Titanium alloys, titanium alloy rods, titanium alloy plates, and engine valves

Abstract

To provide a titanium alloy, a titanium alloy bar, a titanium alloy plate, and an engine valve that maintain ductility and are excellent in fatigue characteristics.SOLUTION: An α+β type titanium alloy of Ti-5Al-2Fe-3Mo base or Ti-6Al-4V base has a two-phase structure composed of an α phase and a β phase, and a maximum area of a segregation abnormal part of the concentration of elements Al, Fe, Mo, and V in a cross section of 2,000 μm2 or less. The segregation abnormal part of the concentration of elements Al, Fe, Mo, and V in a cross section means an aggregate of neighboring measurement points that deviate from average value ±1 mass% of Al and Fe, average value ±1.5 mass% of Mo, and average value ±2 mass% of V in EPMA plane analysis. There are thus provided a titanium alloy, a titanium alloy bar, a titanium alloy plate, and an engine valve that are excellent in fatigue characteristics and ductility.SELECTED DRAWING: None

Description

[Technical Field]

[0001] This invention relates to titanium alloys, titanium alloy rods, titanium alloy plates, and engine valves. [Background technology]

[0002] Titanium alloys are lightweight yet strong and have good corrosion resistance, making them suitable for a wide range of applications. Among them, α+β type titanium alloys, such as Ti-6Al-4V, offer an excellent balance of mechanical properties including strength, ductility, and toughness. They have long been widely used in the aerospace sector, and in recent years, their application to automotive parts has also been increasing.

[0003] On the other hand, there was a demand for α+β type titanium alloys with superior properties and lower costs.

[0004] Patent Document 1 discloses a high-strength α+β type titanium alloy characterized by containing, by mass%, 4.4% to less than 5.5% Al, 1.4% to less than 2.1% Fe, and 1.5% to less than 5.5% Mo, with impurities such as Si being suppressed to less than 0.1% and C to less than 0.01%, and the remainder consisting of Ti and unavoidable impurities. By adding an appropriate amount of Mo to an α+β type titanium alloy containing Al and Fe, an α+β type titanium alloy with high strength, high ductility, and excellent hot and cold workability has been discovered, and by further adding a fourth element, an α+β type titanium alloy with excellent corrosion resistance has been found.

[0005] Patent Document 2 discloses an α+β type titanium alloy member that contains, by mass%, 4.4% to less than 5.5% Al, 1.4% to less than 2.1% Fe, and 2.5% to less than 5% Mo, with impurities of less than 0.1% Si and less than 0.01% C, and the remainder being Ti and unavoidable impurities, characterized in that the area ratio "A" of the protoprecipitation α phase grains is 5% to less than 49% and the Young's modulus is 75 GPa to less than 100 GPa, and the tensile strength is in the 1000 MPa class or higher. This makes it possible to provide an α+β type titanium alloy member with a tensile strength of 1000 MPa or higher, having a Young's modulus of 75 GPa or more and less than 100 GPa, which is comparable to or lower than that of a β type titanium alloy, using an α+β type titanium alloy with a relatively inexpensive alloy composition, and a method for manufacturing the same. Furthermore, it makes it possible to provide a method for manufacturing an α+β type titanium alloy member with a tensile strength of 1000 MPa or higher, which allows for a wider range of adjustment of the Young's modulus so that it is between 75 and 125 GPa, without changing the alloy composition.

[0006] Patent Document 3 discloses an α+β type titanium alloy consisting of 1.4% to less than 2.1% Fe, 4.4% to less than 5.5% Al, the remainder being titanium and impurities; an α+β type titanium alloy in which part of the Fe is replaced with less than 0.15% Ni, less than 0.25% Cr, and less than 0.25% Mn; or an α+β type titanium alloy further containing 0.05% to less than 0.25% Si. This makes it possible to provide an Al-Fe-based α+β type titanium alloy that has fatigue strength equivalent to conventional alloys, as well as superior hot and cold workability compared to conventional alloys.

[0007] Patent Document 4 discloses an α+β type titanium alloy consisting of 0.5% to less than 1.4% Fe, 4.4% to less than 5.5% Al, the remainder being titanium and impurities, or an α+β type titanium alloy in which a portion of Fe is replaced with one or more of the following: less than 0.15% Ni, less than 0.25% Cr, or less than 0.25% Mn, or an α+β type titanium alloy further containing 0.05% to less than 0.25% Si. This invention can provide a titanium alloy with stable fatigue strength with little variation and high hot workability. Alternatively, it can provide a titanium alloy that also possesses even higher creep resistance.

[0008] Patent Document 5 discloses an α+β type titanium alloy formed from α and β phases, wherein the area ratio of the α phase in the microstructure is 40% or less, and the average equivalent circle diameter is 5 μm or less and the average aspect ratio is 3 or more, and the α phase area ratio is 40% or less, and the alloy has excellent machinability. As a manufacturing method, a cast slab is hot forged according to a conventional method, then subjected to processing heat treatment to process it into the desired shape, and then annealed.

[0009] Patent Document 6 discloses a corrosion-resistant Fe-containing titanium material characterized by containing 0.1 to 2.5% by mass of Fe, having a needle-like or lath-like microstructure (structure a), and having an area of ​​5% or more in the concentration distribution analyzed by electron beam microanalyzer (EPMA) where the Fe concentration is concentrated to 1.5 times or more the average concentration. By having structure a and the above-mentioned Fe-enriched area, the corrosion resistance of the titanium material can be improved.

[0010] Patent Document 7 describes a method having predetermined component elements and the maximum value C of each element's measured value. MAX and minimum value C MIN The difference ΔC is 0.2C MINA titanium ingot is disclosed that has an elemental concentration of less than 0.04% or less than 0.04%, and whose microstructure has a circular average grain size of 10 mm or less in the center of the thickness direction of the titanium ingot, and is less than half the thickness, and is easily processable into thin plates or wires. However, the elemental concentration analysis uses conventional chemical analysis methods and does not evaluate microscopic concentration variations on the μm scale, nor does it mention the impact on various properties, including fatigue properties. [Prior art documents] [Patent Documents]

[0011] [Patent Document 1] Japanese Patent Publication No. 2005-320618 [Patent Document 2] Japanese Patent Publication No. 2007-314834 [Patent Document 3] Japanese Patent Application Publication No. 7-62474 [Patent Document 4] Japanese Patent Application Publication No. 7-70676 [Patent Document 5] Japanese Patent Publication No. 2005-320570 [Patent Document 6] Japanese Patent Publication No. 2005-336551 [Patent Document 7] International Publication No. WO2019 / 026251 [Overview of the project] [Problems that the invention aims to solve]

[0012] Further improvements in fatigue properties are required for α+β type titanium alloys. In particular, it is desirable to suppress premature failure in the fatigue fracture of engine valves, which are manufactured by machining or other processes using rod stock as the raw material. Furthermore, increasing the strength to improve fatigue properties reduces ductility, raising concerns that the material may fracture or form voids internally when deformed at room temperature, making it difficult to handle. Therefore, sufficient ductility at room temperature is also required.

[0013] An object of the present invention is to provide a titanium alloy, a titanium alloy bar, a titanium alloy plate, and an engine valve that maintain ductility and have excellent fatigue characteristics.

Means for Solving the Problems

[0014] That is, the gist of the present invention is as follows. [1] A titanium alloy having a two-phase structure composed of an α-phase and a β-phase, containing 2.5 mass% or more of Al and at least 0.8 mass% or more of one or more main elements other than Al, and the maximum area of the segregation abnormal part of the main element concentration in the cross section is 2000 μm 2 The following titanium alloy. Here, the main elements are Al, Fe, Mo, and V, and the segregation abnormal part of the main element concentration in the cross section means an aggregate of measurement points where the main elements other than Mo and V deviate from the average value ±1 mass%, Mo deviates from the average value ±1.5 mass%, and V deviates from the average value ±2% in the EPMA surface analysis and adjacent points. [2] Containing, in mass%, Al: 4.5 to 6.5%, Fe: 1.4 to 2.3%, Mo: 1.5 to 5.5%, with the total of O and N being 0.25% or less, and the balance being Ti and impurities, the titanium alloy according to [1]. [3] Containing, in mass%, Al: 5.0 to 7.0%, V: 3.5 to 5.0%, with the total of O and N being 0.25% or less, and the balance being Ti and impurities, the titanium alloy according to [1]. [4] Containing, in mass%, Al: 4.5 to 6.5%, Fe: 0.8 to 2.3%, Si: 0.0 to 0.50%, with the total of O and N being 0.25% or less, and the balance being Ti and impurities, the titanium alloy according to [1].

[0015] [5] A titanium alloy bar made of the titanium alloy according to any one of [1] to [4]. [6] The titanium alloy bar according to [5], in which the ratio of the acicular structure is 80 area% or less. Here, the proportion of needle-like structures is determined by performing EBSD evaluation on a cross section perpendicular to the axis of the titanium alloy rod, determining the grain size with an orientation difference of 5° or more as the grain boundary, and setting the proportion of needle-like structures as 100 area % if the average grain size is 100 μm or more. If the average grain size is less than 100 μm, data with an Image Quality (IQ) of 20% or more of the average is extracted, and from the extracted data, α grains with an aspect ratio (long axis / short axis) of 3 or less are defined as equiaxed α grains, and the area ratio of grains other than the aforementioned equiaxed α grains to the total measured area is defined as the proportion of needle-like structures.

[0016] A titanium alloy plate made from any one of the titanium alloys described in [7][1] to [4]. [8] A titanium alloy sheet as described in [7], wherein the proportion of needle-like structures is 80 area % or less. Here, the proportion of needle-like structures is determined by performing EBSD evaluation on a cross section perpendicular to the rolling direction of the titanium alloy sheet, determining the grain size with an orientation difference of 5° or more as the grain boundary, and setting the proportion of needle-like structures as 100 area % if the average grain size is 100 μm or more. If the average grain size is less than 100 μm, data with an Image Quality (IQ) of 20% or more of the average is extracted, and from the extracted data, α grains with an aspect ratio (long axis / short axis) of 3 or less are defined as equiaxed α grains, and the area ratio of grains other than the aforementioned equiaxed α grains to the total measured area is defined as the proportion of needle-like structures.

[0017] An engine valve made of a titanium alloy as described in any one of [9][1]~[4].

[10] An engine valve according to [9], wherein the proportion of needle-like structures is 80 area % or less. Here, the proportion of needle-like structures is determined by performing EBSD evaluation on a cross section perpendicular to the axis of the engine valve, determining the grain size with an orientation difference of 5° or more as the grain boundary, and setting the proportion of needle-like structures as 100 area % if the average grain size is 100 μm or more. If the average grain size is less than 100 μm, data with an Image Quality (IQ) of 20% or more of the average is extracted, and from the extracted data, α grains with an aspect ratio (long axis / short axis) of 3 or less are defined as equiaxed α grains, and the area ratio of grains other than the aforementioned equiaxed α grains to the total measured area is defined as the proportion of needle-like structures. [Effects of the Invention]

[0018] This invention relates to an α+β type titanium alloy having a two-phase structure consisting of an α phase and a β phase, with a maximum size of 2000 μm for segregation abnormalities in the concentration of major elements such as Al, Fe, Mo, and V in the cross-section. 2 The following conditions make it possible to provide titanium alloys, titanium alloy rods, titanium alloy plates, and engine valves with excellent fatigue properties. [Modes for carrying out the invention]

[0019] Composition of Titanium Alloys The alloy composition specified in this invention is the average analytical value of the entire product. The analytical sample is prepared by removing the top 1 mm and uniformly sampling from the remaining portion. The analytical methods used are inductively coupled plasma (ICP) emission spectrometry for metallic elements, inert gas fusion infrared absorption spectrometry for O, inert gas fusion thermal conductivity spectrometry for N and H, and high-frequency combustion infrared absorption spectrometry for C. Hereafter, % in the alloy composition refers to mass %.

[0020] 《α+β type titanium alloy》 The titanium alloys covered by this invention are, firstly, broadly α+β type titanium alloys. α+β type titanium alloys have a two-phase structure consisting of an α phase and a β phase, contain 2.5 mass% or more of Al, and further contain 0.8 mass% or more of at least one of Fe, V, or Mo. Specifically, these include titanium alloys such as Ti-6Al-4V(ELI), Ti-3Al-2.5V, Ti-6Al-2Sn-4Zr-2Mo, Ti-8Al-1V-1Mo, Ti-6Al-2Sn-4Zr-6Mo, Ti-6Al-6V-2Sn, and Ti-4.5Al-3V-2Fe-2Mo as listed in the AMS standard, as well as titanium alloys such as Ti-5Al-1Fe, Ti-5Al-1.5Fe-0.25Si, Ti-5Al-2Fe-3Mo, and Ti-8Al-1Fe-1Nb-0.15Si. This invention encompasses all of these α+β type titanium alloys.

[0021] 《Ti-5Al-2Fe-3Mo alloy》 Secondly, the titanium alloy composition targeted by this invention is an α+β type titanium alloy containing Al: 4.5-6.5%, Fe: 1.4-2.3%, and Mo: 1.5-5.5%, and is hereinafter referred to as the "Ti-5Al-2Fe-3Mo alloy." The Ti-5Al-2Fe-3Mo alloy can achieve high strength of more than that of the general-purpose Ti-6Al-4V in the solution state and is a component system that can utilize relatively inexpensive elements. It is a titanium alloy with a component composition similar to the titanium alloys described in Patent Documents 1 and 2.

[0022] (Al: 4.5-6.5%) The higher the Al content, the higher the strength. It also has effects such as controlling the ratio of α-phase to β-phase and suppressing the precipitation of the ω-phase, which causes embrittlement. Therefore, it is common to add 4.5% or more Al, and more preferably 4.8% or more. Adding too much Al reduces hot workability, so the upper limit is 6.5%, and more preferably 6.0% or less.

[0023] (Fe: 1.4~2.3%) The strength increases with increasing Fe content. Furthermore, the β-phase ratio at high temperatures increases, improving hot workability; therefore, Fe should be added at a concentration of 1.4% or more, more preferably 1.8% or more. On the other hand, excessive addition leads to significant segregation in large ingots, which cannot be resolved even with subsequent segregation mitigation measures. Therefore, the upper limit for Fe addition is 2.3%, more preferably 2.1% or less.

[0024] (Mo: 1.5~5.5%) The strength increases with increasing Mo content. Furthermore, the β-phase ratio at high temperatures increases, improving hot workability; therefore, Mo should be added at a concentration of 1.5% or more, more preferably 2.0% or more. On the other hand, adding too much Mo leads to significant segregation in large ingots, and this segregation cannot be resolved even with subsequent segregation mitigation measures. Therefore, the upper limit for Mo addition is 5.5%, more preferably 4.5% or less.

[0025] (O, N, C, H) O, N, C, and H are present in titanium alloys as impurity elements. O and N increase strength but decrease ductility. Therefore, the total amount of O and N should be 0.25% or less. In practice, the amount of N is small and the amount of O is large. There is no particular lower limit, but to ensure strength, it is preferable to have a total of 0.05% or more. As the carbon content increases, the strength increases, but the ductility decreases, as does the hot workability. Therefore, it is preferable to keep it at 0.05% or less. There is no particular lower limit, but it is practically 0.001% or more. H is an element that causes embrittlement if its content is high, so it is preferably kept at 0.013% or less to prevent embrittlement. There is no particular lower limit, but it is practically 0.0001% or more.

[0026] 《Ti-6Al-4V alloy》 The titanium alloy composition targeted by this invention is, thirdly, an α+β type titanium alloy containing Al: 5.0-7.0% and V: 3.5-5.0%, and hereinafter referred to as the "Ti-6Al-4V alloy." This is a component system that encompasses the representative α+β titanium alloy generally known as Ti-6Al-4V.

[0027] (Al: 5.0~7.0%) The more Al is added, the higher the strength. It also has effects such as controlling the ratio of α-phase to β-phase and suppressing the precipitation of the ω-phase, which causes embrittlement. For this reason, it is common to add 5.0% or more Al, and more preferably 5.5% or more. Adding too much Al reduces hot workability, so the upper limit of Al content is 7.0%, and more preferably 6.5% or less.

[0028] (V: 3.5~5.0%) The more V added, the higher the strength. Furthermore, the β-phase ratio increases at high temperatures, improving hot workability; therefore, V should be added at a concentration of 3.0% or more, more preferably 3.5% or more. However, since increasing the amount added increases costs, the upper limit for V addition is set at 5.0%, more preferably 4.5% or less.

[0029] (O, N, C, H) O and N increase strength but decrease ductility. Therefore, their total content should be 0.25% or less. In practice, the amount of N is small, and the amount of O is large. To ensure sufficient strength, a total content of 0.05% or more is required. As the amount of carbon (C) increases, the strength increases, but the ductility decreases, as does the hot workability. Therefore, the amount should be kept below 0.05%. There is no particular lower limit, but in practice, it should be 0.001% or higher. H is an element that causes embrittlement when added in large amounts, so its concentration should be kept below 0.013% to prevent embrittlement. There is no specific lower limit, but in practice it should be 0.0001% or higher.

[0030] 《Ti-5Al-1.5Fe―Si alloy》 The titanium alloy composition targeted by this invention is, fourthly, an α+β type titanium alloy containing Al: 4.5-6.5%, Fe: 0.8-2.3%, and Si: 0.0-0.50%, and is hereinafter referred to as the "Ti-5Al-1.5Fe-Si alloy." This compositional system encompasses the α+β type titanium alloys described in Patent Documents 3 and 4.

[0031] (Al: 4.5-6.5%) The higher the Al content, the higher the strength. It also has effects such as controlling the ratio of α-phase to β-phase and suppressing the precipitation of the ω-phase, which causes embrittlement. Therefore, it is common to add 4.5% or more Al, and more preferably 4.8% or more. Adding too much Al reduces hot workability, so the upper limit is 6.5%, and more preferably 6.0% or less.

[0032] (Fe: 0.8~2.3%) The strength increases with increasing Fe content. Furthermore, the β-phase ratio at high temperatures increases, improving hot workability; therefore, Fe should be added at a concentration of 0.8% or more, more preferably 1.0% or more. On the other hand, if too much Fe is added, segregation becomes significant in large ingots, and even if segregation is mitigated in subsequent processes, the segregation will not be resolved. Therefore, the upper limit for Fe addition is 2.3%, more preferably 2.1% or less.

[0033] (Si: 0.0~0.50%) Si content increases with increasing strength. Furthermore, oxidation at high temperatures is suppressed, so it is added as needed. On the other hand, excessive addition can lead to the formation of coarse precipitates in large ingots, becoming a starting point for fatigue fracture. Therefore, the upper limit for Si content is 0.5%, more preferably 0.4% or less. Si content is not required.

[0034] (O, N, C, H) O, N, C, and H are present in titanium alloys as impurity elements. O and N increase strength but decrease ductility. Therefore, the total amount of O and N should be 0.25% or less. In practice, the amount of N is small and the amount of O is large. There is no particular lower limit, but to ensure strength, it is preferable to have a total of 0.05% or more. As the carbon content increases, the strength increases, but the ductility decreases, as does the hot workability. Therefore, it is preferable to keep it at 0.05% or less. There is no particular lower limit, but it is practically 0.001% or more. H is an element that causes embrittlement if its content is high, so it is preferably kept at 0.013% or less to prevent embrittlement. There is no particular lower limit, but it is practically 0.0001% or more.

[0035] Other element content: less than 0.5% In all of the above titanium alloys 2 through 4, elements other than those listed above are likely to be present as impurities: Sn, Ni, Cr, Mn, Zr, Nb, Cu, (Si, V, Mo). These may be present from the raw materials, especially when scrap or lower-grade sponge titanium is used. These are elements that should be controlled as there is a concern that their content may increase. It is preferable that the total content of these elements be less than 0.5% (preferably less than 0.3%), and that each individual element be 0.1% or less. In addition, there are other elements such as platinum group metals that are contained in corrosion-resistant alloys, but these do not have a significant impact on mechanical properties, so their presence does not pose a problem and therefore do not need to be controlled. Other elements are hardly present in general-purpose alloys, so the possibility of contamination from scrap is low, and contamination from sponge titanium is also low.

[0036] 《Crystal structure of titanium alloys (1)》 The titanium alloy having the component composition of the present invention is an α+β type titanium alloy, and has both an α phase and a β phase in the titanium alloy. In the composition range of α+β type titanium alloys, at high temperatures, two phases, α and β, exist. As the temperature increases, the β phase increases, and the distribution of alloying elements occurs between the α and β phases, resulting in an unstable β phase at room temperature. Therefore, if the alloy is held at low temperatures below 500°C or rapidly cooled, ω, α', and α'' phases may form. The formation of these phases makes the alloy harder and reduces its ductility and toughness at room temperature. Therefore, care must be taken to avoid excessive deformation when handling the alloy at room temperature. The ω phase has a particularly significant impact. The α' and α'' phases are formed by rapid cooling from high temperatures where the β phase is abundant, or by low-temperature holding, and are therefore rarely formed in commonly used manufacturing methods. Even if the α' or α'' phases are formed, they will only make up about 0.1% and will not have any effect. Therefore, it is sufficient to check only the ω phase, and it is sufficient if no peaks are detected by X-ray diffraction. X-ray diffraction is performed using Cu-Kα or Co-Kα rays. 2θ is set to 30-90°, and the scan step is 0.01°. As described above, the titanium alloy of the present invention, in which the ω phase is not confirmed by X-ray diffraction, has a two-phase structure consisting of an α phase and a β phase.

[0037] Uniformity of the compositional composition of titanium alloys There are various causes of fatigue fracture in titanium alloys, but one example is that when there is a difference in strength within the material, that area is likely to become the fracture initiation point. Furthermore, the larger the fracture initiation point, the more likely it is to lead to fracture. Often the fracture initiation point is at the crystal grain level, but it will result in a larger initiation point. There are several reasons for this, one of which is segregation that occurs during the solidification of the alloy. In particular, alloys intended for high strength tend to have a large amount of added elements, making them prone to segregation. To prevent fatigue fracture of a product, it is desirable to minimize segregation in the product. The same applies to maintaining ductility. In addition, large ingots are common to improve productivity and provide products at a low cost, but large ingots are prone to segregation. Therefore, we conceived the idea that even in products manufactured using large ingots, reducing the segregation of alloying elements can improve fatigue properties and maintain ductility. Furthermore, since this phenomenon is caused by a difference in strength, it is desirable to confirm it by the difference in strength, but because titanium is anisotropic, even if evaluated by hardness, the value obtained from the evaluation is not necessarily a useful value in the direction of fatigue load. Therefore, it is important to evaluate the amount of elements that are the source of strength differences, regardless of anisotropy. The elements that are the starting point for fatigue fracture and the source of strength differences are substitutional strengthening elements such as Al, Fe, Mo, and V, and the effect of improving fatigue properties through the uniformity of the component composition has been confirmed. In this invention, these are referred to as the main elements in titanium alloys. Interstitial strengthening elements, which diffuse easily and do not cause segregation, are considered to be excluded, but this does not mean that other substitutional strengthening alloying elements are excluded.

[0038] Large ingots (at least 1 ton in this case) are particularly prone to solidification segregation. Unevenness in the raw material composition is also considered a contributing factor to segregation. Because these segregations are formed over relatively large areas (several millimeters or more) and over long distances, it is difficult to achieve uniformity solely through the diffusion of alloying elements by heat treatment. It is necessary to utilize both metal flow due to processing and diffusion due to heat treatment.

[0039] In particular, the reduction in cross-sectional area due to processing significantly reduces and shortens the diffusion distance during heat treatment, making it easier to obtain a uniform elemental distribution with a short heat treatment time.

[0040] On the other hand, with respect to rods, the reduction in cross-sectional area is smaller compared to plates, so segregation of the alloying elements becomes more pronounced, and fractures caused by this become more apparent. More specifically, in engine valves manufactured using rods as the material, high fatigue strength is required, and premature fracture due to segregation can be a problem, so it is hoped that this will be suppressed.

[0041] Therefore, we aimed to clarify the effect of variations in the composition of titanium alloys (products after hot working) (segregation abnormalities) on their fatigue properties, and to realize a manufacturing method that can obtain sufficient ductility. In particular, we achieved an improvement in fatigue strength for titanium alloy rods, specifically round bars.

[0042] (The maximum size of the segregation anomaly in the concentration of major elements in the cross-section is 2000 μm) 2 below) Hereinafter, "major elements" refers to Al, Fe, Mo, and V. The α+β type titanium alloy of the present invention contains Al and at least one other major element in an amount of 1% by mass or more. As will be shown in the examples described later, the segregation of the main elements in the alloy is 1 × 10⁻⁶ 7 It was found to affect the fatigue limit (hereinafter simply referred to as the fatigue limit).

[0043] First, a Ti-5Al-2Fe-3Mo alloy system will be exemplified. In specimens with a fatigue limit of approximately 650 MPa, there were several test pieces that fractured even at 600 MPa. The vicinity of the fatigue initiation point of these test pieces was mirror-polished, and surface analysis by EPMA was performed. In the EPMA analysis, a region with an accelerating voltage of 10 kV and an angle of 3 mm was measured in 10-μm steps. Regardless of the fatigue life, the alloying elements were distributed approximately in a normal distribution, but in the case of low fatigue life, large regions with enrichment or depletion of Al, Fe, and Mo were formed near the fatigue initiation point. Regarding this region, when regions where Al and Fe deviated from the average value by ±1 mass% and Mo deviated from the average value by ±1.5 mass% were illustrated as segregation regions for Al, Fe, and Mo, the size (area) of the segregation region of any element in the test pieces fractured at 600 MPa was 3000 μm 2 equivalent, resulting in a result corresponding to the fatigue test.

[0044] Next, a Ti-6Al-4V alloy system will be exemplified. In the case of this alloy, in specimens with a fatigue limit of approximately 630 MPa, there were several test pieces that fractured even at 600 MPa. When the same investigation as the above Ti-5Al-2Fe-3Mo titanium alloy was carried out, when regions where Al deviated from the average value by ±1 mass% and V deviated from the average value by ±2% were illustrated as segregation regions, the size (area) of the segregation region of any element in the test pieces fractured at 600 MPa was 3500 μm 2 equivalent, resulting in a result corresponding to the fatigue test.

[0045] Such segregation regions where alloying elements are enriched or depleted are called segregation abnormal parts. The segregation abnormal parts can be evaluated as regions that fluctuated by more than different values for each element with respect to the average value in the surface analysis of EPMA of a predetermined element. In the case of the Ti-5Al-2Fe-3Mo alloy of the present invention, targeting strengthening elements, Al and Fe are ±1 mass%, Mo is ±1.5 mass%, and both enrichment indicated by + and depletion indicated by - are the same, and both can be regarded as segregation regions constituting the segregation abnormal parts. The fatigue initiation point is the weakest part, and the size (area) of the largest segregation abnormal part is important. In the test pieces with a high fatigue life, the size of the segregation abnormal part was at most 2000 μm 2The results were as follows. Therefore, the maximum area of ​​the segregation anomaly in the elemental concentrations of Al, Fe, and Mo in the cross-section was 2000 μm². 2 It was found that fracture due to segregation can be avoided if the following conditions are met. In the case of Ti-6Al-4V alloys, focusing on the strengthening elements, Al is at an average value of ±1 mass%, and V is at an average value of ±2%. Both the enrichment (indicated by +) and the decrease (indicated by -) are the same, and both can be considered as segregation regions constituting the segregation anomaly. Furthermore, in the specimens with a long lifespan, the maximum area of ​​segregation abnormalities was 2000 μm². 2 The results were as follows. Furthermore, these results (segregation of each element and the maximum area of ​​segregation anomalies) were similar for Ti-5Al-1.5Fe-Si alloys.

[0046] Based on the above results, in EPMA surface analysis, a collection of adjacent measurement points where major elements other than Mo and V are outside the mean ±1 mass%, Mo is outside the mean ±1.5 mass%, and V is outside the mean ±2% can be called a segregation region.

[0047] Segregation anomalies are evaluated as a collection of adjacent points in EPMA surface analysis mapping, and can be determined by multiplying the number of adjacent measurement points that satisfy the above criteria for segregation regions by the square of the step size. Adjacent measurement points are defined as those that touch on one side of a square formed by each measurement point.

[0048] When evaluating a bar, the measurement surface should be either a cross-section perpendicular to the bar's axis or a cross-section including the axis. For a plate, the cross-section should be perpendicular to the rolling direction. A wide measurement area is desirable, and measurements should be taken over an area with a corner of 3 mm or more. In addition, the measurement area should be as large as possible relative to the measurement cross-sectional size, either square or rectangular. For rectangular shapes, the ratio of the long side to the short side (long side / short side) should be 1.5 or less. If a corner of 3 mm or more cannot be secured in one field of view, measurements should be taken in multiple fields of view with a corner of 4 mm or more. This is because accurate measurement becomes difficult if segregation anomalies are located at the edges of the field of view, so a wider area than a single field of view is evaluated. When measuring, the top 0.1 mm should not be included in the measurement range because of the influence of contamination near the surface and the difficulty of accurate measurement at edges such as the surface.

[0049] The EPMA acceleration voltage is set to 10-15kV, the step size to 0.01-0.05mm, and the beam diameter during measurement is set to 0.5-1 times the step size.

[0050] Quantification of measurement results is performed by setting the measurement results of pure titanium to an intensity of 0% for each element, and creating a calibration curve using the measurement results of model alloys with the chemical composition of the target α+β alloy, such as Ti-5Al-2Fe-3Mo, Ti-6Al-4V alloys, or Ti-5Al-1.5Fe-Si alloys. For example, a model alloy with the chemical composition of Ti-5Al-2Fe-3Mo is prepared by, for example, creating a small arc-molten ingot (about 100g), heating it to 1100°C to eliminate solidification defects, and processing it to a cross-sectional area of ​​20% or more, resulting in a cross-sectional area of ​​100mm². 2 The material is shaped as described above. It can then be fabricated by annealing at 1100°C for 10-60 minutes and then air-cooling. After removing the top 1 mm, each element is analyzed by ICP emission spectrometry, and this is used as the composition of the model alloy. Therefore, when creating the calibration curve, the top 1 mm should not be included in the measurement range, and measurements should be taken over the widest possible area. Surface analysis is preferable for this measurement, and the average value should be used as the measurement result. The same applies to other alloy compositions.

[0051] When the product is manufactured into an engine valve or similar object, evaluation can be performed on the cut surface obtained by cutting the object to ensure a sufficient measurement area.

[0052] 《Crystal structure of titanium alloys (2)》 (Percentage of needle-like tissue is 80% or less) In α+β type titanium alloys, equiaxed and acicular structures appear as crystalline structures. Equiaxed structures have higher ductility at room temperature and are easier to handle at room temperature compared to acicular structures. Also, since acicular structures have low ductility at room temperature, a smaller amount of acicular structure results in better handling at room temperature. The presence of grain boundary α grains is a reason for the low ductility of needle-like structures. However, in high-alloy materials like the present invention, if the proportion of needle-like structures is high, the ductility at room temperature is low even without the presence of grain boundary α grains. Therefore, it is preferable to keep the proportion of needle-like structures low. The proportion of needle-like structures is expressed as an area percentage and is preferably 80% or less. More preferably 75% or less, and even more preferably 70% or less.

[0053] The needle-like structure in this context refers to regions where finely elongated α-phase particles precipitate within the β-phase. Because the α-phase in the needle-like structure is so fine, it is difficult to recognize it as individual crystal grains. Therefore, the results obtained by EBSD are used to analyze and determine the structure using the method described later. EBSD measurements are performed on a cross-section perpendicular to the axis for titanium alloy rods, such as round bars. The measurement conditions are 500x magnification, covering an area of ​​200 μm or more, with a step size of 0.2 μm. Under these conditions, a total of three or more fields of view are measured: one field of view in the center of the round bar and two fields of view at R / 2. The R / 2 measurements are performed at any point and at a position rotated 90° from that point relative to the axis of the round bar. For plate materials, measurements are performed on a cross-section perpendicular to the rolling direction. Under the same conditions as for the round bar, a total of three or more fields of view are measured at the center of the plate thickness: one field of view in the center of the width, and two fields of view at 1 / 4 width and 3 / 4 width.

[0054] First, in the measurement results, grain boundaries are defined as orientation differences of 5° or more, and the grain size of each individual crystal is determined. Based on these values, the average grain size in the field of view is then calculated.

[0055] If the average grain size is 100 μm or more, the proportion of needle-like structures is considered to be 100%. Needle-like structures are formed when the β phase transforms into the α phase upon cooling. The formed α grains alternate with very thin β grains that have almost the same crystal orientation, forming regions called colonies. However, because the β phase is very thin, it is difficult to detect in measurements, and α grains that should be recognized as different α grains are mistakenly identified as adjacent. As a result, in the analysis to calculate the average grain size, α grains that should be considered as separate are judged as the same grain, and colonies are judged as a single crystal grain. In this case, the crystal size of the α grains is at least about 30% of the β grain size. In a state with 100% β phase, the β grain size easily exceeds 500 μm. However, even with prolonged heating and holding in the presence of equiaxed α grains, it is difficult for the average grain size of the α grains to exceed 100 μm. Therefore, if the average grain size is 100 μm or more, it can be determined that 100% needle-like structures are present. On the other hand, when the β-phase ratio is less than 100%, the β-grains do not coarse due to the pinning effect of the α-grains, and it is difficult for the average particle size of the α-grains after cooling to be 100 μm or more. Therefore, when the average particle size of the α-grains is less than 100 μm, a mixture of needle-like and equiaxed tissues is obtained.

[0056] If the average crystal grain size is less than 100 μm, first extract data with an Image Quality (IQ) of 20% or higher of the average. From the extracted data, determine the α grains (equisaxial α grains) with an aspect ratio (long axis / short axis) of 3 or less, calculate the area percentage of equiaxed α grains within the field of view, and define this as the equiaxed α grain percentage (area %). The crystalline structure is a mixture of equiaxed α grains and needle-like structures, where fine needle-shaped α grains precipitate within the β phase. The sum of the equiaxed α grain percentage (area %) and the percentage of needle-like structures (area %) is taken as 100%, and the percentage of needle-like structures (area %) (labeled "Needle-like Structure Percentage" in Table 3) is determined. The percentage of needle-like structures is calculated as the area ratio to the measured area.

[0057] Specifically, the proportion of needle-like structures is determined by performing EBSD evaluation on a cross section perpendicular to the axis in the case of alloy rods, and on a cross section perpendicular to the rolling direction in the case of plates. The grain size is determined by defining grain boundaries with an orientation difference of 5° or more in the measurement results. If the average grain size is 100 μm or more, the proportion of needle-like structures is set to 100 area %. If the average grain size is less than 100 μm, data with an Image Quality (IQ) of 20% or more of the average is extracted. From the extracted data, α grains with an aspect ratio (long axis / short axis) of 3 or less are defined as equiaxed α grains, and the area ratio of grains other than the aforementioned equiaxed α grains to the total measurement area is defined as the proportion of needle-like structures (area %).

[0058] Manufacturing method for titanium alloys (β-transformation point temperature) The β transformation point temperature Tβ (°C) used in the following manufacturing methods can be expressed as a formula for the major elements and the content of O, N, and C in the titanium alloy. For Ti-5Al-2Fe-3Mo alloys, it can be calculated based on formula (1). For Ti-6Al-4V alloys, it can be calculated based on formula (2). Tβ(℃)=880+23×Al+170×(O+N)+100×C-20×Mo-8×Fe (1) Tβ(℃)=875+23×Al+170×(O+N)+100×C-12×V (2) The element symbol represents the mass percentage value of the alloying element. For α+β type titanium alloys other than the two alloy systems mentioned above, Tβ (°C) can be determined from equation (3). Tβ(℃)=905+20×Al+170×(O+N)+100×C-20×Mo-15×Fe―12V (3)

[0059] (Ingot) The ingot used is manufactured using a conventional method, such as VAR, EB melting, or plasma melting. The ingot may be rectangular or cylindrical, and surface processing such as cutting may be performed as needed. The effects of the present invention can be particularly demonstrated when using an ingot weighing 1 ton or more.

[0060] (Billet forging) The ingot is forged into a billet, which will become the raw material for rolled bars, such as round bars. The shape of the billet can be any shape that allows for the manufacture of rolled round bars or plates, such as a round cross-section, square, rectangular, or octagon. The forging process consists of two stages: rough forging and billet finish forging. Both forging stages must be performed at a temperature higher than the β transformation point.

[0061] Rough forging is performed to eliminate internal defects formed by solidification shrinkage in the ingot, as well as to eliminate coarse solidification structures. Furthermore, it also serves to disperse segregated structures by utilizing metal flow during processing.

[0062] The ingot is heated to a temperature between the β transformation point and 1200°C to form a round cross-section of φ200-250 mm or a square cross-section of 200-250 mm. Alternatively, an octagonal cross-section with a similar cross-sectional area is also acceptable. For plate material, a rectangular cross-section is also acceptable. The heating temperature should be 50°C, more preferably 100°C or more, higher than the β transformation point. The upper limit of the temperature is set at 1200°C because raising it any higher would significantly reduce yield due to oxidation.

[0063] This invention reduces variations in alloy composition by performing one or more upsetting steps during forging. When manufacturing round bar-shaped products using conventional forging and hot rolling, reduction occurs from the radial direction of the round bar, but not from the longitudinal direction, resulting in insufficient metal flow to mitigate segregation. In contrast, performing one or more upsetting steps mitigates segregation and reduces variations in alloy composition. Upsetting is preferably performed at a stage where the ratio of the short side and short axis length Ls to the length L (L / Ls) of the cross-section is 1 to 4 or less. The larger L / Ls, the greater the possibility of buckling, so caution is required when performing upsetting with an L / Ls exceeding 4. The reduction ratio of the upsetting should be 20% to 60%. More preferably 25% or more, and more preferably 30% or more. Upsetting may be performed as a separate heat, or it may be performed in the same heat as forging. The same applies to sheet metal.

[0064] Depending on the ingot size and forging speed, cracks may occur if the temperature drops before the entire length is forged. In this case, reheating is necessary. This reheating is also treated as one heat cycle.

[0065] Furthermore, rough forging is performed over multiple heat cycles. This is because, in large ingots produced in mass production, the surface temperature drops during the process, causing cracks to form, and maintaining a high temperature above the β transformation point for an extended period allows for a more uniform elemental distribution. Preferably, there are four or more heat cycles, and even more preferably, six or more. It is necessary to perform processing that reduces the cross-sectional area by 20% or more in each heat cycle. However, if upsetting is performed in the same heat cycle, the cross-sectional area reduction may be less than 20%. The reason for requiring a cross-sectional area reduction of 20% or more is that repeatedly performing processing that reduces the cross-sectional area by less than 20% in large ingots significantly reduces productivity.

[0066] Next, finish forging is performed. After rough forging, the material may be cooled to near room temperature, or it may be immediately reheated and finish forging performed. Alternatively, a step to remove surface scale and wrinkles may be included before finish forging.

[0067] Finish forging is the process of shaping the material for hot rolling, and the finished size varies depending on the size of the manufactured product. In the case of bar stock, considering the yield reduction due to oxidation, the cross-sectional area should be 2000 mm². 2 That concludes the explanation. There is no specific upper limit, but the cross-sectional area must be 18,000 mm². 2 The following are preferred. The shape may be any shape that can be rolled, such as an octagon or square with a cross-sectional area corresponding to these. In the case of a plate shape, it should be rectangular, and the thickness of the plate should be 50 mm or more. Preferably it should be 70 mm or more, and even more preferably 100 mm or more. The upper limit of the plate thickness and the width size may be any size that can be rolled, and are not particularly limited.

[0068] The heating temperature for finish forging is the same as for rough forging, preferably 50°C, more preferably 100°C or more, higher than the β transformation point. Four or more heats are required, preferably six or more. However, there is no lower limit to the reduction rate of the cross-sectional area per heat, and heats may consist of heating only without reduction. In the finish forging process, forging is performed only by reducing the cross-sectional area from the side, without upsetting.

[0069] (Hot rolling) After finish forging, titanium alloys, particularly titanium alloy rods and titanium alloy sheets, are formed by hot rolling. After finish forging, the material may be cooled to near room temperature, or it may be immediately reheated and hot rolling performed. Alternatively, a process to remove surface scale and wrinkles may be included before hot rolling.

[0070] Hot rolling is performed by heating to a temperature between the β transformation point and 1200°C. Preferably, the temperature is 50°C higher than the β transformation point, and more preferably 100°C or more higher. The upper temperature limit is restricted from the standpoint of yield, as in forging. A preferred upper limit is 1150°C.

[0071] Hot rolling is preferable if the cross-sectional reduction rate is 70% or more per heat, as this reduces the maximum area of ​​segregation abnormalities and the proportion of needle-like structures. Preferably, it is 80% or more, and more preferably 85% or more. There is no particular upper limit, but it is approximately 99% in relation to the product diameter. When manufacturing thin products by hot rolling, the yield decreases significantly due to subsequent descaling and other processes. Therefore, for bar stock, the lower limit of the product diameter of hot-rolled material is set at φ10 mm. For plate stock, the lower limit of the product plate thickness of hot-rolled material is set at 3 mm.

[0072] During manufacturing, if the β phase is held at around 300°C in an unstable state, the ω phase may form significantly. However, in the production of large ingots, the distribution of alloying elements between the α and β phases occurs during hot rolling, stabilizing the β phase, so the ω phase hardly forms. If the presence of the ω phase is still a concern, it can be eliminated by annealing after hot rolling.

[0073] (Annealing and descaling) As mentioned above, annealing of titanium alloy rods and sheets after hot rolling is a necessary process when precipitation of the ω phase is a concern or has been confirmed, but it may also be performed for other purposes such as removing residual stress or strain. Furthermore, in applications where the operating environment is 200-400°C, the material becomes brittle due to precipitation of the ω phase during use, so annealing can suppress the precipitation of the ω phase during use. When annealing is performed, it should be done at 750°C or higher, but below the β transformation point. Higher temperatures are desirable in terms of elemental uniformity. More preferably, it should be 780°C or higher, and even more preferably, 800°C or higher. In order to keep the proportion of needle-like structures to 80% or less, it is preferable not to perform annealing, or if annealing is performed, to do so at a temperature below the β transformation point.

[0074] Scale formed during annealing or hot rolling can be removed by mechanical and chemical methods such as shot blasting, pickling, and polishing.

[0075] Titanium alloy rods and titanium alloy plates As mentioned above, when titanium alloy is in the form of a rod, such as a round bar, the reduction in cross-sectional area during hot working is smaller compared to a plate, so fracture due to segregation of alloying elements becomes apparent. In contrast, the titanium alloy rod and titanium alloy round bar of the present invention have a small maximum area of ​​segregation abnormalities and can achieve excellent fatigue properties. As a titanium alloy rod, the cross-sectional area is 80 mm². 2 ~8000mm 2 Materials within this range can be suitably used. On the other hand, in the case of plate material, the thickness is 3 mm to 50 mm. Preferably it is 30 mm or less, and more preferably 10 mm or less.

[0076] Engine valves As mentioned above, engine valves made of titanium alloy are manufactured using titanium alloy rods as the raw material. Conventional engine valves made of titanium alloy had a large maximum area of ​​segregation abnormalities, but the engine valve of the present invention uses the titanium alloy of the present invention, so the maximum area of ​​segregation abnormalities is small, and excellent fatigue characteristics can be achieved. [Examples]

[0077] For varieties No. A to F having the component compositions shown in Table 1, VAR ingots (1.8 tons) with a diameter of φ720 × 1000 mm were manufactured. For one example (No. 31 in Table 3), the weight was 0.2 tons. Table 2 describes the manufacturing conditions for manufacturing methods No. 1 to 15. As shown in Table 3, using ingots with the components of varieties No. A to F and the weight shown in "Ingot / ton" in Table 3, rough forging, finish forging, and hot rolling were performed under the conditions described for manufacturing methods No. 1 to 15 in Table 2 to produce round bars with the diameter shown in "Hot Rolling / Diameter / After Rolling" in Table 2. In Table 2, values ​​and items that fall outside the preferred range of the present invention are underlined.

[0078] [Table 1]

[0079] [Table 2]

[0080] For rough forging, as indicated in Table 2 under "Rough Forging / Heat No.", rough forging was performed for Heats 1 to 4 with the reduction ratios listed under "Rough Forging / Area Reduction (%)". In cases where "Yes" is indicated in the "Rough Forging / Upsetting" column, upsetting with a reduction ratio of 20% was performed in Heat 1 of rough forging. Nos. 10, 13-15 did not undergo upsetting. The reduction ratios listed for Heat 1 of rough forging take into account the effect of upsetting.

[0081] After rough forging, the material was cooled to room temperature and machined to 5mm / face to create a 240mm square cross-section. After rough forging, the material was divided lengthwise and subjected to finish forging and subsequent processes under various conditions.

[0082] For the finish forging, the material was heated to the temperature indicated in the "Finish Forging / Heating Temperature" column of Table 2, and for heats 1 to 4, the finish forging was performed at the reduction rate indicated in "Finish Forging / Reduced Surface Area (%)" as shown in "Finish Forging / Heat No." of Table 2. For pieces No. 1-7 and 11-15, half the length was forged, and the remaining half was forged after reheating. Therefore, the entire length had a uniform cross-sectional shape after two heats. For example, in piece No. 1, the front half was forged with a 65% reduction in cross-sectional area in the first heat, and the back half was not forged. In the second heat, the front half was not forged, and the back half was forged with a 65% reduction in cross-sectional area. Furthermore, in the third heat, the front half was forged with a 44% reduction in cross-sectional area, and the back half was not forged. In the fourth heat, the front half was not forged, and the back half was forged with a 44% reduction in cross-sectional area. For parts No. 8-10, the entire length was forged in a single heat. For part No. 8, the entire length was forged in the first heat with a 65% reduction in cross-sectional area, and the entire length was forged in the second heat with a 44% reduction in cross-sectional area. For part No. 9, the cross-sectional area was forged in the first to fourth heats with reductions of 45%, 36%, 23%, and 27%, respectively.

[0083] After finish forging, the material was cooled to room temperature, the top 5 mm was machined off, and hot rolling was performed under the conditions described in the "Hot Rolling" column of Table 2. The hot-rolled round bars were then heat-treated under the heat treatment conditions described in the "Annealing" column of Table 2. Examples marked with "-" in the "Annealing" column indicate that annealing was not performed.

[0084] The round bars manufactured as described above were subjected to EPMA analysis, X-ray diffraction, EBSD measurement, tensile testing, and fatigue testing. The quality evaluation results are shown in Table 3. In Table 3, values ​​that fall outside the scope of the present invention are underlined.

[0085] The elemental distribution for evaluating segregation anomalies was measured by EPMA on a mirror-polished cross-section perpendicular to the axis of a round bar. The EPMA measurement was performed at an acceleration voltage of 10kV, in a 5mm square area, at the center of the cross-section, with a step size of 0.02mm and a beam diameter of 0.02mm. In the Ti-5Al-2Fe-3Mo titanium alloy, in EPMA surface analysis, clusters of adjacent measurement points where Al and Fe are outside the mean ±1 mass%, and Mo is outside the mean ±1.5 mass%, were defined as areas of elemental segregation anomalies in Al, Fe, and Mo in the cross-section. The maximum area of ​​these elemental segregation anomalies in Al, Fe, and Mo is shown in Table 3. In Ti-6Al-4V titanium alloys, in EPMA surface analysis, clusters of adjacent points where Al is outside the mean ±1 mass% and V is outside the mean ±2 mass% were defined as areas of elemental segregation abnormalities in Al and V within the cross-section. The maximum area of ​​these areas of elemental segregation abnormalities in Al and V is shown in Table 3. For α+β type titanium alloys other than the two-component systems described above, the clusters of adjacent measurement points where the average value ±1 mass% of the main elements Al and Fe were deviated from were defined as areas of segregation abnormalities in the elemental concentrations of Al and Fe in the cross-section, respectively.

[0086] In X-ray diffraction, measurements were taken using Cu-Kα rays in the range of 2θ from 30 to 90° with a scan step of 0.01° to confirm the presence or absence of the ω phase, and the results are shown in Table 3.

[0087] EBSD measurements were performed on a cross section perpendicular to the cylindrical rod axis, polished with colloidal silica, using an acceleration voltage of 15kV. Three fields of view were used: the center of the cross section, the position corresponding to the center of the line segment connecting the radial surface and the center (R / 2 position), and the position obtained by rotating the R / 2 position 90° clockwise in the circumferential direction with respect to the center of the cylindrical rod axis. The measurement conditions were a magnification of 500x, a 200μm square area, and a step size of 0.2μm. After measurement, grain size was determined with an orientation difference of 5° or more as the grain boundary. If the average grain size was 100μm or more, the proportion of needle-like structures was set to 100 area %. If the average grain size was less than 100μm, data with an Image Quality (IQ) of 20% or more of the average was extracted. From the extracted data, α grains with an aspect ratio (long axis / short axis) of 3 or less were defined as equiaxed α grains, and the area ratio of grains other than the aforementioned equiaxed α grains to the total measurement area was defined as the proportion of needle-like structures, which is listed as "Needle-like Structure Ratio" in Table 3.

[0088] Tensile tests were performed using a specimen of a round bar with a parallel section diameter of φ6.25 mm and a parallel section length of 28 mm, with a gauge length of 25 mm. The test was conducted at a strain of 0.4% / min up to 2% and then at 25% / min until fracture, and the elongation at fracture was measured using the butt joint method.

[0089] For the fatigue test, a rotating bending fatigue test specimen with a parallel section of φ6 mm was prepared, and the surface of the parallel section was finished with #1000 emery paper in the circumferential direction of the parallel section. Using this specimen, a rotating bending fatigue test was performed with a stress amplitude of either 560 or 640 MPa. The stress amplitude was selected to be within the range of 0.54 to 0.56 times the tensile strength. The stress amplitude was also set by changing it in 20 MPa increments until the predetermined stress amplitude was reached. If multiple stresses fell within the range of 0.54 to 0.56 times the tensile strength, the test was performed with the higher stress amplitude. Within this range, the stress level is slightly lower than the fatigue limit, making it possible to determine whether the specimen was good or bad. The repeating speed was 50-60 Hz. Under these test conditions, if the specimen did not break, 1 × 10⁻⁶ 7 If the required number of repetitions was reached, it was considered a pass and was marked with a circle (○). If the number was not reached, it was marked with a cross (×) and recorded in the "Fatigue Characteristics" column of Table 3.

[0090] [Table 3]

[0091] Examples No. 1 to 14 of the present invention have the component compositions of the present invention shown in Table 1, and as a result of manufacturing under the preferred manufacturing conditions of the present invention as shown in Table 2, the maximum area of ​​segregation abnormalities for all alloying elements was within the range of the present invention as shown in Table 3, and the fatigue properties were good. In Example No. 4 of the present invention (manufacturing method No. 4), although the proportion of needle-like structures (needle-like structure ratio) was outside the preferred range of the present invention because the annealing temperature after hot rolling was outside the upper limit of the preferred range (β transformation point temperature), the fatigue properties were good.

[0092] Comparative Examples No. 15 to 31 are comparative examples. Comparative Examples No. 15 and 23 (Manufacturing Method No. 8) involved only two heat cycles in the finishing forging process. As a result of fewer processing cycles and fewer times the material was held at high temperatures, elemental diffusion was insufficient, and the maximum area of ​​the segregation abnormality fell outside the scope of the present invention. In Comparative Examples No. 16 and 24 (Manufacturing Method No. 9), the heating temperature during the finish forging was outside the lower limit of the preferred range. In Comparative Examples No. 17 and 25 (Manufacturing Method No. 10) and Nos. 20-22 and 28-31 (Manufacturing Methods 13-15), upsetting was not performed during rough rolling. In Comparative Examples No. 18 and 26 (Manufacturing Method No. 11), the heating temperature during hot rolling was outside the lower limit of the preferred range (β transformation point + 50°C). In Comparative Examples No. 19 and 27 (Manufacturing Method No. 12), the reduction ratio during hot rolling was outside the lower limit of the preferred range (70%). Therefore, in all of Nos. 16-22 and 24-31 (Manufacturing Methods No. 9-15), the maximum area of ​​any segregation abnormality of any alloying element was outside the range of the present invention. As a result, Comparative Examples No. 15 to 31 (Manufacturing Methods No. 8 to 15) all exhibited poor fatigue characteristics.

[0093] Comparative Example No. 19 (Manufacturing Method No. 12) had a large diameter after hot rolling and a low reduction ratio during hot rolling, which affected the cooling conditions after hot rolling, and the proportion of needle-like structures was outside the preferred range of the present invention.

[0094] Comparative Examples No. 15 and 23 (Manufacturing Method No. 8) showed an increase in the proportion of needle-like structures because the unstable β phase formed by annealing transformed into the α phase during cooling, thus falling outside the preferred range of the present invention.

[0095] Comparative Example No. 31 had a casting weight of 0.2 tons, which is less than 1 ton. However, the manufacturing method was Manufacturing Method No. 14, which falls outside the preferred range of the present invention. As a result, the maximum area of ​​the segregation abnormality was outside the range of the present invention, and the fatigue characteristics were also poor. [Examples]

[0096] For varieties No. G and H having the component compositions shown in Table 4, VAR ingots (1.8 tons) with dimensions of φ720 × 1000 mm were manufactured. Table 5 shows the manufacturing conditions for manufacturing methods No. 16 and 17. The ingots were subjected to rough forging, finish forging, and hot rolling under the conditions described for manufacturing methods No. 16 and 17 in Table 5 to obtain the sheet material shown in "Hot Rolling / Sheet Thickness / After Rolling" in Table 5. In Table 5, values ​​and items that fall outside the preferred range of the present invention are underlined.

[0097] [Table 4]

[0098] [Table 5]

[0099] [Table 6]

[0100] For rough forging, as shown in Table 5 under "Rough Forging / Heat No.", rough forging was performed for Heats 1 to 4 with the reduction ratios listed under "Rough Forging / Area Reduction (%)". In cases where "Yes" is indicated in the "Rough Forging / Upsetting" column, upsetting with a reduction ratio of 20% was performed in Heat 1 of rough forging. No. 17 did not undergo upsetting. The reduction ratios listed for Heat 1 of rough forging take into account the effect of upsetting. After rough forging, the material was cooled to room temperature and then machined to a thickness of 5 mm per surface, resulting in a cross-section of 290 mm x 490 mm. After rough forging, it was divided along its length.

[0101] For the finish forging, the material was heated to the temperature indicated in the "Finish Forging / Heating Temperature" column of Table 5. As indicated in the "Finish Forging / Heat No." column of Table 5, finish forging was performed for Heats 1 to 4 with the reduction ratio indicated in "Finish Forging / Reduced Surface Area (%)" to achieve a width of 500 mm. Similar to manufacturing methods No. 1-7 in Example 1, half the length was forged, and the remainder was forged after reheating. Therefore, a uniform cross-sectional shape was achieved throughout the entire length after two heat treatments.

[0102] After finish forging, the material was cooled to room temperature, the top 5 mm was cut off, and hot rolling was performed under the conditions described in the "Hot Rolling" column of Table 5. The hot-rolled sheet material was not annealed.

[0103] Various evaluations were performed in the same manner as in Example 1. For X-ray diffraction, measurements were taken in a cross section perpendicular to the rolling direction at the center of the plate thickness at the center of the width. For tensile and fatigue tests, a rectangular bar was cut from the center of the plate thickness at the center of the width so that the rolling direction was the longitudinal direction, and the same test specimens were processed and tested. For EPMA surface analysis samples for segregation anomaly evaluation, the cross section was perpendicular to the rolling direction. For EBSD measurement, a total of three fields were measured in a cross section perpendicular to the rolling direction: one field of view at the center of the plate thickness, and two fields of view at 1 / 4 width and 3 / 4 width.

[0104] The round bars manufactured as described above were subjected to EPMA analysis, X-ray diffraction, EBSD measurement, tensile testing, and fatigue testing, as in Example 1. The quality evaluation results are shown in Table 6. In Table 6, values ​​that fall outside the scope of the present invention are underlined. Example No. 32 of the present invention has the component composition of the present invention shown in variety No. G in Table 4, and as a result of manufacturing under the preferred manufacturing conditions of the present invention as shown in manufacturing method No. 16 in Table 5, the maximum area of ​​segregation abnormalities for all alloying elements was within the range of the present invention, as shown in No. 32 in Table 6, and the fatigue properties were good. Comparative Example No. 33 (Manufacturing Method No. 17) was made using rough forging without upsetting, and the maximum area of ​​the Fe and Mo segregation abnormalities fell outside the scope of the present invention. As a result, the fatigue properties were poor.

Claims

1. In mass percent, Al: 4.5-6.5%, Fe: 1.4-2.3%, Mo: Contains 1.5-5.5%, The total amount of O and N is 0.25% or less, and the remainder consists of Ti and impurities. A titanium alloy having a two-phase structure consisting of α-phase and β-phase, and in which the maximum area of ​​the segregation abnormality of the main element concentration in the cross-section is 2000 μm² or less. Here, the major elements are Al, Fe, and Mo, and the segregation anomaly in the concentration of major elements in the cross-section refers to a collection of adjacent points in EPMA surface analysis where the concentration of major elements other than Mo is outside the average value ±1 mass%, and the concentration of Mo is outside the average value ±1.5 mass%.

2. In mass percent, Al: 5.0-7.0%, V: 3.5-5.0%, The total amount of O and N is 0.25% or less, and the remainder consists of Ti and impurities. A titanium alloy having a two-phase structure consisting of α-phase and β-phase, and in which the maximum area of ​​the segregation abnormality of the main element concentration in the cross-section is 2000 μm² or less. Here, the major elements are Al and V, and the segregation anomaly in the concentration of the major elements in the cross-section refers to a collection of adjacent measurement points in EPMA surface analysis where the concentration of major elements other than V is outside the mean ±1 mass%, and V is outside the mean ±2%.

3. In mass percent, Al: 4.5-6.5%, Fe: 0.8-2.3%, Si: 0.0 to 0.50%, The total amount of O and N is 0.25% or less, and the remainder consists of Ti and impurities. A titanium alloy having a two-phase structure consisting of α-phase and β-phase, and in which the maximum area of ​​the segregation abnormality of the main element concentration in the cross-section is 2000 μm² or less. Here, the major elements are Al and Fe, and the segregation anomaly in the concentration of the major elements in the cross-section refers to a collection of adjacent points in EPMA surface analysis where the concentration of the major elements falls outside the mean ±1 mass% range.

4. A titanium alloy rod made of the titanium alloy described in any one of claims 1 to 3.

5. The titanium alloy rod according to claim 4, wherein the proportion of needle-like structures is 80 area % or less. Here, the proportion of needle-like structures is determined by performing EBSD evaluation on a cross section perpendicular to the axis of the titanium alloy rod, determining the grain size with an orientation difference of 5° or more as the grain boundary, and setting the proportion of needle-like structures as 100 area % if the average grain size is 100 μm or more. If the average grain size is less than 100 μm, data with an Image Quality (IQ) of 20% or more of the average is extracted, and α grains with an aspect ratio (long axis / short axis) of 3 or less are defined as equiaxed α grains, and the area ratio of grains other than the aforementioned equiaxed α grains to the total measured area is defined as the proportion of needle-like structures.

6. A titanium alloy plate made of the titanium alloy described in any one of claims 1 to 3.

7. The titanium alloy plate according to claim 6, wherein the proportion of needle-like structures is 80 area % or less. Here, the proportion of needle-like structures is determined by performing EBSD evaluation on a cross section perpendicular to the rolling direction of the titanium alloy sheet, determining the grain size with an orientation difference of 5° or more as the grain boundary, and setting the proportion of needle-like structures as 100 area % if the average grain size is 100 μm or more. If the average grain size is less than 100 μm, data with an Image Quality (IQ) of 20% or more of the average is extracted, and α grains with an aspect ratio (long axis / short axis) of 3 or less are defined as equiaxed α grains, and the area ratio of grains other than the equiaxed α grains to the total measured area is defined as the proportion of needle-like structures.

8. An engine valve made of a titanium alloy according to any one of claims 1 to 3.

9. The engine valve according to claim 8, wherein the proportion of needle-like tissue is 80 area % or less. Here, the proportion of needle-like structures is determined by performing EBSD evaluation on a cross section perpendicular to the engine valve axis, determining the grain size with an orientation difference of 5° or more as the grain boundary, and setting the proportion of needle-like structures as 100 area % if the average grain size is 100 μm or more. If the average grain size is less than 100 μm, data with an Image Quality (IQ) of 20% or more of the average is extracted, and α grains with an aspect ratio (long axis / short axis) of 3 or less are defined as equiaxed α grains, and the area ratio of grains other than the equiaxed α grains to the total measured area is defined as the proportion of needle-like structures.

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