JP2729011B2 - TiAl-based intermetallic compound alloy having high strength and method for producing the same - Google Patents

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION The present invention utilizes the transformation heat treatment and superplastic deformation, TiAl-based intermetallic compound if having high strength
The present invention relates to gold and a method for producing the same, and is used for application to heat resistant structural members having high specific strength.

[0002]

2. Description of the Related Art The intermetallic compound TiAl, which is expected to be put to practical use as a heat-resistant material, is difficult to process because of poor ductility. Techniques for improving the low workability, which are the biggest obstacles to the practical use of TiAl, are roughly classified into the application of the working process and the alloy design. Low workability mainly refers to lack of ductility at room temperature, and TiAl cannot be directly applied at room temperature to conventional working methods such as rolling and forging.

[0003] In the case of applying a processing process, conventional rolling, from near net shaping represented by powder processing,
Including processing methods such as forging. High temperature sheath rolling (1373) using Co-based superalloy (S-816)
K, rolling speed: 1.5 m / min)
213361), 800 ° C. or higher, strain rate 10
^-2 sec ^-1 or less at constant temperature forging (JP-A-63-1718)
No. 62) has been reported. The feature of such a processing method is to utilize the development of the rolling performance of TiAl at 800 ° C. or higher.
Molding is enabled by using the strain rate dependency on the mechanical properties of. However, there is a drawback that the application of large-scale equipment is not always easy because the processing conditions for performing sufficient molding are a high temperature of 1273 K or more and the strain rate must be reduced as much as possible.

On the other hand, it has been reported that after mixing and compacting of Ti and Al, compacting by high-temperature and high-pressure treatment (Japanese Patent Application Laid-Open No. 63-163).
-140049). This method is different from the above-mentioned processing process, and has an advantage that it can be formed into various shapes simultaneously with molding. On the other hand, there is a problem in that impurities are inevitable due to the use of an active metal such as Ti or Al. Is pointed out.

[0005] On the other hand, reports on improvement of room temperature ductility by addition elements are described in Mn addition (Japanese Patent Application Laid-Open No. 61-41740) and Ag addition (Japanese Patent Application Laid-Open No.
No. 23847), and General Elec
tric Corp. Addition of Si (US Patent: 48
36983), and Ta addition (U.S. Pat.
7), Cr addition (US Pat. No. 4,842,819), and B addition (US Pat. No. 4,842,820). Among them, General Electric Corp. The ranges of the components of each alloy system of Si, Ta, Cr, and B are determined from the ductility evaluation by a four-point bending test. In each case, titanium is equivalent to aluminum or higher than aluminum. Further, in order to improve the high-temperature ductility, 0.1%
005 to 0.2 wt. % B (Japanese Unexamined Patent Publication No. 63-12563)
No. 4) or 0.02 to 0.3 wt. % B and 0.
2 to 5.0 wt. % Si as a composite (Japanese Patent Laid-Open No. 63-125)
634). Until now, there is only one patent example by composite addition, but in order to improve a plurality of characteristics, it is necessary to consider the fourth and fifth additional elements. That is, the effects of these additional elements require a wide range of alloy component adjustments, including improvement in oxidation resistance and improvement in creep resistance, in addition to improvement in rolling performance. It is said that the tensile elongation at room temperature is a value of 3.0%, but this has not yet been achieved by the component design method by selecting any of the added elements, and the microstructure control such as refining by combined use with the processing process. It is considered essential to respond through this.

[0006]

SUMMARY OF THE INVENTION An object of the present invention is to control the structure of a TiAl-based alloy by selecting a component system and processing conditions, and to design a material having superplastic deformability,
Forming is performed using the superplastic deformability of the design material to form a product close to the final product shape, and a high-strength product is produced by heat treatment using phase transformation.

[0007]

In order to achieve the above object, the present inventors have developed basic mechanical properties of a multi-component TiAl-based intermetallic compound alloy, mechanical properties of a structure controlling material obtained by processing recrystallization, and the present invention. As a result of an experimental and theoretical analysis of the phase stability of the constituent phases that strongly affect the mechanical properties of the material, the following means for solving the problems were found as effective methods. In other words, the target microstructure control is not limited to microstructure refinement by simple recrystallization, but the β phase expected to be a metastable phase is precipitated at
The first step is to form a + γ two-phase structure, relax the introduced strain to a β phase that is rich in deformability, and apply superplastic deformability without impairing the excellent strength of TiAl.
Then the second step, strength properties, creep characteristics, hydrogen embrittlement and to improve the oxidation resistance, by utilizing this structure control phase transformation superplastic deformation material, to gamma + alpha ₂ dual phase structure. And, by establishing a process for controlling the structure of the working structure incorporating this series of processes,
This is a solution to the above problem. The details will be described below. The TiAl-based intermetallic alloy of the present invention has a composition
Is expressed by the following formula in atomic fraction, and β precipitated at the γ grain boundary
The phase has a volume fraction of 2 to 25% and has superplastic deformability
Β + γ TiAl-based intermetallic compound produced by transformation heat treatment
Alloy at room temperature to 1073 K
Composed of α ₂ + γ two-phase structure with a strength of more than 00MPa
Has high strength. Ti _a Al _100-abc Cr _b X _c X: one or more of Nb, Mo, Hf, Ta, W, provided that 47.5 ≦ a ≦ 52 1 ≦ b ≦ 5 0.5 ≦ c ≦ 3 b ≧ c 2a + b + c ≧ 100 The other TiAl-based intermetallic alloys of the present invention have the original composition.
It is expressed by the following formula in the subfraction. _{Ti a Al 100-abd Cr b} Y d Y: Si, B of one, two or where 47.5 ≦ a ≦ 52 1 ≦ b ≦ 5 0.1 ≦ d ≦ 2 2a + b + yet another TiAl of d ≧ 100 invention The base intermetallic compound alloy has the composition
Is represented by the following formula in atomic fraction. _{Ti a Al 100-abcd Cr b} X c Y d X: Nb, Mo, Hf, Ta, W, one of V or two or more Y: Si, one or two provided that 47.5 ≦ a ≦ 52 1 of B ≦ b ≦ 5 0.5 ≦ c ≦ 3 b ≧ c 0.1 ≦ d ≦ 2 2a + b + c + d ≧ 100

[0008] The precipitation of the grain boundary β phase is the superplasticity of the first stage described above.
This is an absolute condition for imparting deformability. With the third additive element
And the β-stabilizing elements Mo, V, Nb,
As a result of selecting six types of Fe and Mn and controlling the structure,
Only Cr was able to observe a clear grain boundary precipitation phase.
Therefore, Cr was selected as the third additive element. Sa
In addition, the TiAlC
r To compensate for the lack of strength that was a disadvantage of the ternary system,
Attempted to add high melting point elements. Room temperature before tissue control
The deformation characteristics of Cr at room temperature
Mo without loss of performance improvement effect, Nb, W, Hf,
Ta's5Improvement of strength was observed in some additional elements, and T
Solid solution strengthening in the iAlγ phase was confirmed. Further strength
Increase in strength not only at room temperature but also at high temperature
Was. From the above results, Mo as the fourth additive element, N
b, W, Hf and Ta were selected. In this quaternary
However, the above-mentioned first stage grain boundary β phase is basically satisfied.
The content of the third additive element, Cr, is also in the fourth additive range.
It is a problem if the relation within a certain range is satisfied for the additive element
It was shown that there was no. Further, as a fifth additive element,
Β phase formation ability by Cr of three additive elements, and fourth additive element
Is5 aboveSeed elementOr VDissolves the solid solution effect
A microphone that aims for even higher strength without being heard
When the lower alloying was tried by adding a small amount of B and Si,
Significant improvement in strength from room temperature to 1073K
Was.

The composition range of the component system is such that the effect of adding Cr is T
If the i-excess side and the fourth additive element exceed a certain range, even if the grain boundary β phase is precipitated, the addition of Cr and the fourth additive element due to the lack of superplastic deformability due to the high strength of the matrix. All elements are added in the direction of substitution with Al, and the amount of added Cr is 1% or more (atomic fraction, hereinafter the same) in order to precipitate a β phase. If it is less than 1%, the amount of the grain boundary β phase is not sufficient to cause superplastic deformation,
Is exceeded, a precipitation phase mainly composed of Ti and Cr appears in the matrix, and Cr is no longer distributed to the formation of the grain boundary β phase.

The most important point of the component range of the fourth additive element is that the amount of Cr added as the third additive element should not be exceeded. In particular, Mo (1/30/1990, 53rd Superplasticity Study Group, Osaka International Exchange Center), W (Metal
l. Trans. A (1983) 2170) enables precipitation of the β phase inside the matrix, as is known in the art, and the angle-formed grain boundary β phase is impaired by the strengthening of the matrix. That is, the precipitation site of the β phase does not need to be a matrix, and the formation of the β phase is limited to the grain boundaries. This is because the present inventors have shown that even if the β phase precipitates in the matrix, it contributes to improvement in strength but does not contribute to securing deformability. From the above points, the amount of the fourth additive element is always 0.5 to 3% within a range smaller than the Cr addition amount.
This is because solid solution strengthening by these fourth addition elements is not clear at 0.5% or less. Also, the reason for setting it to 3% or less is that
It is not necessary to strengthen the matrix more than necessary to ensure high-temperature deformability due to the grain boundary β phase, and even if the strengthening is not so great, the strengthening can be sufficiently compensated for by the transformation heat treatment in the second stage. It is.

The fourth or fifth additive element may be
All of Si and B are already added for the purpose of improving the strength at an intermediate temperature or lower, and are aimed at strengthening the solid solution by adding a small amount and precipitating the fine dispersed phase. Therefore, the addition amount is determined so as not to impair the formation of the important grain boundary β phase in the present alloy and the effect of the fourth additive element on the solid solution of the matrix. If it is 0.1% or less, these strengthening effects are not remarkable, and if it exceeds 2%, the matrix strengthening is conversely strengthened by the precipitated phase, and the relaxation of the accumulated strain is not possible even by the deformation of the grain boundary β phase. .

The TiAl-based intermetallic compound molded article of the present invention
The integrated manufacturing method covers the component system determined as described above.
And after smelting the raw material of the alloy, a non-oxidizing atmosphere or 5 × 1
1173K to solidus temperature in a vacuum atmosphere higher than 0 ^-3 Torr
Temperature, the initial strain rate is 5 × 10 ^{-5 to} 5 × 10 ^-1 s
High temperature processing of ec ^-1 and processing rate of 60% or more
Includes grain boundary β phase with a volume fraction of 2 to 25% of precipitated β phase
Become a β + γ two-phase alloy with superplastic deformability, then 10
Cooled down to 873K at a cooling rate faster than K / min
After that, it is processed into a product molded body by superplastic processing,
117 in a neutral atmosphere or in a vacuum higher than 5 × 10 ^-5 Torr
Transformation maintained at 3K to solidus temperature for 2 hours or more
Heat treated, 400MP at room temperature to 1073K
a Processed molded products with α ₂ + γ dual phase structure
High strength TiAl-based intermetallic compound moldings to be manufactured
To manufacture consistently. Consistency of molded products of TiAl-based intermetallic compounds
In the manufacturing method, after melting the raw material of the alloy, the non-oxidizing
In an atmosphere or a vacuum atmosphere higher than 5 × 10 ⁻⁵ Torr
At a temperature of 173K to the solidus temperature, the initial strain rate is 5 × 1
High temperature processing of 0 ^{-5 to} 5 × 10 ^-1 sec ^-1 and processing rate of 60% or more
And the volume fraction of the β phase precipitated at the γ grain boundary is 2 to 25.
Β + γ dual phase with superplastic deformability containing 3% grain boundary β phase
Gold and then at least 87 at a cooling rate higher than 10K / min
After the temperature has dropped to 3K, the product is compacted by superplastic processing.
At 1173K to solidus temperature continuously in the processing equipment
Temperature for 2 to 24 hours.
At room temperature to 1073K
To produce processed products with α ₂ + γ dual phase structure
Integrated production of TiAl-based intermetallic compound molded products with high strength
You may make it. The TiAl-based intermetallic compound composition
In the method of manufacturing a shaped article, a γ phase in which a β phase is precipitated at a grain boundary by high-temperature processing is used as a matrix to form a γ + β fine two-phase structure containing a slight α ₂ phase. However, this α ₂ phase is a part that could not be completely transformed into a β phase formed by working recrystallization, has no meaning in the present invention, and has a very small volume fraction of several percent or less. The determination of the high temperature processing conditions, must be recrystallized gamma phase to destroy gamma + alpha ₂ dual phase structure after initial melting and casting. At the processing temperature and degree of processing required to cause recrystallization of the γ phase, the precipitated β phase already formed by heat treatment before thermal transformation or high-temperature processing can sufficiently withstand deformation, and finally It is considered that the crystal γ phase became a structure in which the β phase segregated at the γ phase grain boundary with the β phase deformed in the grain growth process as a barrier. Such a mechanism has been proposed based on the experimental results so far. Based on this hypothesis, necessary high-temperature processing conditions will be examined. First, temperature, but when Cr is the third additive element, the β phase has already been
Since being formed on the alpha ₂ phase of the initial lamellar structure is revealed by the inventors, it thermal processing recrystallization in the formation of β phase is not necessarily required condition is indicated,
The processing temperature was set to 1173 K or more necessary for recrystallization of the γ phase. If the temperature is lower than this temperature, recrystallization of the γ grains does not sufficiently occur, and it is difficult to crystallize the β phase at the γ grain boundaries. Further, in order to obtain a uniform structure, the working ratio was set to 60% or more. If the working ratio is lower than this, an unrecrystallized region is formed,
Insufficient β + γ dual phase structure containing grain boundary β phase, γ
This is because the β phase remains inside the matrix and it is difficult to impart superplastic deformability. On the other hand, when the initial strain rate is 5 × 10 ⁻¹ sec ⁻¹ or more, a work deformation structure is formed in addition to a recrystallization structure, and a grain boundary β phase cannot be obtained. The initial strain rate is 5 × 10 ⁻⁵ s
If it is slower than ec- ¹ , fine recrystallized γ grains undergo grain growth, and the effect of fine-grain superplasticity is significantly reduced, making it impossible to achieve superplasticity as in the present invention.

On the other hand, in high-temperature processing, the reason why the non-oxidizing atmosphere or the degree of vacuum is higher than 5 × 10 ⁻³ Torr is that, in the case of an oxidizing atmosphere or a vacuum lower than this vacuum, a TiAl-based atmosphere is used. This is because the intermetallic compound alloy is oxidized to deteriorate various properties. The reason why the cooling rate was faster than 10K / min is the first stage floor, after obtaining a γ-phase having a grain boundary β phase thermomechanical treatment at a high temperature, the superplastic forming using the β-phase Must be applied, if 10K / min
When it cooled at a slower cooling rate, a portion of the β-phase in order to impair the superplastic deformability and transformed into alpha ² phase and γ-phase. Moreover, the transformation heat treatment of the material (beta + gamma) subjected to superplastic forming in a second stage, something that increases the strength by the alpha ² phase and gamma-phase, the transformation heat treatment is critical temperature and time, cooling rate There is no major problem in that the β phase disappears. That is, in consideration of the economics of the process, it is not necessary to unnecessarily reduce the cooling rate, and
If the speed is higher than 0 K / min, the purpose of the transformation heat treatment can be achieved. Further, the reason for setting the cooling temperature to 873K is that lowering the cooling rate and lowering the cooling temperature has the same value as stabilizing the lamellar structure on the TTT diagram, and the β phase required for superplastic deformation. Is to be as stable as possible. That is, the cooling temperature was set as high as possible, and the set maximum value was 873K. When the temperature is lower than this temperature, the lamellar structure is further stabilized, and at the same time, it is necessary to reheat the transformation heat treatment to be performed subsequently.

At the same time, the structure having excellent superplastic deformability formed by high-temperature processing in the first stage is simultaneously processed into a molded product,
The second step is a step of eliminating the β phase by transformation heat treatment. The transformation heat treatment conditions at this time are as follows: from the phase stability of β phase, in the temperature range of 1173K or more and the solidus temperature or less,
A heat treatment for 2 to 24 hours is sufficient. The reason for setting the temperature range in this manner is that the β phase formed in the first stage is in a thermally metastable state, and can easily be transformed into a γ + α ₂ two-phase structure under the set conditions. 1173K
If it is lower than this, the time required for the transformation becomes longer, which is uneconomical. Further, the volume fraction of the α ₂ phase formed by the transformation heat treatment depends on the volume fraction of the initial grain boundary β phase. The grain boundary β phase requires 2% to 25% to cause superplastic deformation without impairing the strength of γ. When the β phase disappears by the transformation heat treatment, the α ₂ phase formed necessarily becomes 5% or more and 40% or less depending on the amount of the initial β phase and transformation heat treatment conditions. If 5
% Or less, the amount of the initial β phase is lower than 2%, or the transformation heat treatment condition causes the phase transformation to be equal to or less than the disappearance of the β phase. That is, the β phase is partially left, which is synonymous with no improvement in strength. If the α ₂ phase is 40% or more, the amount of the initial β phase is higher than 25%, or the transformation heat treatment is shifted to a longer time and higher temperature than the above conditions. This is not practically meaningful, since further increase in strength cannot be expected. The reason for this is that the mechanism for increasing the strength is due to the phase transformation of the grain boundary β phase, and no other factor acts on it. That is, as long as the amount of the grain boundary β phase is within 25%, the volume fraction of the α ₂ phase formed by the phase transformation is 40%.
% Cannot be inevitably exceeded.

On the other hand, the reason why the sample was inserted into a Ti alloy capsule in the isothermal forging, hot extrusion, and rolling, and the inside of the capsule was evacuated to a vacuum higher than 5 × 10 ⁻³ Torr is that each of the subsequent high-temperature This is because, in the processing, the sample itself is prevented from being oxidized so that the sample is not brought into contact with the atmosphere so that the sample itself can be oxidized.

Further, the reason why the sample was sheathed with the Ti alloy in the isothermal forging, hot extrusion, and rolling is that in each subsequent high-temperature working, the minimum necessary for controlling the working structure by the sheath of the Ti alloy. Possible to prevent oxidation
This is because simplicity is recognized in industrial use.

The reason for using a Ti alloy for the capsule or the case in these treatments is that the reactivity at the contact interface with the material is low and the strength ratio at the processing temperature is suitable for processing. by. That is, in the strength ratio of both the sample and these capsules or case,
When the sample strength is extremely high, the capsule or case bears processing strain, and cannot be pressurized in a state close to the hydrostatic pressure. In the worst case, the capsule or the case is broken before controlling the sample structure. If the capsule or case strength is higher than the sample strength, the processing strain is spent on the deformation of the capsule or case, and the load on the sample is reduced. Is destroyed. In the above-mentioned high-temperature processing,
Process to a molded product, and continuously in the processing equipment
Alternatively, a transformation heat treatment may be performed. In addition, the sample
5 × 10 ^-3 Torr inside Ti alloy case
Also degassed in high vacuum, then electron beam welding Ti alloy
The case may be sealed.

[0018]

【Example】

(Example 1) 50.6Ti-46.5Al-2.88Cr intermetallic compound at 5% atomic%, 60% workability at 1473K (1200 ° C), and constant temperature forged material with an initial strain rate of 5 × 10 ^-4 s ^-1 High purity Ti (99.9 wt.%), Al (99.99 wt.%)
%) And Cr (99.3 wt.%) As melting raw materials, and a TiAl-based intermetallic compound containing Cr and about 80 mmφ × 300 mm was alloyed by plasma melting. 1373
As a result of performing a homogenizing heat treatment in vacuum at K (1050 ° C.) for 96 hours, an equiaxed grain structure having a crystal grain size of 80 μm was obtained.
Table 1 shows the chemical analysis values after the homogenization heat treatment. A cylindrical ingot of 35 mmφ × 42 mm was cut out from the ingot by electric discharge machining and subjected to constant temperature forging. Forging is performed in a vacuum atmosphere at an initial strain rate of 5 × 10 ^-4 s ^-1 and a sample temperature of 1
The pressure was reduced by 60% at 473K (1200 ° C). FIG. 1 shows a photograph of the structure of this sample after constant temperature forging. Along with the structure consisting of equiaxed fine crystal grains having an average grain size of 18 μm, several μm
A grain boundary precipitated phase having the following thickness was observed. This grain boundary phase was later identified as β phase. From ingot material after forging, gauge part size 11.5 × 3 × 2 by wire cutting
A tensile test piece of mm ³ was cut out, and a tensile test was performed in a vacuum atmosphere while changing a strain rate and a test temperature. The test was performed until the sample was broken at a constant test temperature and strain rate for each sample, and a true stress-true strain diagram was obtained. As an example of the result showing superplasticity, an elongation value of about 480% was obtained at a test temperature of 1473 K (1200 ° C.) and a strain rate of 5 × 10 ⁻⁴ s ⁻¹ . In the sample showing superplasticity, it was observed that the gauge portion was uniformly deformed without showing necking, and it was observed that the grain boundary β phase was stretched after tension. Further, a strain rate sensitivity index (hereinafter referred to as m) calculated from the strain rate dependency of stress.
Value) is 1273K (10
At 00 ° C.), 0.31 was obtained, and at 1473 K (1200 ° C.), 0.49 was obtained. Table 2 shows m-values calculated from these true stress-true strain diagrams to show the temperature dependence. From this table, it is clear that at 1273 K (1000 ° C.) or more, the m value exceeds 0.3 which is an index of the appearance of superplasticity.

Table 3 shows the temperature dependence of the elongation value as a result of these high-temperature tensile tests. From Table 3 1273K (10
It can be seen that the elongation value is significantly improved when the temperature is higher than (00 ° C.). The thus obtained β + γ biphasic structure was further
Heat treatment is performed at 323K for 12 hours. FIG. 2 shows the structure after the heat treatment. Although the form of the grain boundary β phase became unclear, the γ grains did not become coarse, indicating that the initial grain size was around 18 μm. 1473 of the sample subjected to this heat treatment
Table 4 shows the tensile test results at K (1200 ° C.) and a strain rate of 5 × 10 ⁻⁴ s ⁻¹ . From this table, it is clear that the elongation value decreases and the strength increases with the disappearance of the grain boundary β phase. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the β phase disappears and the α ₂ phase is formed by the heat treatment.

[0020]

[Table 1]

[0021]

[Table 2]

[0022]

[Table 3]

[0023]

[Table 4]

(Example 2) A constant temperature forging of 51.6 Ti-43.5 Al-4.87 Cr intermetallic compound in atomic% at 1473 K (1200 ° C.), 60% workability, and an initial strain rate of 5 × 10 ^-4 s ^-1 . The title components were melted by the same plasma melting method as in Example 1,
The sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
C) at a constant pressure of 60%. Table 5 shows the component analysis values after the plasma melting heat treatment. An equiaxed microstructure with an average grain size of about 25 μm is obtained by microstructure control, and several μm
m or less was observed. This grain boundary phase was later identified as β phase as in Example 1. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. As an example of the result showing superplasticity, 1473K (1
200 ℃), 5 × 10 at a strain rate of ^-4 s ^-1 to about 470% or more, 1273K (1000 ℃), 5 × 10 -4 s elongation still of greater than or equal to about 470% at a strain rate of ^-1 is obtained Was. In the sample exhibiting superplasticity, it was observed that the gauge portion was uniformly deformed without showing necking, and it was observed that the grain boundary phase was stretched after tension. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress is a true stress of 0.1.
Is 0.37 at 1273K (1000 ° C)
At 3,1473K (1200 ° C), a figure of 0.46 was obtained. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is also shown in Table 2. From this table, 1273
At K (1000 ° C.) or higher, it is apparent that the m value exceeds 0.3 which is an index of the appearance of superplasticity.

As a result of the high temperature tensile test, the temperature dependence of the elongation value is also shown in Table 3. From Table 3 1273K
It is understood that the elongation value is significantly improved at (1000 ° C.) or higher. The β + γ two-phase structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. Although the form of the grain boundary β phase became unclear, the γ grains were not coarsened, and the initial grain size was around 25 μm. 1473K (1200 ° C.), 5 × 10 ⁻⁴ s ^{−1 of the} heat-treated sample
Table 4 shows the results of the tensile test at a strain rate of. From this table, it was clarified that the elongation value decreased and the strength increased with the disappearance of the grain boundary β phase. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the β phase disappears and the α ₂ phase is formed by the heat treatment.

[0026]

[Table 5]

Example 3 50.1 Ti-46.6 Al-2.80 Cr-atomic%
0.57Si intermetallic compound A constant temperature forging material having a working ratio of 60% at 1473 K (1200 ° C.) and an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ The title component was melted by the same plasma melting method as in Example 1,
The sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
C) at a constant pressure of 60%. Table 6 shows the component analysis values after the plasma melting heat treatment. An equiaxed microstructure having an average particle size of about 30 μm is obtained by controlling the structure, and several μm
m or less was observed. This grain boundary phase was later identified as β phase as in Example 1. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. As an example of the result showing superplasticity, 1473K (1
An elongation value of about 470% or more was obtained at a strain rate of 5 × 10 ⁻⁴ s ⁻¹ at 200 ° C.). In the sample exhibiting superplasticity, it was observed that the gauge portion was uniformly deformed without showing necking, and it was observed that the grain boundary phase was stretched after tension. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress is 1273 when the value of the true stress 0.1 is used.
For K (1000 ° C), 0.30, 1473K (1200
° C) gave a number of 0.41. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is shown in Table 2. From this table, it is clear that at 1273 K (1000 ° C.) or more, the m value exceeds 0.3 which is an index of the appearance of superplasticity.

As a result of these high-temperature tensile tests, the temperature dependence of the elongation value is also shown in Table 3. From Table 3 1273K
It is understood that the elongation value is significantly improved at (1000 ° C.) or higher. The β + γ two-phase structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. Although the form of the grain boundary β phase became unclear, the γ grains were not coarsened and the initial grain size was around 30 μm. 1473 K (1200 ° C.), 5 × 10 ⁻⁴ s ^{−1 of the} heat-treated sample
Table 4 shows the results of the tensile test at a strain rate of. From this table, it was clarified that the elongation value decreased and the strength increased with the disappearance of the grain boundary β phase. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the β phase disappears and the α ₂ phase is formed by the heat treatment.

[0029]

[Table 6]

Example 4 48.9 Ti-47.0 Al-2.83 Cr-atomic%
1.33B intermetallic compound A constant temperature forged material having a working ratio of 60% at 1473 K (1200 ° C.) and an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ The title component was melted by the same plasma melting method as in Example 1,
The sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
C) at a constant pressure of 60%. Table 7 shows the component analysis values after the plasma melting heat treatment. An equiaxed microstructure having an average particle size of about 23 μm is obtained by microstructure control, and several μm
m or less was observed. This grain boundary phase was later identified as β phase as in Example 1. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. As an example of the result showing superplasticity, 1473K (1
An elongation value of about 440% or more was obtained at a strain rate of 5 × 10 ⁻⁴ s ⁻¹ at 200 ° C.). In the sample exhibiting superplasticity, it was observed that the gauge portion was uniformly deformed without showing necking, and it was observed that the grain boundary phase was stretched after tension. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress is 1273 when the value of the true stress 0.1 is used.
For K (1000 ° C.), 0.28, 1473 K (1200
° C) gave a number of 0.47. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is also shown in Table 2. From this table, it is clear that at 1273 K (1000 ° C.) or more, the m value exceeds 0.3 which is an index of the appearance of superplasticity.

As a result of these high-temperature tensile tests, the temperature dependence of the elongation value is also shown in Table 3. From Table 3 1273K
It is understood that the elongation value is significantly improved at (1000 ° C.) or higher. The β + γ two-phase structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. Although the form of the grain boundary β phase became unclear, the γ grains were not coarsened, and the initial grain size was around 23 μm. 1473 K (1200 ° C.), 5 × 10 ⁻⁴ s ^{−1 of the} heat-treated sample
Table 4 shows the results of the tensile test at a strain rate of. From this table, it was clarified that the elongation value decreased and the strength increased with the disappearance of the grain boundary β phase. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the β phase disappears and the α ₂ phase is formed by the heat treatment.

[0032]

[Table 7]

Example 5 49.2 Ti-47.0 Al-2.85 Cr-atomic%
0.99 Nb intermetallic compound A constant temperature forged material having a working ratio of 60% at 1473 K (1200 ° C.) and an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ The title component was melted by the same plasma melting method as in Example 1,
The sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
C) at a constant pressure of 60%. Table 8 shows the component analysis values after the plasma melting heat treatment. An equiaxed microstructure with an average grain size of about 16 μm was obtained by microstructure control, and several μm
m or less was observed. This grain boundary phase was later identified as β phase as in Example 1. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. As an example of the result showing superplasticity, 1473K (1
An elongation value of about 380% or more was obtained at a strain rate of 5 × 10 ⁻⁴ s ⁻¹ at 200 ° C.). In the sample exhibiting superplasticity, it was observed that the gauge portion was uniformly deformed without showing necking, and it was observed that the grain boundary phase was stretched after tension. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress is 1273 when the value of the true stress 0.1 is used.
For K (1000 ° C), 0.30, 1473K (1200
° C) gave a value of 0.42. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is shown in Table 2. From this table, it is clear that at 1273 K (1000 ° C.) or more, the m value exceeds 0.3 which is an index of the appearance of superplasticity.

As a result of these high-temperature tensile tests, the temperature dependence of the elongation value is also shown in Table 3. From Table 3 1273K
It is understood that the elongation value is significantly improved at (1000 ° C.) or higher. The β + γ two-phase structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. Although the form of the grain boundary β phase became unclear, the γ grains were not coarsened, and the initial grain size was around 16 μm. 1473 K (1200 ° C.), 5 × 10 ⁻⁴ s ^{−1 of the} heat-treated sample
Table 4 shows the results of the tensile test at a strain rate of. From this table, it was clarified that the elongation value decreased and the strength increased with the disappearance of the grain boundary β phase. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the β phase disappears and the α ₂ phase is formed by the heat treatment.

[0035]

[Table 8]

Example 6 48.8 Ti-46.8 Al-2.60 Cr-atomic%
1.05 Nb-0.75 Si intermetallic compound Isothermal forging with 60% workability at 1473 K (1200 ° C.) and initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ The title component was melted by the same plasma melting method as in Example 1. Made
The sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
C) at a constant pressure of 60%. Table 9 shows the component analysis values after the plasma melting heat treatment. An equiaxed microstructure with an average particle size of about 20 μm is obtained by microstructure control, and several μm
m or less was observed. This grain boundary phase was later identified as β phase as in Example 1. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. As an example of the result showing superplasticity, 1473K (1
An elongation value of about 350% or more was obtained at a strain rate of 5 × 10 ⁻⁴ s ⁻¹ at 200 ° C.). In the sample exhibiting superplasticity, it was observed that the gauge portion was uniformly deformed without showing necking, and it was observed that the grain boundary phase was stretched after tension. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress is 1273 when the value of the true stress 0.1 is used.
For K (1000 ° C.), 0.29, 1473 K (1200
° C) gave a number of 0.40. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is also shown in Table 2. From this table, it is clear that at 1273 K (1000 ° C.) or more, the m value exceeds 0.3 which is an index of the appearance of superplasticity.

As a result of these high-temperature tensile tests, the temperature dependence of the elongation value is also shown in Table 3. From Table 3 1273K
It is understood that the elongation value is significantly improved at (1000 ° C.) or higher. The β + γ two-phase structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. Although the form of the grain boundary β phase became unclear, the γ grains were not coarsened and the initial grain size was around 20 μm. 1473 K (1200 ° C.), 5 × 10 ⁻⁴ s ^{−1 of the} heat-treated sample
Table 4 shows the results of the tensile test at a strain rate of. From this table, it was clarified that the elongation value decreased and the strength increased with the disappearance of the grain boundary β phase. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the β phase disappears and the α ₂ phase is formed by the heat treatment.

[0038]

[Table 9]

(Comparative Example 1) Isothermal forging of 48.2Ti-48.6Al-3.2V intermetallic compound in atomic% at 1473K (1200 ° C), 60% workability, and initial strain rate of 5 × 10 ^-4 s ^-1 The title components were melted by the same plasma melting method as in Example 1,
The sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
C) at a constant pressure of 60%. FIG. 3 shows a micrograph of the structure of this sample after constant temperature forging. A structure composed of equiaxed fine crystal grains having an average particle size of 25 μm was observed, but no grain boundary phase as observed in the examples was observed. Table 10 shows the component analysis values after the plasma melting heat treatment. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. 1473 of the test conditions under which the superplastic elongation was obtained in the examples.
K (1200 ° C.), about 170 at a strain rate of 5 × 10 ^-4 s ^-1
% Elongation was obtained. The tensile specimen showed necking. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress was 0.20 at 1473 K (1200 ° C.) when the true stress value of 0.1 was used. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is shown in Table 2. From this table, it became clear that this sample did not show superplasticity.

As a result of these high-temperature tensile tests, the temperature dependence of the elongation value is shown in Table 3 together with Example 1. From this table, it is clear that even at a high temperature, the plastic elongation as in the examples was not obtained. The structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. FIG. 4 shows the structure after the heat treatment. It can be seen that the γ grains are coarsened by the heat treatment. Table 4 shows the results of a tensile test of the heat-treated sample at 1473 K (1200 ° C.) and a strain rate of 5 × 10 ⁻⁴ s ⁻¹ . From this table, it is clear that the elongation value decreases and the strength decreases due to the coarsening of the γ grains. Table 13 shows α ₂ before and after the heat treatment determined by the image analysis processing.
2 shows a change in volume fraction of the phase and the β phase. It can be seen that the α ₂ phase exists independently of the heat treatment, and the volume fraction of the β phase is very small.

[0041]

[Table 10]

(Comparative Example 2) Isothermal forging of 50.2Ti-48.6Al-1.2Mn intermetallic compound in atomic% at 1473K (1200 ° C) with 60% workability and initial strain rate of 5 × 10 ^-4 s ^-1 The title components were melted by the same plasma melting method as in Example 1,
The sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
C) at a constant pressure of 60%. An equiaxed fine grain structure of about 32 μm was obtained. Table 11 shows the component analysis values after the plasma melting heat treatment. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. In the example, 1473K (1200
° C), an elongation value of about 120% was obtained at a strain rate of 5 × 10 ^-4 s ^-1 , and the tensile test piece showed necking. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress is 1473K when the true stress value of 0.1 is used.
At (1200 ° C.), a value of 0.20 was obtained. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is shown in Table 2. From this table, it became clear that this sample did not show superplasticity.

As a result of these high-temperature tensile tests, the temperature dependence of the elongation value is shown in Table 3 together with Example 1. This Table 3
From this, it is clear that even at a high temperature, the plastic elongation as observed in the examples was not obtained. The structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. The γ grains were coarsened by the heat treatment. 1473K (1200 ° C.), 5 × 10 ^{−4 of the} heat-treated sample
Table 4 shows the tensile test results at a strain rate of s ^-1 . From this table, it is clear that the elongation value decreases and the strength decreases due to the coarsening of the γ grains. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the α ₂ phase exists independently of the heat treatment, and the volume fraction of the β phase is very small.

[0044]

[Table 11]

(Comparative Example 3) 50.5 Ti-49.5 Al intermetallic compound in atomic%, 74% workability at 1473 K (1200 ° C.), and constant temperature forging material with an initial strain rate of 5 × 10 ⁻⁴ s ^−1. Melted by the same plasma melting method as in Example 1,
A sample subjected to the same heat treatment was subjected to an initial strain rate of 5 × 10 ⁻⁴ s ⁻¹ and a sample temperature of 1473 K (1200) in a vacuum atmosphere.
Constant temperature forging was performed at 60 ° C. under a pressure of 60%. An equiaxed fine grain structure of about 26 μm was obtained. Table 12 shows the component analysis values after the plasma melting heat treatment. A high-temperature tensile test was performed in the same manner as in Example 1 to obtain a true stress-true strain diagram. In the example, 1473K (1200
° C), an elongation value of about 120% was obtained at a strain rate of 5 × 10 ^-4 s ^-1 , and the tensile test piece showed necking. The strain rate sensitivity index (hereinafter referred to as m value) calculated from the strain rate dependency of the stress is 1473K when the true stress value of 0.1 is used.
(1200 ° C.) gave a value of 0.20. The m value was calculated from the true stress-true strain diagram, and the temperature dependence is shown in Table 2. From this table, it became clear that this sample did not show superplasticity.

As a result of these high-temperature tensile tests, the temperature dependence of the elongation value is shown in Table 3 together with Example 1. This Table 3
From this, it is clear that even at a high temperature, the plastic elongation as observed in the examples was not obtained. The structure thus obtained is further subjected to a heat treatment at 1323 K for 12 hours. The γ grains were coarsened by the heat treatment. Table 4 shows the results of a tensile test of the heat-treated sample at 1473 K and a strain rate of 5 × 10 ⁻⁴ s ⁻¹ . From this table, it is clear that the elongation value decreases and the strength decreases due to the coarsening of the γ grains. Table 13 shows changes in the volume fraction of the α ₂ phase and the β phase before and after the heat treatment determined by the image analysis processing. It can be seen that the α ₂ phase exists independently of the heat treatment, and the volume fraction of the β phase is very small.

[0047]

[Table 12]

[0048]

[Table 13]

The above are examples and comparative examples relating to the component systems. Table 14 shows other comparative examples relating to the component systems and examples and comparative examples relating to the atmosphere, temperature, time and cooling rate during the transformation heat treatment. Show. Table 14 also shows the tensile strength at 1073K after the transformation heat treatment of all Examples and Comparative Examples.

[0050]

[Table 14]

[Table 15]

[Table 16]

[Table 17] As is clear from Table 14, the examples of the present invention show excellent material properties with high strength and elongation. On the contrary,
In the comparative example, only one of the strength and the elongation was high, and was not suitable as a structural material.

[0051]

The TiAl-based intermetallic compound alloy of the present invention has a high superplasticity, a high specific strength and a high heat resistance.

[Brief description of the drawings]

FIG. 1 is a photograph of a metallographic structure of a sample obtained in Example 1 of the present invention after isothermal forging.

FIG. 2 is a photograph of the metal structure of the sample shown in FIG. 1 after heat treatment.

FIG. 3 is a photograph of a metal structure of a sample obtained in Comparative Example 1 after isothermal forging.

FIG. 4 is a metallographic photograph of the sample shown in FIG. 3 after heat treatment.

Claims (11)

(57) [Claims] 1. The composition is represented by the following formula in atomic fraction,
the volume fraction of the β phase precipitated at the γ grain boundary is 2 to 25%,
An alloy produced by subjecting a β + γ TiAl-based intermetallic compound having superplastic deformability to a transformation heat treatment,
Α ₂ + having a strength of 400 MPa or more in a temperature range of 3K
A high-strength TiAl-based intermetallic alloy having a gamma two-phase structure. _{Ti a Al 100-abc Cr b} X c X: Nb, Mo, Hf, Ta, one or two or more of W, however 47.5 ≦ a ≦ 52 1 ≦ b ≦ 5 0.5 ≦ c ≦ 3 b ≧ c 2a + b + c ≧ 100 2. The composition is represented by the following formula in atomic fraction,
the volume fraction of the β phase precipitated at the γ grain boundary is 2 to 25%,
An alloy produced by subjecting a β + γ TiAl-based intermetallic compound having superplastic deformability to a transformation heat treatment,
Α ₂ + having a strength of 400 MPa or more in a temperature range of 3K
A high-strength TiAl-based intermetallic alloy having a gamma two-phase structure. _{Ti a Al 100-abd Cr b} Y d Y: Si, B of one, two or where 47.5 ≦ a ≦ 52 1 ≦ b ≦ 5 0.1 ≦ d ≦ 2 2a + b + d ≧ 100 3. The composition is represented by the following formula in atomic fraction,
the volume fraction of the β phase precipitated at the γ grain boundary is 2 to 25%,
An alloy produced by subjecting a β + γ TiAl-based intermetallic compound having superplastic deformability to a transformation heat treatment,
Α ₂ + having a strength of 400 MPa or more in a temperature range of 3K
A high-strength TiAl-based intermetallic alloy having a gamma two-phase structure. _{Ti a Al 100-abcd Cr b} X c Y d X: Nb, Mo, Hf, Ta, W, one of V or two or more Y: Si, one or two provided that 47.5 ≦ a ≦ 52 1 of B ≦ b ≦ 5 0.5 ≦ c ≦ 3 b ≧ c 0.1 ≦ d ≦ 2 2a + b + c + d ≧ 100 4. The TiAl-based intermetallic alloy according to claim 1, wherein the volume fraction of the α ₂ phase is 5 to 40%. 5. The combination according to claim 1, wherein
After smelting the gold material , use a non-oxidizing atmosphere or 5 × 10 ^-3 To
rr, in a high vacuum atmosphere, at a temperature of 1173 K to the solidus temperature, the initial strain rate is 5 × 10 ⁻⁵ to 5 × 10 ⁻¹ sec ⁻¹ ,
Perform high temperature processing at a processing rate of 60% or more and precipitate at the γ grain boundary.
Β + γ two-phase alloy having superplastic deformability including a grain boundary β phase having a volume fraction of 2 to 25% of the β phase, and then 10 K / min.
After lowering the temperature to at least 873K at a faster cooling rate, ultra-塑
The product is processed to a product molded body by a hot working process, and subjected to a transformation heat treatment for 2 hours or more at a temperature of 1173 K to a solidus temperature in a non-oxidizing atmosphere or in a vacuum higher than 5 × 10 ⁻⁵ Torr, An integrated production method of a high-strength TiAl-based intermetallic compound molded product for producing a processed molded product having an α ₂ + γ two-phase structure having a strength of 400 MPa or more in a temperature range of room temperature to 1073 K. 6. The method according to claim 5, wherein the high-temperature processing is isothermal forging, the sample is inserted into a Ti alloy case, and isothermal forging is performed in the atmosphere. 7. The constant temperature forging is performed by degassing the inside of the Ti alloy case into which the sample is inserted under a vacuum higher than 5 × 10 ⁻³ Torr, and then sealing the Ti alloy case by electron beam welding. Integrated production method of a TiAl-based intermetallic compound molded article of the present invention. 8. The method according to claim 1, wherein the high-temperature processing is rolling, and the specimen is made of Ti.
6. The integrated production method of a TiAl-based intermetallic compound molded product according to claim 5, wherein the product is inserted into an alloy case and rolling is performed in the atmosphere. 9. The TiAl according to claim 8, wherein the inside of the Ti alloy case into which the sample is inserted is evacuated with a vacuum higher than 5 × 10 ⁻³ Torr, and the Ti alloy case is hermetically sealed by electron beam welding and rolled. An integrated production method for molded products of base intermetallic compounds. 10. The integrated production method of a TiAl-based intermetallic compound molded product according to claim 5, wherein the high-temperature processing is hot extrusion, the sample is inserted into a Ti alloy case, and the hot extrusion is performed in the atmosphere. 11. The hot extruding is performed by degassing the inside of the Ti alloy case into which the sample is inserted under a vacuum higher than 5 × 10 ⁻³ Torr and then sealing the Ti alloy case by electron beam welding. An integrated production method of the TiAl-based intermetallic compound molded article according to the above.