setting the temperature in the heating furnace to a temperature higher than the β transformation point temperature by 30°C or higher and a temperature higher than the β transformation point temperature by 100°C or lower (β transformation point temperature +30°C to β transformation point temperature +100°C). temperature range). When the temperature in the heating furnace is 30° C. higher than the β transformation temperature, even if there are portions where the temperature is uneven in the heating furnace or the size of the titanium alloy billet is large, This is preferable because the entire billet is heated to a temperature equal to or higher than the β transformation temperature. Further, when the temperature in the heating furnace is 100° C. or lower than the β transformation point temperature, oxidation of the surface layer of the titanium alloy billet is suppressed, and coarsening of the metal structure in the titanium alloy billet is suppressed. Therefore, a high-quality titanium alloy bar can be obtained.

In the first step, after heating to the β-single-phase temperature range, the titanium alloy billet is preferably taken out from the heating furnace and rapidly cooled, or it is preferably processed and then rapidly cooled. Rapid cooling is generally performed by immersing the titanium alloy billet in a sufficient amount of water in order to obtain a sufficient cooling rate. good. Rapid cooling is preferably continued until the surface temperature of the titanium alloy billet reaches 300° C. or lower.

In the first step, strain may be imparted to the titanium alloy billet by heating it to the β single-phase temperature range, removing it from the heating furnace, and then working it. Recrystallization is caused by applying strain, and coarsening of crystal grains in the metal structure is suppressed.

Next, in the second step, the titanium alloy billet after the first step is heated to a temperature in the α+β two-phase region, and the titanium alloy billet is heated over its entire length in a direction intersecting the longitudinal direction of the titanium alloy billet. forging and then cooling. In the second step, a titanium alloy billet, which is a material to be worked, is worked in a two-phase state of α phase and β phase. In particular, it is preferable to work in a temperature range where the β phase exists in the structure at a rate of about 50%.

In the second step, the temperature in the heating furnace for heating the titanium alloy billet is 60° C. lower than the β transformation point temperature or higher and lower than the β transformation point temperature (β transformation point temperature −60° C. to less than the β transformation point temperature. temperature range). Considering the temperature rise due to the heat generated during processing, the upper limit of the heating temperature is preferably less than 20° C. lower than the β transformation temperature (β transformation temperature −20° C.).

When the temperature in the heating furnace is at least 60°C lower than the β transformation point temperature, the deformation resistance of the titanium alloy billet during hot working can be prevented from becoming too large, and hot working can be performed easily and efficiently. It can be performed. In addition, when the temperature in the heating furnace is lower than the β transformation point temperature, α crystal grains are sufficiently precipitated in the metal structure of the titanium alloy billet, so grain growth is suppressed, and in the α + β two-phase temperature range, A sufficient effect can be obtained by applying hot working.

Since the surface temperature of the titanium alloy billet gradually decreases during forging, if the surface quality deteriorates or surface cracks are likely to occur, the forging is temporarily interrupted before the end of the second step, and then forged again. , It is preferable to forge after heating the titanium alloy billet.

The second step will be further described with reference to FIG. FIG. 4 is a diagram showing a titanium alloy billet and an anvil, FIG. 4(a) is a side view before rolling, FIG. 4(b) is a plan view before rolling, and FIG. It is a side view after the reduction, and FIG. 4(d) is a plan view after the reduction. FIG. 4(e) is a side view after rotating the titanium alloy billet around the axis, and FIG. 4(f) is a plan view after rotating the titanium alloy billet around the axis. In addition, in FIG. 4, the code|symbol 1 is an anvil, and the code|symbol 2 is a billet.

In the second step, a pair of anvils is applied from a direction substantially perpendicular to the longitudinal direction of the billet. The second step promotes the division and refinement of α crystal grain aggregates (α colonies) in the titanium alloy, and the longitudinal direction of the bar and the circumferential direction of the bar in the cross section perpendicular to the major axis. suppresses an increase in the anisotropy of the (0001) plane orientation with respect to each of

More specifically, the second step includes a first forging step of forging a titanium alloy billet with an anvil under predetermined conditions, a first rotating step of rotating the titanium alloy billet around its major axis, and a longitudinal rotation of the titanium alloy billet. and a first feeding step of moving in the axial direction. In this case, the first forging process is performed at least twice, the first rotating process is performed between the first forging processes, and then the first feeding process is performed. Thus, the titanium alloy billet is forged over its entire length.

Specifically, as shown in FIGS. 4(a) and 4(b) to 4(c) and 4(d), a first forging step is performed to partially form the outer peripheral surface of the titanium alloy billet. After pressing down a part to be processed with an anvil, as shown in FIGS. 4(c) and 4(d) to FIGS. Rotate around the long axis. Then, the titanium alloy billet is rolled down from a new direction. After the reduction in the first forging step is performed two or more times in this manner, the first feeding step is performed to relatively move the titanium alloy billet in the longitudinal direction by a predetermined feeding amount. When the first feeding step is performed, the anvil will face a new portion to be processed. The first forging process, the first rotating process, and the first feeding process are performed again on this new part to be machined. This operation is performed sequentially from one longitudinal end of the titanium alloy billet to the other end, and if necessary, gripping is changed, and the entire titanium alloy billet is forged. The rotation of the titanium alloy billet around the long axis in the first rotation step may be, for example, every 90° for a rectangular cross section and every 45° for an octagonal cross section.

The upper limit of the number of times of the first forging process is restricted by the forging ratio before and after the second process. Forging and stretching are repeated within the range of the forging ratio before and after the second step being 1.6 or less.

In the first forging step, as shown in FIG. 5, in the cross section perpendicular to the longitudinal direction of the titanium alloy billet, the maximum width passing through the center of gravity of the cross section is W1, and the minimum width passing through the center of gravity of the cross section is W2. The width Wini of the titanium alloy billet before forging (hereinafter sometimes referred to as Wi) and the width Wafter after forging (hereinafter referred to as Wa It is preferable to perform forging so that the ratio ΔW (ΔW=Wa/Wi) to 1.05 or less.
The width Wini of the titanium alloy billet before forging is the maximum projected width of the titanium alloy billet when viewed from the pressing direction of the anvil, as shown in FIG. 4(b). Further, the width Wafter of the titanium alloy billet after forging is the maximum projected width of the titanium alloy billet after completion of reduction.
FIG. 5(a) is a diagram showing W1 and W2 when the orthogonal cross-section shape is substantially rectangular, FIG. 5(b) is a diagram when the orthogonal cross-section is substantially circular, and FIG. This is the case where the cross section is substantially hexagonal.

All of these conditions are for making the degree of accumulation of the crystal orientation of the α crystal grains in the titanium alloy closer to the long axis direction of the bar and the circumferential direction of the bar in the plane perpendicular to the long axis direction. It is a condition. After the titanium alloy billet is rolled down, the titanium alloy billet is rotated and rolled down again to increase the contact area between the anvil and the surface of the titanium alloy billet in the longitudinal direction. Elongation is suppressed, and the accumulation of α crystal grains in the direction normal to the (0001) plane in the long axis direction increases. On the other hand, in the orthogonal section, the accumulation of α crystal grains in the direction normal to the (0001) plane increases in the circumferential direction perpendicular to the rolling direction. In addition, increasing the contact area between the anvil and the surface of the titanium alloy billet increases the accumulation of strain inside the titanium alloy billet, promoting the division of α crystal grains and α crystal grain aggregates (α colonies). Furthermore, by defining the upper limit of the forging ratio, it is possible to suppress an increase in the aspect ratio of the α crystal grain aggregate (α colony). By these actions, the degree of accumulation of coarse α crystal grains in the (0001) plane normal direction with respect to the long axis direction and the circumferential direction can be brought closer, and the aspect ratio of α crystal grain aggregates (α colonies) can be reduced. can. That is, W1/W2 after forging is 1.3 or less, and the ratio ΔW (ΔW=Wa/Wi) between the width Wini of the titanium alloy billet before forging and the width Wafter after forging is 1.05. A titanium alloy billet is forged so that the forging ratio is 1.6 or less. If the forging ratio exceeds 1.6, the anisotropy of the crystal orientation of the α crystal grains increases.

In order to efficiently change the degree of integration or the area ratio, the titanium alloy billet in which the crystal orientation is randomized by cooling after the β heat treatment in the first step is processed first. Control is important.

Next, in the third step, the titanium alloy billet is heated to a temperature in the α+β two-phase region that is equal to or lower than the heating temperature in the second step, and the titanium alloy billet is separated from the longitudinal direction of the titanium alloy billet. It is forged along its entire length and cooled at least at the end. In the third step, the second forging step is performed so that W1/W2 after forging is 1.5 or less. When performing the second forging step, the titanium alloy billet is rotated around its longitudinal axis to change the forging direction each time. This reduces the anisotropy of the (0001) plane orientation of the α crystal grains.

In the third step, the temperature in the heating furnace for heating the titanium alloy billet is preferably at least 80° C. lower than the β transformation point temperature and at most the heating temperature in the second step. Considering the temperature rise due to the heat generated during processing, the upper limit of the heating temperature is preferably less than 20° C. lower than the β transformation temperature (β transformation temperature −20° C.).

When the temperature in the heating furnace is at least 80°C lower than the β transformation point temperature, the deformation resistance of the titanium alloy billet during hot working can be prevented from becoming too large, and hot working can be performed easily and efficiently. It can be performed. Further, if the temperature in the heating furnace is equal to or higher than the temperature in the second step, the degree of integration of the (0001) plane orientation is lowered, which is not preferable.

Also in the third step, the temperature of the titanium alloy billet gradually decreases during forging, so if the surface quality is deteriorated or surface cracks are likely to occur, forging should be performed before the end of the third step. is suspended once, and the titanium alloy billet is heated again before forging.

The third step includes a second forging step of forging the titanium alloy billet with an anvil under predetermined conditions, a second rotating step of rotating the titanium alloy billet around its long axis, and moving the titanium alloy billet in its long axis direction. and a second feeding step. In this case, the second forging process is performed at least twice, the second rotating process is performed between the second forging processes, and then the second feeding process is performed. Thus, the titanium alloy billet is forged over its entire length.

More specifically, after the second forging step is performed to press down a part of the outer peripheral surface of the titanium alloy billet to be machined with an anvil, the second rotating step is performed to rotate the titanium alloy billet around its long axis. Let Then, the titanium alloy billet is rolled down from a new direction. After the reduction in the second forging step is performed two or more times in this manner, the second feeding step is performed to relatively move the titanium alloy billet in the longitudinal direction by a predetermined amount. When the second feeding step is performed, the anvil faces a new portion to be processed. The second forging process, the second rotating process, and the second feeding process are again performed on this new part to be machined. This operation is performed sequentially from one longitudinal end of the titanium alloy billet to the other end, and if necessary, gripping is changed, and the entire titanium alloy billet is forged. The rotation of the titanium alloy billet about its major axis in the second rotation step may be, for example, every 90° for a rectangular cross section and every 45° for an octagonal cross section.

In the third step, the second forging step is repeated until the forging ratio reaches 1.5 to 2.5 so that W1/W2 after forging is 1.5 or less. If the forging ratio is less than 1.5, the size of the α crystal grains cannot be refined, and the fatigue life deteriorates. On the other hand, if the forging ratio exceeds 2.5, the α crystal grain aggregates are elongated in the long axis direction of the billet, the aspect ratio is increased, and the fatigue life is deteriorated, which is not preferable.

As described above, according to the titanium alloy bar of the present embodiment, the area ratio of the α crystal grain aggregates having an equivalent circle diameter of less than 100 μm and an aspect ratio of less than 3.0 is 60% or more, and the α crystal grains The angles θ and ω formed between the normal direction of the (0001) plane of the close-packed hexagonal crystal that constitutes Since the area ratio of α crystal grains with an equivalent circle diameter of more than 20 μm in the range of 0° or more and 55° or less is 15% or less, the anisotropy of the Dwell fatigue characteristics is reduced, and the quality of the titanium alloy bar is improved. can be made

Further, according to the present embodiment, by sequentially performing the first step, the second step and the third step, it is possible to manufacture a titanium alloy bar having excellent Dwell fatigue characteristics.

The titanium alloy bar of the present embodiment can be suitably used as a material for turbine blades and turbine disks of aircraft engines, for example. That is, by further working the titanium alloy bar of the present embodiment to form a turbine blade or turbine disk, a turbine blade or turbine disk having excellent Dwell fatigue characteristics can be obtained.

Next, examples of the present invention will be described.
Titanium alloy bars were produced and evaluated by the method described below.

(pre-process)
A cylindrical ingot having a diameter of about 750 mm and having the composition shown in Table 1 obtained by melting is heated to the β single-phase temperature range in a heating furnace heated to 1020° C. or higher and 1200° C. or lower, which is higher than the β transformation temperature. After that, β forging, which is forged after being removed from the heating furnace, and α + β forging, which is removed from the heating furnace and forged after being heated to a two-phase region of α + β from 900 ° C to 980 ° C below the β transformation temperature, are performed once. Alternatively, the process was repeated several times to obtain a rod-shaped billet having a cross-sectional shape shown in Table 1 perpendicular to the longitudinal direction. The rod-shaped billet was used as an intermediate billet (titanium alloy billet). The β transformation temperature of the titanium alloy billet shown in Table 1 was in the range of 990°C to 1010°C.

In the column of the shape of the intermediate billet in Table 1, "ψ510" and "ψ380" mean that the cross-sectional shape is circular with a diameter of 510 mm and 380 mm, respectively, and "480 octagon" and "600 octagon" mean that the cross-sectional shape is It is octagonal and means an octagonal shape in which the distance between two opposing parallel sides is 480 mm or 600 mm, and "260 * 260" means that the cross-sectional shape is a square with a side length of 260 mm, "610*320" means that the cross-sectional shape is a square with a length of 610 mm and a width of 320 mm, and "600*300" means that the cross-section is a square with a length of 600 mm and a width of 300 mm.

(First step)
After heating the intermediate billet obtained in the preliminary step in a heating furnace at a heating temperature shown in Table 2, it is taken out of the heating furnace and subjected to forging (processing) followed by water cooling or forging ( It was water-cooled without processing). Water cooling was performed by immersion in a water tank containing a sufficient amount of water. Further, water cooling was performed until the ingot surface temperature became at least less than 300°C.

(Second step)
After heating the titanium alloy billet obtained in the first step in a heating furnace at the heating temperature shown in Table 2, it was removed from the heating furnace and forged. For the forging, the first forging process with W1/W2 and ΔW shown in Table 2 was performed 2 to 6 times. Table 2 shows the number of times of the first manufacturing process. By performing the first rotation process between each first forging process, the rolling direction of the first forging process was changed each time. Then, the 1st sending process was performed. By repeating the first forging process, the first rotating process, and the first feeding process a plurality of times, the titanium alloy billet was forged over its entire length. In addition, when the temperature dropped during forging, it was reheated in a heating furnace. Thus, a billet having a circular or polygonal cross-sectional shape was obtained. The heating temperature in the second step was in the range from the β-transformation temperature −60° C. to less than the β-transformation temperature (the temperature in the α+β two-phase region) for all samples.

(Third step)
After the titanium alloy billet after the second step was heated in a heating furnace at the heating temperature shown in Table 2, it was forged under the conditions shown in Table 2. For forging, the second forging step with W1/W2 shown in Table 2 was performed 2 to 6 times. Table 2 shows the number of times of the second manufacturing process. By performing the second rotation process between each second forging process, the rolling direction of the second forging process was changed for each time. After that, the second feeding process was performed. By repeating such a second forging process, a second rotating process and a second feeding process a plurality of times, the titanium alloy billet was forged over its entire length. In addition, when the temperature dropped during forging, it was reheated in a heating furnace. Thus, a titanium alloy bar having a circular cross section was produced. The heating temperature in the third step is No. All of the samples other than No. 10 had a temperature in the α+β two-phase region and a temperature lower than the heating temperature in the second step. The heating temperature of No. 10 exceeded the heating temperature of the second step.

After the second step and the third step, the ingot was air-cooled (allowed to cool) until the surface temperature of the ingot fell below 300°C.

The crystal structure of the obtained titanium alloy bar was measured.
First, a test piece having a cross section parallel to the longitudinal direction as an observation surface was taken from the longitudinal center of a titanium alloy bar. The measurement point on the observation plane was a position at a depth of r/2 from the surface for a circular sample with a cross section of radius r. Next, measurement was performed using EBSD at a measurement interval of 2.0 μm, an acceleration voltage of 15 kV, and a rectangular region of 3 mm in length and 3 mm in width at the measurement point on the observation surface of the test piece.
The obtained measurement results were analyzed using OIM (crystal orientation analysis software manufactured by TSL Solutions Co., Ltd.). First, a Partiton targeting only the α phase was created and used as an analysis target. α crystal grains were determined by setting the angular difference (misorientation angle) between the orientations (c-axis direction) of adjacent EBSD measurement points to 5° or less.

For the anisotropy in the orthogonal cross section, a Crystal Direction Map is created, and the orientation difference between the radial direction and the circumferential direction in the orthogonal cross section and the normal direction of the (0001) plane of the α crystal grain is within 25 °. The area ratio (Total Fraction) of crystal grains was determined.
In addition, using the Crystal Direction Map, the area ratio of α crystal grains ( Total Fraction) was obtained.

Furthermore, after setting the grain size to more than 20 μm in the grain properties of the partition, a crystal direction map is created, and the normal direction of the (0001) plane of the α crystal grain and the radial direction and the circumferential direction of the titanium alloy bar are formed. The area ratio (Total Fraction) of α-crystal grains having an angle in the range of 0° or more and 55° or less was obtained.

In addition, the Dwell fatigue properties of the obtained titanium alloy bar were measured.
As the test pieces, tensile test pieces and fatigue test pieces were taken such that the radial direction and the circumferential direction of the titanium alloy bar were the longitudinal directions.

The measurement conditions of the tensile test were as follows.
Specimen shape: parallel part φ5×30 mm, gauge length 25 mm, strain rate: 8.3×10 ⁻⁵ s ⁻¹ .

The measurement conditions of the fatigue test were as follows.
Fatigue test piece shape: Parallel part φ5.08 mm×15.24 mm, gauge length 12 mm.
Fatigue test method: axial force, pulsating, stress ratio 0.05. Maximum stress = 95% of the 0.2% proof stress of the same material (same direction).
Normal fatigue: triangular wave, load 1s, unload 1s
Dwell fatigue: trapezoidal wave, load 1s, hold 120s, unload 1s

Table 3 shows the area ratio of α crystal grain aggregates having an equivalent circle diameter of less than 100 μm and an aspect ratio of less than 3.0 in a cross section parallel to the major axis direction of the titanium alloy bar, and the number of (0001) planes of the α crystal grains. The area ratio of α-crystal grains with an equivalent circle diameter of more than 20 μm, where the angle θ between the normal direction and the major axis direction of the titanium alloy rod is in the range of 0° to 55°, and the normal to the (0001) plane The angle ω between the direction and the circumferential direction of the titanium alloy bar is in the range of 0° or more and 55° or less, and shows the area ratio of α crystal grains with an equivalent circle diameter of more than 20 μm. Table 3 also shows the ratio of the longer of the Dwell fatigue life divided by the shorter of the radial and circumferential Dwell fatigue lives in the orthogonal cross section. In the examples within the scope of the present invention, the Dwell fatigue rupture life was greater than 5000 cycles.

As shown in Table 3, in the examples within the scope of the present invention, the ratio of the Dwell fatigue life in the orthogonal cross section is 2.0 or less, indicating that the anisotropy is small. On the other hand, in the comparative example, which is outside the scope of the present invention, it can be seen that the anisotropy of the Dwell fatigue characteristics is increased.

1... Anvil, 2... Billet.

Claims (5)

The chemical components are Al: 5.50 to 6.75% by mass, V: 3.5 to 4.5% by mass, Fe: 0.05 to 0.40% by mass, O: 0.05 to 0.25% by mass % with the balance being Ti and impurities, or the chemical components are Al: 5.50 to 6.50% by mass, Sn: 1.75 to 2.25% by mass, Zr : 3.5 to 4.5% by mass, Mo: 1.8 to 2.2% by mass, Fe: 0.02 to 0.25% by mass, O: 0.02 to 0.15% by mass, An α+β-type titanium alloy bar, the balance of which is Ti and impurities ,
In a cross section parallel to the major axis direction of the α+β type titanium alloy bar, the area ratio of α crystal grain aggregates having an equivalent circle diameter of less than 100 μm and an aspect ratio of less than 3.0 is 60% or more,
The normal direction of the (0001) plane of the close-packed hexagonal crystal that constitutes the α crystal grains, the longitudinal direction of the α+β-type titanium alloy rod, and the circumferential direction of the α+β-type titanium alloy rod in the cross section perpendicular to the major axis. The α + β type titanium alloy bar, characterized in that the area ratio of α crystal grains having an equivalent circle diameter of more than 20 μm in the range of 0 ° to 55 ° and ω formed with each is 15% or less . A first step of heating a titanium alloy billet obtained by hot working an ingot to a temperature in the β single phase region and then quenching it;
a second step of heating the titanium alloy billet to a temperature in the α+β two-phase region, forging the titanium alloy billet over its entire length in a direction intersecting the longitudinal direction of the titanium alloy billet, and then cooling;
The titanium alloy billet is heated to a temperature in the α + β two-phase region and not higher than the heating temperature in the second step, and the titanium alloy billet is extended from a direction intersecting the major axis direction of the titanium alloy billet to the entire length of the billet. and a third step of forging over and at least finally cooling, in that order,
In the second step, in the cross section perpendicular to the longitudinal direction of the titanium alloy billet, the maximum width passing through the center of gravity of the perpendicular cross section is W1, the minimum width passing through the center of gravity of the perpendicular cross section is W2, and W1/W2 after forging. is 1.3 or less, and the ratio ΔW (ΔW= Wafter/Wini ) between the width Wini of the titanium alloy billet before forging and the width Wafter after forging is 1.05 or less; a step of performing a first rotation step of rotating the titanium alloy billet around the longitudinal axis and a first feeding step of moving the titanium alloy billet in the longitudinal direction, wherein the first forging step is performed at least twice; A step of forging the titanium alloy billet over its entire length by performing the above and performing the first feeding step after performing the first rotating step between the first forging steps,
The third step includes a second forging step of performing a reduction such that W1/W2 after forging is 1.5 or less, a second rotating step of rotating the titanium alloy billet around the long axis, and the titanium a step of performing a second feeding step of moving the alloy billet in the longitudinal direction, wherein the second forging step is performed at least twice, and after performing the second rotating step between the second forging steps a step of forging the titanium alloy billet over its entire length by performing the second feeding step;
The forging ratio in the second step is 1.6 or less, and the forging ratio in the third step is 1.5 or more and 2.5 or less,
The method for producing an α+β type titanium alloy bar according to claim 1, characterized in that: 3. The method for producing an α+β type titanium alloy bar according to claim 2 , wherein said first step is a step of heating said titanium alloy billet to a temperature in the β single phase region, working it, and then quenching it. The chemical components of the titanium alloy billet are Al: 5.50 to 6.75 mass%, V: 3.5 to 4.5 mass%, Fe: 0.05 to 0.40 mass%, O: 0.05. 4. The method for producing an α+β type titanium alloy bar according to claim 2 or 3 , wherein the Ti content is 0.25% by mass and the balance is Ti and impurities. The chemical components of the titanium alloy billet are Al: 5.50 to 6.50% by mass, Sn: 1.75 to 2.25% by mass, Zr: 3.5 to 4.5% by mass, Mo: 1.8 2.2% by mass, Fe: 0.02 to 0.25% by mass , O: 0.02 to 0.15% by mass, and the balance consisting of Ti and impurities. A method for producing an α + β type titanium alloy bar.

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