JP2686020B2 - Superplastically deformable β + γTiAl-based intermetallic alloy and method for producing the same - Google Patents

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a TiAl-based alloy having superplastic deformability, and a processed product into a complex shape by superplastic working and a method for producing the same, which is applied to a high specific strength heat resistant structural member. Used for.

[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), Ta added (US Patent: 484281)
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). So far, this is the only example of invention by compound addition, but it is also necessary to examine the fourth and fifth additive elements in order to improve a plurality of characteristics. 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,
It is a process of forming using the superplastic deformability of the design material to form a product close to the final product shape.

[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
+ Γ two-phase structure, the relaxation of the introduced strain is converted to the β phase rich in deformability, and the superplastic deformability is loaded without impairing the excellent strength of TiAl. The details will be described below.

Precipitation of the grain boundary β phase is an absolute condition for imparting the superplastic deformability. Therefore, including the reported Cr, six types of β-stabilizing elements Mo, V, Nb, Fe, and Mn were selected as the third additive element with respect to Ti, and as a result of controlling the structure, clear grain boundaries were obtained. I was able to observe the precipitation phase
Since it was only Cr, it was decided to select Cr as the third additive element. Further, addition of refractory elements was attempted at any time in order to compensate for the lack of strength, which was a drawback of the TiAlCr ternary system, without impairing the precipitation morphology of the grain boundary β phase. As a result of investigating the deformation characteristics at room temperature before applying the tissue control,
Without compromising the room temperature compressive deformability improving effect of Cr, M
It was confirmed that the six additive elements of o, V, Nb, W, Hf, and Ta improved in strength, and solid solution strengthening in the TiAlγ phase was confirmed. Further, the strength was increased not only at room temperature but also at high temperature.

From the above results, M was selected as the fourth additive element.
o, V, Nb, W, Hf and Ta were selected. Also in this quaternary system, the above-mentioned grain boundary β phase is basically satisfied, and its component range also satisfies the relationship that the amount of Cr, which is the third additive element, is within a certain range with respect to the fourth additive element. If it was done, it was shown that there was no problem. Further, as a fifth additive element, a microalloy aiming at further strengthening without impairing the β phase forming ability of Cr of the third additive element and the solid solution effect of the six elements of the fourth additive element. Was tried by adding a small amount of B and Si.
Since a remarkable improvement in the strength up to 3K was observed, the final component system was changed from the ternary system posted in claim 3 to the claim 6
It was decided to be up to the five-element system listed in.

In the composition range of the component system, the effect of Cr addition 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 β phase in the grain boundary is not enough to cause superplastic deformation, and it is 5%.
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 addition amount of Cr, which is the third additive element, is not exceeded. Especially Mo (1/30/1990, 53rd Superplasticity Research 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
In order to secure the high temperature deformability by the grain boundary β phase, it is not necessary to strengthen the matrix more than necessary, and even if the strengthening is not so great, the strengthening component can be sufficiently supplemented by the transformation heat treatment in the subsequent process. Is.

The fifth additive elements Si and B have already been added for the purpose of improving the strength at an intermediate temperature or lower, and are aimed at solid solution strengthening by addition of a small amount and precipitation effect by a fine dispersion precipitation 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.
At 0.1% or less, these reinforcing effects are not remarkable and 2
On the contrary, when the content exceeds%, the matrix strengthening becomes stronger due to the precipitation phase, and the deformation of the grain boundary β phase does not allow relaxation of the accumulated strain.

After the components thus determined have been melted, a high temperature processing treatment is carried out to form a γ + β fine two-phase structure having a γ phase with a β phase precipitated at the grain boundaries as a matrix. After melting the sample, homogenize it, but set the temperature range to 1
273 K to solidus temperature or less, heating time 2 to 100
The reason for setting the time is to remove macrosegregation by melting and to stabilize the lamellar phase in which the β phase is partially precipitated in the initial α ₂ phase due to the equilibration of the structure. The fine two-phase structure consisting of the γ phase and the metastable β phase contains some α ₂ phase, but this α ₂ phase could not be transformed into the metastable β phase formed by work recrystallization. In the present invention, it does not make any sense in the present invention, and the volume fraction is a very small amount of several% or less. The reason for this latter will be described in detail below.

To determine the high temperature processing conditions, it is necessary to destroy the γ + α ₂ two-phase structure after the initial melt casting and recrystallize the γ phase. 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. Although such a mechanism has been proposed based on the experimental results so far, the necessary high temperature processing conditions will be examined based on this hypothesis.
First, regarding the temperature, when Cr was used as the third additive element, the β phase had already formed in the initial lamellar structure α ₂ at the stage of the melting heat treatment.
It became clear by the inventors that they were formed in a phase,
The processing temperature was set to 1173 K or higher, which is necessary for the recrystallization of the γ phase, because it was shown that the thermal processing and recrystallization is not always necessary for the formation of the β phase. If the temperature is lower than this temperature, recrystallization of γ grains does not sufficiently occur, and it is difficult to crystallize the β phase at γ grain boundaries. Further, in order to obtain a uniform structure, the working ratio was set to 60% or more. If the workability is lower than this, a non-recrystallized region is formed, the β + γ two-phase structure containing the grain boundary β phase cannot be sufficiently formed, and the β phase remains inside the γ matrix, so that the superplastic deformability is imparted. Because it is difficult.

On the other hand, the initial strain rate is 5 × 10 ^-1 se
This is because when c- ¹ or more, a work-deformation structure is formed in addition to the recrystallization structure, and the grain boundary β phase cannot be obtained. When the initial strain rate is slower than 5 × 10 ⁻⁵ sec ⁻¹ , fine recrystallized γ grains cause grain growth, and the effect of fine grain superplasticity is remarkably reduced, and the superplasticity of the present invention is reduced. This is because expression is impossible and productivity is significantly reduced. When such processing conditions are set, the volume fraction of the grain boundary β phase inevitably becomes 2 to 25%, but if it is lower than 2%, the β phase necessary for superplastic working is absolutely It is not suitable because it will be insufficient. Further, if it is higher than 25%, the material strength as TiAl decreases.

Further, in the high temperature processing, the reason why the non-oxidizing atmosphere or the degree of vacuum is higher than 5 × 10 ⁻³ Torr is that the TiAl group is used when the oxidizing atmosphere or the degree of vacuum is lower than this degree. This is because the intermetallic compound alloy is oxidized and various characteristics are deteriorated. The reason why the cooling rate is higher than 10 K / min is that after the γ phase having the grain boundary β phase is obtained by thermomechanical treatment at high temperature, the β phase must be used for superplastic working. This is because if cooling is performed at a cooling rate lower than 10 K / min, a part of the β phase transforms into the α ₂ phase and the γ phase and impairs superplastic deformability. Further, the reason why the temperature lowering temperature is set to 873K is that lowering the cooling rate and lowering the temperature lowering temperature is the same value as stabilizing the lamellar structure on the TTT diagram, and β necessary for superplastic deformation
This is to keep the phases as stable as possible. That is,
Set the cooling temperature as high as possible and set the maximum value to 8
73K. 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.

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

Furthermore, the reason why the sample is sheathed with a Ti alloy in constant temperature forging, hot extrusion, and rolling is that the minimum required for controlling the texture by the sheath of Ti alloy in each subsequent high-temperature machining. It is possible to prevent the limit of oxidation,
This is because simplicity is recognized in industrial use.

The reason why the Ti alloy is used for the capsule or the case in these treatments is that the reactivity at the contact interface with this 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. When the strength of the capsule or the case is higher than the strength of the sample, the processing strain is consumed for the deformation of the capsule or the case, the load on the sample is reduced, and at the same time, the process recrystallization does not proceed. This is because the case will be destroyed. The present invention exhibits superplastic deformability at 1173 K or higher.

[0020]

【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. 1323
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 the results of these high temperature tensile tests. Table 3 to 127
It can be seen that the elongation value is remarkably improved at 3 K (1000 ° C.) or higher.

[0021]

[Table 1]

[0022]

[Table 2]

[0023]

[Table 3]

(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 4 shows 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
Elongation value of about 470% or more at a strain rate of 5 × 10 ^-4 s ^-1 at 200 ° C, and about 470% or more at a strain rate of 1273K (1000 ° C) of 5 × 10 ^-4 s ^-1. It 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. In addition, the strain rate sensitivity index (m value below) calculated from the strain rate dependence of stress is 10% of true stress.
Value is 0.3 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. Table 3 also shows the temperature dependence of the elongation values as the results of these high temperature tensile tests.

[0025]

[Table 4]

(Example 3) 50.1Ti-46.6Al-2.80Cr-in 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 5 shows the component analysis values after the plasma melting heat treatment. By controlling the structure, an equiaxed fine structure with an average grain size of about 30 μm can be obtained, and several μ at the grain boundaries
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 also shown in Table 2. From this table, it is clear that at 1273 K (1000 ° C.) or higher, the m value exceeds 0.3, which is an index of the superplasticity expression value. Table 3 also shows the temperature dependence of the elongation values as the results of these high temperature tensile tests.

[0027]

[Table 5]

(Example 4) 48.9Ti-47.0Al-2.83Cr- in 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 6 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. Table 3 also shows the temperature dependence of the elongation values as the results of these high temperature tensile tests.

[0029]

[Table 6]

(Example 5) 49.2Ti-47.0Al-2.85Cr- in 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 7 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
A value of 0.42 was obtained in ° C. 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. Table 3 also shows the temperature dependence of the elongation values as the results of these high temperature tensile tests.

[0031]

[Table 7]

(Embodiment 6) 48.8Ti-46.8Al-2.60Cr- in 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 8 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. Table 3 also shows the temperature dependence of the elongation values as the results of these high temperature tensile tests.

[0033]

[Table 8]

Comparative Example 1 48.2 Ti-48.6 Al-3.2 V intermetallic compound in atomic% 60% workability at 1473 K (1200 ° C.), isothermal forging with initial strain rate of 5 × 10 ^-4 s ^-1 The material described above 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%. FIG. 2 shows a microstructure photograph of this sample after isothermal 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 9 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. Table 2 also shows m values calculated from these true stress-true strain diagrams and showing the temperature dependence. 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.

[0035]

[Table 9]

Comparative Example 2 50.2 Ti-48.6 Al-1.2 Mn intermetallic compound in atomic% 60% workability at 1473 K (1200 ° C.), constant temperature forging with an initial strain rate of 5 × 10 ^-4 s ^-1 The material described above 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%. An equiaxed fine grain structure of about 32 μm was obtained. 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. 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. Table 2 also shows m values calculated from these true stress-true strain diagrams and showing the temperature dependence. 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. It is clear from Table 3 that the plastic elongation as seen in the examples is not obtained even at high temperatures.

[0037]

[Table 10]

Comparative Example 3 50.5 Ti-49.5 Al intermetallic compound in atomic% 74% workability at 1473 K (1200 ° C.), constant temperature forging material with initial strain rate of 5 × 10 ^-4 s ^-1 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 26 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. Table 2 also shows m values calculated from these true stress-true strain diagrams and showing the temperature dependence. From this Table 3, it became clear that this sample does not exhibit superplasticity. As a result of these high temperature tensile tests, the temperature dependence of the elongation value is shown in Table 3 in Example 1
It is shown together with. It is clear from Table 3 that the plastic elongation as seen in the examples is not obtained even at high temperature.

[0039]

[Table 11]

The above are examples and comparative examples relating to the component system, but other comparative examples relating to the component system are shown in Table 12 together with the above examples and comparative examples.

[0041]

[Table 12]

[Table 13]

Table 13 shows examples and comparative examples concerning homogenization conditions and processed structure control conditions (temperature, strain rate, processing rate, atmosphere, cooling rate, processing method) in the production.

[0043]

[Table 14]

[Table 15] According to Table 13, in the present invention examples, all m values exceed 0.3 which is an index of superplasticity development. On the other hand, in each of the comparative examples, the m value is less than 0.3.

[0044]

Since the TiAl-based intermetallic compound alloy of the present invention has a high superplastic deformability, the formability is improved and a shape close to the final product can be obtained with high accuracy.

[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 metallographic photograph of the sample obtained in Comparative Example 1 after isothermal forging.

─────────────────────────────────────────────────── ─── Continuation of the front page (51) Int.Cl. ⁶ Identification code Internal reference number FI Technical display location C22F 1/00 683 8719-4K C22F 1/00 683 691 8719-4K 691B 8719-4K 691C 692 8719- 4K 692A 694 8719-4K 694A 8719-4K 694B

Claims (10)

(57) [Claims] 1. A ternary system containing Cr as an additive element in a TiAl-based intermetallic compound , and further containing Nb, Mo, Hf, and T.
Add one or more of a, W, and V at the following atomic fractions:
The composition expressed by the following formula consisting of the added multi-component system , characterized in that it is a β + γ two-phase structure exhibiting superplastic deformability containing 2 to 25% of β phase in the γ phase grain boundary in volume fraction TiAl
Base intermetallic compound alloy. 2. A ternary system containing Cr as an additive element in a TiAl-based intermetallic compound, and one or two of Si and B.
The composition represented by the following formula consisting of a multi-element system in which seeds are added in the following atomic fractions, and β phase in the γ phase grain boundary in a volume fraction of 2 to 25
% TiAl-based intermetallic compound alloy having a β + γ two-phase structure exhibiting superplastic deformability. 3. A ternary system containing Cr as an additive element in a TiAl-based intermetallic compound, and further Nb, Mo, Hf, T.
a, W, V, one or more, and Si, B
The composition expressed by the following formula consisting of a multi-component system in which one or two kinds are added in the following atomic fractions, and the β phase is integrated into the γ phase grain boundary.
A TiAl-based intermetallic compound alloy having a β + γ two-phase structure showing a superplastic deformability of 2 to 25% in terms of content. 4. A formation according to any one of claims 1 to 3
1273 after melting a TiAl-based intermetallic compound containing
K to solidus temperature range for 2 to 100 hours
Qualitative treatment, then non-oxidizing atmosphere or 5x10
^-3 Torr, higher vacuum atmosphere, 1173K ~ solidus temperature
Initial strain rate is 5 × 10 ^-1 sec ^-1 or less at
High-temperature processing with a rate of 60% or more, and cooling faster than 10K / min
Γ characterized by lowering the temperature to a minimum of 873K at the cooling speed
And a method for producing a β two-phase TiAl-based intermetallic compound alloy. 5. The high temperature processing is isothermal forging,
Insert into a Ti alloy case and perform the above isothermal forging in the atmosphere.
And having superplastic deformability above 1173K.
A method for producing a TiAl-based intermetallic compound as described above. 6. A Ti alloy case in which the sample is inserted
Part is degassed at a vacuum higher than 5 × 10 ^-3 Torr, and then Ti alloy case
The method for producing a TiAl-based intermetallic compound according to claim 5, wherein the base is sealed by electron beam welding, and isothermal forging is performed in the atmosphere. 7. The high temperature working is rolling, and the sample is made of Ti.
Inserted in an alloy case and rolled in the atmosphere,
The Ti according to claim 4, which has superplastic deformability at 73 K or more.
A method for producing an Al-based intermetallic compound. 8. The Ti alloy case in which the sample is inserted is degassed under a vacuum higher than 5 × 10 ⁻³ Torr, and then the Ti alloy case is removed.
The method for producing a TiAl-based intermetallic compound according to claim 7, wherein the base is sealed by electron beam welding and rolling is performed in the atmosphere. 9. 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.
Performed, claim having superplastic deformability above 1173K
4. The method for producing a TiAl-based intermetallic compound as described in 4 . 10. A Ti alloy case in which the sample is inserted
After degassing the inside with a vacuum higher than 5 × 10 ^-3 Torr, Ti alloy
The case is sealed by electron beam welding and the hot pressing
The method for producing a TiAl-based intermetallic compound according to claim 9, wherein the deposition is performed in the atmosphere.