JPH07150316A - Manufacture of (alpha+beta) type ti alloy forged material - Google Patents

Abstract

PURPOSE:To manufacture a Ti alloy forged material excellent in both strength and ductility and also in homogeneity that is necessary for thick parts such as a turbine blade. CONSTITUTION:This manufacturing method of an (alpha+beta) type Ti alloy forged material is composed of stages in which, after the (alpha+beta) type Ti alloy consisting of 3.0-5.0% Al, 2.1-3.7% V, 0.85-3.15% Mo, 0.85-3.15% Fe, 0.06-0.20% O (wt.% in all case) and the blance Ti with inevitable impurities is heated to a temp. in the (alpha+beta) region and forged, the forged material heated and held at T deg.C of Tbeta-75<=T<Tbeta deg.C (Tbeta: beta transformation point, deg.C) and successively cooled. After the above heating and forging, the solution heat treatment is executed at T deg.C of Tbeta-100<=Tbeta-25 deg.C (Tbeta: beta transformation point, deg.C) and successively aging treatment is executed in the temp. range of 400-600 deg.C.

Description

Detailed Description of the Invention

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a method for producing a Ti alloy forged material, and more particularly to a method for producing an (α + β) type Ti alloy forged material which has both excellent strength and ductility and high homogeneity.

[0002]

2. Description of the Related Art The Ti alloy is widely used as a material for aerospace equipment because of its high specific strength, but in recent years, various equipment parts have become large in size and the need for thick parts has increased. In the case of Ti alloy, thick parts are manufactured by the die forging method, but there is a problem that a high strength member cannot be obtained and a homogeneous member cannot be obtained.

Strength Problems of Thick Parts (above): Ti-6
The (α + β) -type Ti alloy represented by the Al-4V alloy is a Ti alloy that has excellent mechanical properties such as strength and ductility as well as excellent weldability and workability, and is widely used. Although variously used as a forging material, since high strength has been achieved by solution aging, high strength in thick parts has not yet been realized due to restrictions of cooling conditions after solution heat treatment. Conventional (α + β)
This is because, in the case of type alloys, in order to increase the strength, it is necessary to perform cooling with a high cooling rate such as water cooling after solution heat treatment, while for thick parts, it is difficult to rapidly cool to the center of the parts even by water cooling. .

Problem of homogeneity of die forging material (above): In the die forging material, internal variations occur in the microstructure, strain distribution and mechanical properties of the member after forging due to the following reasons. (1) When a heated material is set in a die and forged, heat is removed through the die or atmosphere in contact with the die. Therefore, the temperature of the material cannot be kept uniform, and the strain distribution and microstructure after forging become non-uniform. This problem becomes particularly noticeable when the forging process is complicated and the time from the start of forging to the end of forging greatly differs depending on the part.

(2) The distribution of processing strain caused by forging largely depends on the shape of parts. As the shape becomes more complex, the degree of heterogeneity in microstructure and mechanical properties also increases. (3) The temperature of the material during forging rises remarkably due to the heat generated during processing. Although the amount of heat generation varies depending on the temperature of the material itself and the amount of processing, in combination with the problems (1) and (2), each part of the heavy-weight component undergoes a thermal history that differs significantly depending on its shape and site.

Various solutions to the above problems have been proposed so far, but regarding the problem of strength (), in order to avoid solution treatment and subsequent quenching treatment, (α + β) type Instead of, use β-type Ti alloy
Japanese Patent Application Laid-Open No. 4-63239 discloses a method of obtaining a Ti alloy having high strength by forging in a temperature range of 650 to 1200 ° C. and then directly aging in a temperature range of 350 to 650 ° C.

However, in the aging treatment shown here, since the aging temperature is low, the work strain generated during forging and the inhomogeneity of the microstructure inside the member remain in the material after aging, and the mechanical properties depend on the part. We cannot solve the problem of variation.

Further, Japanese Patent Laid-Open No. 5-59510 discloses an alloy having the same composition as the alloy used in the present invention (β transformation point-
After heating to a temperature in the range of 150 ℃) -β transformation point, 0.5-
There is disclosed a technique for obtaining a high tensile strength of 1105 MPa or more by cooling at a cooling rate of 10 ° C./sec, performing solution treatment, and further performing aging treatment at a temperature in the range of 400 to 600 ° C. However, the technique disclosed here is not related to the forged material, and therefore, there is no mention of controlling the variation in mechanical properties peculiar to the forged material, which is a problem in the present invention.

Regarding the problem () of variation in mechanical properties depending on parts, Japanese Patent Laid-Open No. 63-157846 discloses Ti-6.
After forging the Al-4V alloy (α + β), the (β transformation point −18
There is disclosed a technique for obtaining a forged member having a small internal variation by annealing in a temperature range of (0 ° C) to β transformation point and cooling at a cooling rate of air cooling or less. As an example, a Ti-6Al-4V alloy having an oxygen content of 0.20% (% by weight, the same applies hereinafter) has variations due to (α + β) forging followed by annealing at 900 ° C. for 2 hours and subsequent furnace cooling. It is described that a small and high strength forged member was obtained.

However, comparing the numerical values of the examples with the numerical values of the comparative examples, although the variation itself is certainly small, the minimum value of the examples is the same as the minimum value of the comparative examples in terms of both strength and ductility. It is lower than that, and the minimum value of elongation is 12.4%, and it cannot be said that the mechanical properties are rather deteriorated when viewed as the entire member.

Further, in the above embodiment, the tensile strength (928) of the (α + β) type forged Ti alloy is relatively high as a forging material.
.About.956 MPa), which is due to the high oxygen content of the working alloys. According to AMS standard 4928K, 0.
The value of 20% corresponds to the maximum allowable value of oxygen, and the amount above this is considered to be out of specification due to the relationship with ductility. Incidentally, the tensile strength when the oxygen content is suppressed to about 0.08% in order to secure sufficient ductility is actually about 900 to 930 MPa under the same recrystallization annealing conditions as in the above embodiment. (For example, AEROSPACE STRUCTURAL METALS HANDBOOK, Vol.
4, code 3707, p.10, Army Materials and Mechanics R
(eserch Center, USA, 1980)

[0012]

As described above, Ti
Despite various studies on alloy composition and heat treatment conditions for alloys and their forged materials, while maintaining the original strength and ductility of each alloy composition, at the same time suppressed internal variations in microstructure and mechanical properties to a small extent. A technique for obtaining a thick forged material with high homogeneity has not been known yet.

In view of the present situation of the forging material manufacturing technology for Ti alloys, the present invention has a sufficiently high strength even when the cooling rate inside the member is slow in cooling after heat treatment like a thick forging material. Of the Ti alloy forgings, which can achieve high temperature homogeneity in mechanical properties by controlling the processing strain and microstructure, especially the optimum alloy composition and heat treatment conditions. The purpose is to clarify.

From an industrial point of view, it is needless to say that it is desirable that high strength and homogenization can be achieved at the same time by a simple heat treatment such as annealing instead of a two-step heat treatment such as solution aging treatment. .

[0015]

Means and Actions for Solving the Problems The inventors of the present application have conducted extensive studies to solve these problems, and as a result, a part of the alloy composition described in JP-A-5-59510 is forged. Is an optimal composition for simultaneously solving the above two problems of ensuring strength and ductility and controlling internal variation of mechanical properties. If an alloy having this composition is subjected to an appropriate heat treatment, Ti having the desired material properties can be obtained. It has been found that an alloy forged material can be obtained.

That is, the first invention of the present application relates to a material to be forged.
Al: 3.0-5.0%, V: 2.1-3.7%, Mo: 0.85-
An (α + β) -type Ti alloy containing 3.15%, Fe: 0.85 to 3.15%, O: 0.06 to 0.20% (all by weight, the same applies hereinafter) and consisting of the balance Ti and unavoidable impurities is used. ) After heating to the region temperature and forging, Tβ-75
(Α + β) type T consisting of a process of heating and holding at a temperature T ° C. where T ° C. ≦ T <T β (T β: β transformation point, ° C.) and then cooling.
A method for producing an i alloy forged material, which enables production of an (α + β) type Ti alloy forged material having excellent strength and ductility and high homogeneity.

Here, the cooling following the heating and holding at T ° C. may be air cooling or water cooling, and the cooling rate is not particularly limited, but in order to simplify the expression, In the description of the first invention, this heat treatment is 1
The case of step annealing will be described.

On the other hand, in the second invention of the present application, an (α + β) type Ti alloy having the same composition range as that of the first invention is used as a material to be forged, and this is heated to a temperature in the (α + β) range for forging,
Tβ-100 ° C ≤ T ≤ Tβ-25 ° C (Tβ: β transformation point, ° C)
A method for producing an (α + β) type Ti alloy forged material, which comprises the steps of solution treatment at a temperature of T ° C., followed by aging treatment at a temperature range of 400 to 600 ° C., which has the same strength and ductility as the first invention. This enables the production of (α + β) type Ti alloy forgings with excellent and high homogeneity.

The cooling rate after solution treatment is as follows.
Similar to the cooling rate in the first invention, there is no particular limitation.
As typical compositions of the above alloys, Al: 4.5%, V: 3.0%, Mo: 2.0%, Fe:
An alloy composed of 2.0%, O: 0.08%, the balance Ti and inevitable impurities can be mentioned. The β transformation point of this representative composition is 900 ° C, and the recrystallization temperature is 800 ° C.

First of all, the problem of high strength, which is the first problem, will be described. In the case of Ti-6Al-4V alloy, the transformation from the β phase stable at high temperature to the α phase stable at low temperature occurs relatively easily. Therefore, when the cooling rate is slow like air cooling, during the cooling after solution treatment. There is a problem in that the pro-eutectoid α phase is coarsened and a large amount of new transformation precipitation of α phase occurs, and in the subsequent aging, it is not possible to generate sufficient α phase to contribute to the strength increase.

On the other hand, the alloy used in the present invention increases the stability of the β phase and suppresses the rapid generation / increase of the α phase after the solution treatment, and at the time of the aging treatment, an appropriate amount of the β phase is contained. By precipitating the α phase finely, a large difference can be obtained in that a high material strength can be obtained even under a condition where the cooling rate is slow such as air cooling. This solves the problem of cooling rate that is characteristic of thick parts.

This alloy is an alloy excellent in superplastic forming ability in the first place as described in JP-A-3-274238, but V, Mo and Fe which are β-phase stabilizing elements.
Is contained in a predetermined amount, and the stability of the β phase is extremely high. These β-type stabilizing elements have a function of suppressing rapid and large amount of α-phase precipitation in the cooling process after the solution treatment. Also, among these, Mo has a slow diffusion rate in Ti unlike other additive elements, has the effect of refining the entire structure, and prevents coarsening of the α phase precipitated in the β phase during the cooling process. . By these actions, the structure is miniaturized throughout, and high strength as a material is achieved.

Due to the above characteristics, unlike the Ti-6Al-4V alloy, in the alloy used in the present invention, it is possible to increase the strength of the forged material by the solution aging treatment (second).
invention). Furthermore, the solution aging treatment itself is omitted, and the heat treatment after forging is an extremely simple heat treatment of only annealing,
It also becomes possible to secure sufficient strength and ductility (first invention).

Next, the control of the internal variation of material characteristics, which is the second problem, will be described. Inside the thick forged material, inhomogeneity occurs in processing strain and microstructure, and as a result, mechanical properties vary depending on the part.
It is an unavoidable problem after going through the forging process step, and it is considered that there is no means for solving this variation other than removal by performing heat treatment in the post process of forging.

With respect to this point, the first invention of the present application solves this problem by a simple heat treatment in which only the heat treatment temperature of the material to be forged is limited to the range of Tβ-75 ° C. ≦ T <Tβ, for example, annealing. That is, by heat-treating an alloy having the above-mentioned composition in a specific temperature range above the recrystallization temperature, the processing strain is released, and at the same time, the microstructure is controlled to reduce the variation in mechanical properties inside the member. .

In the second invention, the forged material is Tβ-100.
As in the case of the first invention, the solution treatment is carried out in the limited temperature range of ℃ ≤ T ≤ Tβ-25 ℃, and the aging treatment is subsequently carried out in the temperature range of 400 to 600 ℃, while aiming at high strength. It has realized the reduction of variations in mechanical properties inside the members.

In the first invention of the present application, even if the cooling rate after heating and holding at T ° C. is slow, there is no problem in controlling the processing strain and the microstructure. The same applies to the cooling rate after solution treatment in the second invention.
The fact that the alloy composition of the present invention plays a major role in enabling such heat treatment is a finding that, together with the large role of the alloy composition in achieving high strength of the material, is the core of the present invention. is there.

The components of the Ti alloy used in the present invention are limited to the above range, that is, the role of each component in the alloy is as follows. Al: A typical α-stabilizing element, which is an essential additional element for the (α + β) -type Ti alloy. If the Al content is less than 3.0%, it becomes difficult to form an (α + β) type alloy and sufficient strength cannot be obtained as a material. On the other hand, when the Al content exceeds 5.0%, Ti3Al, which is an intermetallic compound, is formed, and the toughness is remarkably reduced.
Therefore, the Al content is limited to 3.0 to 5.0%.

V: An important additional element that stabilizes the β phase and at the same time lowers the β transformation point. Suppresses the rapid formation and increase of α phase after annealing or solution treatment,
It has the effect of finely precipitating phases. If the V content is less than 2.1%, the β transformation point cannot be sufficiently lowered, and the effect of stabilizing the β phase becomes small, so that the effect of suppressing the formation of the α phase during annealing or after solution treatment is obtained. I can't get it.
On the other hand, if the V content exceeds 3.7%, the stability of the β phase becomes too large, and a preferable two-phase structure of (α + β) cannot be obtained, resulting in insufficient strength. Therefore,
The V content is limited to the range of 2.1 to 3.7%.

Mo: has the effect of stabilizing the β phase and, at the same time, suppressing grain growth. Therefore, like V, it is important not only for suppressing the rapid formation and increase of the α phase after annealing or solutionizing and finely precipitating the α phase, but also for the effect of refining the entire structure. And is an additive component that occupies an important position in enhancing strength. Mo content is 0.85
If it is less than%, the crystal grains become coarse during annealing or after solution treatment, and the desired effect described above cannot be obtained. On the other hand, when the Mo content exceeds 3.15%, the β phase is excessively stabilized and a preferable two-phase structure cannot be obtained, so that an increase in strength cannot be expected. Therefore, the Mo content is limited to the range of 0.85 to 3.15%.

Fe: An additive component that effectively increases the stability of the β phase. Combined with the effects of V and Mo, it suppresses the rapid formation and increase of the α phase after annealing or after solution treatment, and also precipitates fine acicular α phase in the β phase during the cooling stage.

If the Fe content is less than 0.85%, the stability of the β phase is not sufficient, and the formation and increase of the α phase cannot be suppressed during cooling after annealing or solution heat treatment. High strength cannot be achieved by chemical aging treatment. On the other hand, if the content exceeds 3.15%, a brittle intermetallic compound is formed between Fe and Ti, or a segregation phase called β-fleck is formed, and the mechanical properties of the alloy deteriorate. Therefore, the Fe content is 0.85 to 3.
Limited to 15% range.

O: Oxygen content is normal (α + β) type Ti
The amount is similar to that of the alloy. If the O content is less than 0.06%, a problem occurs in strength, and if it exceeds 0.20%, the ductility sharply decreases. Therefore, the O content is limited to the range of 0.06 to 0.20%.
From the standpoint of strength-ductility balance, 0.08-0.1%
It is desirable to limit the range.

Next, the reason for limiting the heat treatment conditions will be described by taking the case of the representative component as an example. The heat treatment of the first invention is as follows. First, as a result of investigating the relationship between the annealing temperature and the mechanical properties of the representative alloys used in the present invention, it was found that if the heat treatment was performed at a temperature equal to or higher than the recrystallization temperature, the variation in the mechanical properties due to the site could be removed. However, on the other hand, the strength is reduced by releasing the processing strain, and the diameter specified in AMS4928K
It may fall below the strength specification value (895 MPa) of 100 mm products. Therefore, in order to remove the processing strain and secure the strength, the heat treatment in the above-mentioned temperature range is required.

The β transformation point Tβ of this alloy is 900 ° C., but when the annealing temperature is changed in the range of 800 ° C. which is the recrystallization temperature to Tβ, β which exists in thermal equilibrium during the annealing process.
The volume fraction of the phase increases with increasing temperature. (Α + β)
The improvement of the strength in the type Ti alloy is borne by the needle-like α phase which is finely precipitated in the β phase during cooling, but if the volume fraction of the β phase during annealing increases, it will precipitate during cooling. The volume ratio of the fine acicular α-phase also increases, and high strength can be secured. Moreover, since the heat treatment is performed in the (α + β) 2 phase region, both α and β phases coexist in a balanced manner, so that the structure is kept fine.

Such an effect is obtained when the annealing temperature is (Tβ-75
It begins to appear at 825 ℃, which is ℃), and the strength rises up to just below the β transformation point at 900 ℃. Furthermore, when the annealing temperature is raised to a temperature exceeding Tβ, the entire surface becomes a β structure while the crystal grains are rapidly coarsened. Therefore, although the strength is maintained by the precipitation of the needle-shaped fine α-phase, the ductility sharply decreases. Therefore, the annealing temperature should be above (Tβ-75) ° C and below Tβ.

Next, the heat treatment of the second invention will be described. As a result of examining various conditions for strengthening by solution aging treatment, an alloy having the same composition as that of the first invention was subjected to solution treatment within a temperature range of recrystallization temperature to (β transformation point -25 ° C), and was continuously subjected to 400 It was found that high strength can be achieved by aging in the temperature range of ~ 600 ℃. In this case, if the solution treatment temperature is lower than the recrystallization temperature, internal variations remain in the mechanical properties after aging, while if it exceeds (β transformation point −25 ° C.), the ductility lowers, which is a problem. . If the aging temperature is less than 400 ° C, no increase in strength can be expected after aging, and if the aging temperature is more than 600 ° C, the strength is increased and the strength is immediately softened, making it difficult to control the strength. For the above reasons, the solution aging treatment conditions are limited to the predetermined temperature range.

As described above, (α + β) of the present invention
In the method for producing the type Ti alloy forged material, a simple heat treatment step such as only annealing can be selected, or a solution aging treatment can be selected for further strengthening.

As an important technical feature of the present invention, even if a thick forged material whose cooling rate is difficult to control is cooled at a relatively slow rate such as furnace cooling or cooling outside the furnace, sufficient strength and ductility as a member are obtained. Can be secured, and at the same time, it is possible to control internal variations. Therefore, there is no problem even if a cooling method having a high cooling rate such as water cooling is used for cooling after annealing or solution treatment, and the cooling rate is not particularly limited.

[0040]

【Example】

(Example 1) Al: 4.5%, V: 3.0%, Mo: 2.0
%, Fe: 2.0%, O: 0.08%, C: 0.02%, N: 0.01
%, H: 0.01%, the balance consisting of Ti (α +
β) type Ti alloy [β transformation point: 900 ℃, recrystallization temperature: 800
℃] ingot is heated to 1100 ℃ β region, forged,
A round billet with a diameter of 220 mm was obtained. This round billet (α
After heating to 800 ° C which is the + β) range and forging into a round billet with a diameter of 152 mm, it is cut at the center in the length direction to give a diameter of 15
Two round billets with a length of 2 mm and a length of 1200 mm were obtained. (Hereinafter, this one is called the # 1 billet and the other is called the # 2 billet).

Next, the # 1 billet is set to 800 in the (α + β) range.
It was heated to ℃, processed from four directions in a plane perpendicular to the length direction, forged from the top side to a length of 1/2, and the diameter of this portion was set to 94 mm. After that, the entire billet was returned to the heating furnace and heated again to 800 ° C.
Similarly, forging the length of 2 to 94 mm, the length is 3000 m
Finished billet with m and 94mm outer diameter.

The # 2 billet was also finished to be the same size as the # 1 billet, with a diameter of 94 mm and a length of 3000 mm. However,
In case of # 2, only one heat was applied, and after heating to 760 ° C., forging was performed in the order of 1/2 length from the top side and 1/2 length from the bottom side. The above billet manufacturing process is shown collectively in FIG.

As a comparative example, a typical (α + β) type Ti alloy, Ti-6Al-4V alloy, was forged by the same method to obtain a round billet (# 3 billet) having a diameter of 94 mm and a length of 3000 mm. It was The composition of the Ti-6Al-4V alloy used is
Al: 6.15%, V: 4.02%, Fe: 0.17%, O: 0.09
%, C: 0.02%, N: 0.01%, H: 0.01%, balance Ti
[Β transformation point: 1000 ° C, recrystallization temperature: 870 ° C]. T
In the case of the i-6 Al-4 V alloy, the β transformation point and the recrystallization temperature are different from those of the alloy used in the present invention, so the heating temperature was set to 930 ° C. Further, the number of heats was 1 heat, which is the same as that of the # 2 billet. The manufacturing process of this Ti-6Al-4V alloy is also shown in FIG.

A cylindrical test piece having a length of 70 mm was cut out from the top side and the bottom side of these three billets, annealed at various temperatures as shown in Table 1 for 2 hours, and then air-cooled.
At this time, a hole having a diameter of 2 mm was opened in the center of the test body, a thermocouple was inserted, and the cooling rate was measured. The cooling rate in the center was 0.5 ° C / sec.

[0045]

[Table 1]

Tensile test pieces having a parallel part diameter of 6.25 mm and a gauge length of 25 mm were cut out from the center of each of the test pieces thus obtained, and a room temperature tensile test was carried out under the conditions of a crosshead speed of 0.15 mm / min. . The test results are shown in Table 2 and FIG.

[0047]

[Table 2]

As can be seen from FIG. 2, in the example of the present invention, the strength-ductility balance is far superior to that in the comparative example of the Ti-6Al-4V alloy having almost the same oxygen content. Is shown. This is because the Ti alloy used in the present invention is more stable in the β phase than the Ti-6Al-4V alloy, and therefore, even when the cooling rate is slow as in this example, coarsening of the pro-eutectoid α phase during cooling is observed. In addition to being suppressed, Ti-6Al-4 in the β phase during the cooling process
This is because a fine needle-like α phase which is not seen in the V alloy is generated and the average grain size of the entire structure can be controlled to be small by the effect of Mo.

Also, from FIG. 3, when the annealing temperature is 800 ° C. or lower, large variations in mechanical properties are recognized depending on the sampling position of the cylindrical test body even under the same heat treatment conditions. This is because the strain energy stored during forging is different in each part of the billet due to the difference in thermal history, and this strain energy is not completely released even after annealing. On the other hand, almost no difference is recognized in the examples of the present invention, as long as the condition of annealing in a specific temperature range is satisfied, the strain introduced by forging is released regardless of other thermal history, This is because the mechanical properties of the member are determined only by the annealing temperature.

(Example 2) Under the same conditions as in Example 1, forged billets # 1 to # 3 were obtained. # 1 and # 2
From the top and bottom sides of the billet, a 70 mm long columnar test piece was cut out, solution-treated at various temperatures as shown in Table 3 (soaking for 1 hour), air-cooled, and subsequently kept at 510 ° C. After aging for 6 hours, a solution aging treatment of air cooling was performed. # 3
A similar cylindrical specimen was cut out from the billet, which was solution-treated (soaked for 1 hour) at 950 ° C and then water-cooled, followed by aging and cooling under the same conditions as # 1 and # 2 for solution-aging treatment. gave.

A tensile test piece was cut out from the central portion of each of the cylindrical test bodies subjected to the above-mentioned treatment in the same manner as in Example 1, and a normal temperature tensile test was carried out under the same conditions as in Example 1. The test results are shown in Tables 3 and 4 and FIG.

[0052]

[Table 3]

[0053]

[Table 4]

As can be seen from FIG. 4, in the case of the example of the present invention, high strength was achieved by the aging treatment, and compared with the comparative example of the Ti-6Al-4V alloy having almost the same oxygen content, it was far more. It exhibits excellent strength-ductility properties. This is because the Ti alloy of the example of the present invention has a more stable β phase than the Ti-6Al-4 V alloy, so that the coarsening of the pro-eutectoid α phase is suppressed even when the cooling rate is slow like air cooling. , Ti-6Al-4V again
This is because fine acicular α-phase which is not seen in the alloy is generated in β-phase, and the average grain size of the entire structure can be controlled to be small by the effect of Mo.

Further, in the comparative example, even if the solution aging treatment temperature is the same, the mechanical properties greatly vary depending on the site. This is probably because the difference in strain energy stored during forging is inherited even after aging.

[0056]

According to the present invention, it is possible to manufacture a (α + β) type Ti alloy forging having uniform mechanical properties over the entire member and having an excellent strength-ductility balance. Further, in the production, a simple heat treatment step such as only annealing can be selected, or a solution heat treatment aging treatment can be selected in order to increase the strength. The Ti forged material manufactured by this method can be widely used for thick parts such as turbine blades that require high reliability.

[Brief description of drawings]

FIG. 1 is a view showing a forging condition of a test material billet.

FIG. 2 is a diagram showing the relationship between annealing temperature and tensile strength and elongation.

FIG. 3 is a diagram showing a relationship between an annealing temperature and mechanical properties of various parts of a forged material.

FIG. 4 is a diagram showing a relationship between solution aging conditions and mechanical properties of each portion of the forged material.

 ─────────────────────────────────────────────────── ─── Continuation of the front page (72) Inventor Chiaki Ouchi 1-2-1, Marunouchi, Chiyoda-ku, Tokyo Nihon Steel Pipe Co., Ltd.

Claims (2)

[Claims] 1. Al: 3.0 to 5.0%, V: 2.1 to 3.7.
%, Mo: 0.85 to 3.15%, Fe: 0.85 to 3.15%, O: 0.
After heating an (α + β) type Ti alloy containing 06 to 0.20% (both by weight) of Ti and unavoidable impurities to the temperature in the (α + β) range and forging, Tβ-75 ° C ≤ T <Tβ (Tβ: β transformation point, ° C) A method for manufacturing an (α + β) type Ti alloy forged material, which comprises a step of heating and holding at a temperature T ° C of which the temperature is T ° C, and subsequently cooling. 2. An (α + β) type Ti alloy having the same composition range as that of the alloy according to claim 1 is heated to a temperature in the (α + β) range for forging, and then Tβ-100 ° C. ≤ T ≤ Tβ- A method for producing an (α + β) type Ti alloy forged material, which comprises a step of solution treatment at a temperature T ° C of 25 ° C (Tβ: β transformation point, ° C) and subsequent aging treatment in a temperature range of 400 to 600 ° C.