α+β type titanium alloy and production method therefor

Abstract

The present invention provides an α+β type titanium alloy and a production method therefor, which has an ultrafine structure causing superplasticity under low temperatures and has a high deformation ratio compared to conventional α+β type Ti alloys. The alloy has an ultrafine structure made of equiaxial crystals in which an area ratio of crystals having a grain diameter of 1 μm or less is 60% or more, and maximum frequency grain diameter is 0.5 μm or less, wherein a portion in which the integration degree of plane orientation of the hexagonal close-packed crystal is 1.00 or more exists within a range of 0 to 60 degrees with respect to a normal line of a processed surface of the alloy.

Claims

1. An α+β titanium alloy comprising: an ultrafine structure consisting of equiaxial crystals in which area ratio of crystals having a grain diameter of 1 μm or less is 60% or more, and a maximum frequency grain diameter is 0.5 μm or less; wherein a portion in which an integration degree of plane orientation (0001) of a hexagonal close-packed crystal is 1.00 or more exists within a range of 0 to 60 degrees with respect to a normal line of a processed surface of the alloy.

2. The α+β titanium alloy according to claim 1, wherein the alloy exhibits superplasticity phenomenon when the alloy is deformed at a temperature of 650 to 950° C. with a tensile strain ratio of 1×10.sup.−4 to 10.sup.−2/sec.

3. The α+β titanium alloy according to claim 1, wherein the alloy consists of 4 to 9 mass % of Al, 2 to 10 mass % of V, and a balance of Ti and inevitable impurities.

4. The α+β titanium alloy according to claim 3, wherein the alloy is a Ti-6Al-4V.

5. A production method for an α+β titanium alloy according to claim 1, the method comprising: heating a material at a temperature of 1000° C. or more and maintaining for 1 second or more, cooling the material to room temperature at a cooling rate of 20° C./sec or more, heating the material to a temperature of 700 to 850° C. at a temperature increase rate of 3.5 to 800° C./sec and maintaining for less than 10 minutes, hot working the material at a strain rate of 1 to 50/sec with a strain of 1 or more; and cooling the material at a cooling rate of 5 to 400° C./sec.

6. The α+β titanium alloy according to claim 1, wherein the alloy has 200% or more of fracture elongation.

7. The α+β titanium alloy according to claim 6, wherein the alloy exhibits superplasticity phenomenon when the alloy is deformed at a temperature of 650 to 750° C. with a tensile strain ratio of 1×10.sup.−4 to 10.sup.−2/sec.

Description

BRIEF EXPLANATION OF DRAWINGS

(1) FIG. 1 shows a profile of X-ray diffraction of a material of the present invention.

(2) FIG. 2A shows structures of practical examples of the present invention and graphs showing crystal grain diameter distributions, which were measured by the electron backscatter diffraction (EBSD) method, FIG. 2B shows distribution of integration degree of plane orientation (0001) of the hcp crystal with respect to a normal direction (working direction) of a worked surface of practical examples.

(3) FIG. 3A shows structures of comparative examples and graphs showing crystal grain diameter distributions, which were measured by the electron backscatter diffraction (EBSD) method, FIG. 3B shows distribution of integration degree of plane orientation (0001) of the hcp crystal with respect to a normal direction (working direction) of a worked surface of comparative examples.

(4) FIG. 4 shows appearances of test pieces and fracture elongation.

(5) FIG. 5 is a graph showing the relationship between hot working strain (c) introduced in working of the practical example and fracture elongation in a tensile test in which tensile strain rate was 1×10.sup.−2/sec.

(6) FIG. 6 is a graph showing the relationship between tensile strain rate and fracture elongation in each tensile test temperature.

(7) FIG. 7A shows structures of practical examples of the present invention and graphs showing crystal grain diameter distributions, which were measured by the electron backscatter diffraction (EBSD) method, and FIG. 2B shows distribution of integration degree of plane orientation (0001) of the hcp crystal with respect to a normal direction (working direction) of a worked surface of practical examples.

EXAMPLES

(8) 1. Structure

(9) A plate with 4 mm thick Ti-6Al-4V alloy was prepared and subjected to solid solution treatment at 1100° C. for 30 minutes, and was quenched in water at a cooling rate of 20° C./sec or more, thereby forming an acicular α′ martensite structure. Then, the plate was placed into a furnace and was heated at a temperature increase rate of 3.5 to 800° C./sec. When the temperature of the plate reached 700 to 850° C., the plate was immediately removed from the furnace and was subjected to hot rolling in one pass so that the thickness of the plate was 1.4 mm or less (condition in which applied strain was 1 or more). The peripheral velocity of the roll was set so that the strain rate at exit from the roll was 1 to 50/sec. The plate was cooled at a cooling rate of 5 to 400° C./sec after rolling.

(10) The cross section of the plate was analyzed using an X-ray diffraction (XRD) apparatus. An example of the XRD profile is shown in FIG. 1. FIG. 1 is an XRD profile of Practical Example 1, which was processed at a processing temperature of 800° C., processing strain of 1.05, and a processing strain rate of 7/sec. It can be understood that the phase was almost a single α phase from FIG. 1.

(11) Then, the structural form was observed by an electron backscatter diffraction (EBSD) device (OIM ver. 4.6 produced by TSL Solutions). Specifically, a grain boundary map was made, and crystal grain diameter distribution of the α phase, which was the main structure, was measured. A typical example of the structural form of the plate after processing is shown in FIG. 2A. In FIG. 2A, Practical Example 2 was processed at a processing temperature of 800° C., processing strain of 1.05, and a processing strain rate of 7/sec. In FIG. 2A, the upper row shows grain boundary maps obtained by the EBSD method, showing structures of the rolled surfaces (processed surfaces) of Practical Examples 1 and 2, and the lower row shows graphs showing distributions of the crystal diameter of the α phase corresponding to the structures of Practical Examples 1 and 2. It should be noted that “RD” indicates the rolling direction and “TD” indicates cross direction in the grain boundary maps.

(12) According to the grain boundary maps shown in FIG. 2A, although some amount of forms in which crystal grains elongate toward the rolling direction exist, it was understood that a large amount of the forms are occupied by fine equiaxial crystals. According to the graphs shown in FIG. 2A, it was found that a peak of maximum frequency of grain diameters appeared at 0.5 μm or less, respectively, and the area ratio of crystals in which the grain diameter was 1 μm or less is 60% or more. These results show that an ultrafine structure consisting of equiaxial crystals in which area ratio of crystals having a grain diameter of 1 μm or less was 60% or more, and maximum frequency grain diameter was 0.5 μm or less was formed by the hot rolling.

(13) FIG. 2B shows distribution of the integration degree (crystal orientation) of plane orientation (0001) of the hcp crystal with respect to a normal line direction (processing direction) of the processed surface. As is understood from FIG. 2B, as characteristics of Practical Examples 1 and 2, a portion in which the integration degree of plane orientation (0001) of the hcp crystal is 1.00 or more exists within a range of 0 to 60 degrees with respect to a normal line of a processed surface. Thus, the material of the present invention has crystals having specific orientation within a specific range of angles in a high frequency.

(14) As a comparison, a plate with 4 mm thick Ti-6Al-4V alloy was subjected to solid solution treatment at 1100° C. for 30 minutes, and was quenched by water at a cooling rate of 20° C./sec or more, thereby forming an acicular α′ martensite structure. Then, the plate was placed in a furnace and was heated at a temperature increase rate of 100° C./sec. When the temperature of the plate reached 700 to 800° C., the plate was immediately removed from the furnace. The plate was subjected to hot rolling in one pass so that the thickness of the plate was 2.37 mm and the peripheral velocity of the roll was set so that the strain rate at exit from the roll was 10/sec, and was subjected to hot rolling in one pass so that the thickness of the plate was 1.85 mm and the peripheral velocity of the roll was set so that the strain rate at exit from the roll was 1/sec. After the rolling, the plate was cooled at a cooling rate of 5 to 400° C./sec after rolling, thereby obtaining comparative examples. Comparative Example 1 was processed under conditions of a processing temperature of 700° C., a processing strain of 0.77, and a processing strain rate of 1/sec, and Comparative Example 2 was processed in conditions of a processing temperature of 800° C., a processing strain of 0.77, and a processing strain rate of 1/sec. In FIG. 3A, the upper row shows grain boundary maps obtained by the EBSD method, showing structures of the rolled surfaces (processed surfaces) of Comparative Examples 1 and 2, and the lower row shows graphs showing distributions of the crystal diameter of the α phase corresponding to the structures of Comparative Examples 1 and 2. FIG. 3B shows distribution of the integration degree (crystal orientation) of plane orientation (0001) of the hcp crystal with respect to a normal line direction (processing direction) of the processed surface. As is understood from FIGS. 3A and 3B, although equiaxial crystals in which area ratio of crystals having a grain diameter of 1 μm or less is 60% or more, and maximum frequency grain diameter is 0.5 μm or less, was obtained, portions in which the integration degree of plane orientation (0001) of the hcp crystal is 1.00 or more is distributed in a wide range of angles, crystal orientation was low and was nearly random. This is because the introduced strain was low at 0.77, and as mentioned below, the fracture elongation was less than 200% when a tensile test was performed at a tensile test temperature of 650° C. (Comparative Example 1) and 700° C. (Comparative Example 2) and at a tensile strain rate of 0.01/sec.

(15) 2. Tensile Test

(16) Practical examples were produced under the conditions as above, and were formed in a shape shown in FIG. 4 as tensile test pieces (Practical Examples 3 to 13). The tensile test was performed at a predetermined test temperature while changing tensile strain rate from 1×10.sup.−4 to 1×10.sup.−2, and exhibition of the superplasticity was evaluated. Tensile test temperature was set at 650° C., 700° C., or 750° C., which are lower than the temperature at which the superplasticity is caused. For example, in conventional Ti-6Al-4V alloys (crystal diameter: 3 to 10 μm, equiaxial crystals α+β structure)), the superplasticity appears at 800 to 950° C., but the tensile test was performed at at least 150° C. lower than this temperature. In the tensile test, when strain rate sensitivity index m of deforming stress is 0.3 or more and fracture elongation (plastic elongation) is 200% or more, it was judged that the superplasticity was caused based on the general difinishon. For comparison, a plate with 4 mm thick Ti-6Al-4V alloy was processed under the same conditions shown in Table 1 through the same process as Comparative Examples 1 and 2, thereby obtaining Comparative Examples 3 to 6.

(17) Appearances of fracture elongations in test pieces after the tensile test are shown in FIG. 4. As shown in FIG. 4, the Ti-6Al-4V alloy plates of the present invention (maximum frequency crystal diameter d.sub.a=0.5 μm or more) showed high fracture elongation of 200% or more in all test conditions, and it was confirmed that the superplasticity was caused at a tensile test temperature of 650 to 750° C. and at a tensile strain rate of 1×10.sup.−4 to 1×10.sup.−2.

(18) Processing conditions, structure forms, tensile test conditions, and results thereof are shown in Table 1. Area ratio of crystals with diameters of 1 μm or less and maximum frequency crystal diameter were measured by an EBSD method. In Table 1, the case in which a portion in which the integration degree of plane orientation (0001) of the hcp crystal was 1.00 or more exists within a range of 0 to 60 degrees with respect to a normal line of the processed surface was observed is indicated as “Yes” and the case in which the superplasticity was caused is indicated as “Exists”. As shown in Table 1, in Practical Examples 3 to 13, the area ratio of crystals having a grain diameter of 1 μm or less was 60% or more, and the maximum frequency crystal diameter was 0.5 μm or less, and a portion in which the integration degree of plane orientation (0001) of the hcp crystal was 1.00 or more exists within a range of 0 to 60 degrees with respect to a normal line of the processed surface, and comprised fine crystal structure. As a result, it may be recognized that the superplasticity was caused at a low temperature of 650 to 750° C. and at a high tensile strain rate of 1×10.sup.−4 to 1×10.sup.−2. In contrast, in Comparative Examples 3 and 6, the processing strain was low at less than 1, a portion in which the integration degree of plane orientation was 1.00 or more did not exist within a range of 0 to 60 degrees with respect to a normal line of the processed surface, and maximum frequency crystal diameter was greater than 0.5 μm. In Comparative Examples 4 and 5, the processing strain was low at less than 1, a portion in which the integration degree of plane orientation was 1.00 or more did not exist within a range of 0 to 60 degrees with respect to a normal line of the processed surface, and as a result, the strain rate sensitivity index m of deforming stress was less than 0.3, and the superplasticity was not caused.

(19) TABLE-US-00001 TABLE 1 Area ratio Integration Strain Condition of processing (Particle Maximum degree Condition of rate starting material diameter frequency (1 or more tensile test Fracture sensi- Presence of Processing Processing is 1 μm particle exists within Strain elon- tivity super - temperature Processing strain rate or less) diameter a range of 0 to Temperature rate gation index plasticity (° C.) strain (/sec) (%) (μm) 60 degrees) (° C.) (/sec) (%) m phenomenon Comparative 750 0.52 10 84 0.75 No 650 0.01 130 0.15 None Example 3 Comparative 700 0.77 1 99 0.45 No 650 0.01 170 0.15 None Example 4 Practical 800 1.05 10 68 0.49 Yes 650 0.001 271 0.31 Exists Example 3 Practical 800 1.05 22 66 0.45 Yes 650 0.01 220 0.31 Exists Example 4 Practical 800 1.55 22 78 0.40 Yes 650 0.01 250 0.31 Exists Example 5 Comparative 800 0.77 1 92 0.45 No 700 0.01 180 0.18 None Example 5 Practical 800 1.05 7 71 0.49 Yes 700 0.0001 552 0.37 Exists Example 6 Practical 750 1.05 10 73 0.45 Yes 700 0.001 400 0.37 Exists Example 7 Practical 750 1.05 22 74 0.43 Yes 700 0.01 250 0.37 Exists Example 8 Practical 800 1.55 39 68 0.49 Yes 700 0.01 280 0.37 Exists Example 9 Practical 800 1.24 10 69 0.49 Yes 700 0.01 270 0.37 Exists Example 10 Practical 850 1.05 10 69 0.48 Yes 750 0.001 600 0.50 Exists Example 11 Comparative 800 0.52 10 81 0.70 No 750 0.01 160 0.21 None Example 6 Practical 700 1.05 7 70 0.38 Yes 750 0.01 350 0.50 Exists Example 12 Practical 850 1.55 10 64 0.48 Yes 750 0.01 410 0.50 Exists Example 13

(20) FIG. 5 shows a relationship between processing strain introduced by hot working at a temperature of 750 to 850° C. for obtaining practical examples and fracture elongation in the tensile test at a tensile strain rate of 1×10.sup.−2/sec of the practical examples obtained by the tensile test. As shown in FIG. 5, if the processing strain was less than 1, the fracture elongation was less than 200% due to difference of the structural form and that a portion in which the integration degree of plane orientation was 1.00 or more did not exist within a range of 0 to 60 degrees with respect to a normal line of the processed surface, and thus, the superplasticity was not caused.

(21) 3. Comparison with the Conventional Material

(22) Fracture elongations of an invented material, a severe deformation material, which is refined by severe deformation process according to “METALLURGICAL AND MATERIALS TRANSACTIONS” (Y. G. K O et al., 2006, 37A, p. 381-391), and a conventional material of Ti-6Al-4V alloy, are compared. The conventional material had an average crystal diameter d of 11 μm and was subjected to anneal at 850° C. for 2 hours. The severe deformation material was produced by the ECAP method with a processing strain of 3.92 and had an average crystal diameter d of 0.3 μm. FIG. 6 is a graph showing the relationship between tensile strain rate (1×10.sup.−4 to 1×10.sup.−2) of the invented materials (Practical Examples 3, 4, 6 to 8, 11, and 12) which were obtained by hot working performed at a temperature of 750 to 850° C. at a processing strain of 1.05 and fracture elongation. As shown in FIG. 6, in the invented materials, the fracture elongation was greatly improved compared to the conventional material in tensile strain rate (1×10.sup.−4 to 1×10.sup.−2) at the tensile test temperatures. Furthermore, the invented materials had fracture elongation equal to or greater than that of the severe deformation materials at each tensile test temperature at each tensile strain rate. Specifically, the fracture elongation of the severe deformation material at a tensile test temperature of 650° C. in a strain rate of 1×10.sup.−2 was less than 200%, but the fracture elongation of the invented material was good at more than 200%.

(23) Table 2 shows strain rate sensitivity indexes m of invented materials (Practical Examples 4, 8, and 12), the severe deformation material according to “METALLURGICAL AND MATERIALS TRANSACTIONS”, and the conventional material at each plastic deformation temperature (tensile test temperature) in a strain rate of 1×10.sup.−2. In general, value m in ordinary plastic deformation is about 0.1 to 0.2 or more, but m is large within 1>m≧0.3 in a region of the superplasticity. The materials of the present invention showed higher values m than the severe deforming material and the conventional material, and exceeded 0.3, and showed superior superplasticity characteristics.

(24) TABLE-US-00002 TABLE 2 Plastic deformation Strain rate sensitivity index m temperature (Tensile Severe test temperature) Invented deformation Conventional (° C.) material material material 650 0.31 0.24 0.10 700 0.37 0.28 0.11 750 0.50 — —

(25) FIG. 7A shows a structural form of the invented material after a tensile test performed at a temperature of 700° C. at a strain rate of 1×10.sup.−2. It should be noted that the material of the present invention was produced by the same process as for Practical Examples 1 to 13, and the temperature increase rate in hot rolling was 12° C./sec and the hot rolling was performed in one pass such that the thickness of the material was 1.4 mm when the material temperature was 700° C. The rolling was performed setting the peripheral velocity of the roll such that the strain rate at exit of the roll was 7/sec. The cooling rate after rolling was 100° C./sec. In FIG. 7A, the upper row shows grain boundary maps obtained by the EBSD method, showing structures of the rolled surfaces (processed surfaces) of the invented material, and the lower row shows graphs showing distributions of the crystal diameter of the α phase. FIG. 7B shows distribution of the integration degree (crystal orientation) of plane orientation (0001) of the hcp crystal with respect to a normal line direction (processing direction) of the processed surface. As shown in FIG. 7A, the invented material had a uniform and fine equiaxial structure having crystal grain diameter was about 1 μm after the tensile test. Although the maximum frequency crystal diameter was 1.15 μm and crystal orientation was degraded according to FIG. 7B compared to the invented material before the tensile test, it may be understood that high strength was maintained before deformation since uniform equiaxial crystals with a diameter of about 1 μm were formed.

(26) Thus, according to the present invention, a Ti-6Al-4V alloy plate composed of approximately a single α phase, and comprising a fine equiaxial crystals structure in which area ratio of crystals having a grain diameter of 1 μm or less is 60% or more, and maximum frequency grain diameter is 0.5 μm or less, wherein a portion in which the integration degree of plane orientation (0001) of the hcp crystal is 1.00 or more exists within a range of 0 to 60 degrees with respect to a normal line of a processed surface of the alloy can be obtained by performing plastic working in suitably controlling the processing temperature and processing rate using an α′ martensite structure as a starting structure. In the processing, an ultrafine structure can be obtained only by processing strain of 1 or more (for example, a 4 mm thick plate is worked to 1.4 mm thick or less by rolling). The reason for this may be said to be that non-contiguous dynamic recrystallization, which hardly acts conventionally, actively acts by hot working at a high strain rate using an α′ martensite as a starting structure. Therefore, the processing can be more practically performed compared to a severe deforming process, and production cost can be restricted to the same amount as the cost for production of existing Ti alloy plates. Therefore, Ti-6Al-4V alloy plates having ultrafine crystal grains causing superplasticity at low temperatures and high rates of deformation can be produced by a simple production method using existing machinery.

(27) In the present invention, since crystal grains are refined by hot working under suitable processing conditions using an α′ structure of a Ti alloy as a starting structure, the method can be applied not only to Ti-6Al-4V alloys, but also other α+β type alloys, and superplasticity at low temperatures with high deformation rate in other α+β type alloys. For example, as other α+β type alloys, Ti-8Mn, Ti-3Al-2.5V, Ti-6Al-6V-2Sn, Ti-7Al-1Mo, Ti-6Al-2Sn-4Zr-6Mo, Ti-5Al-2Cr-1Fe, and Ti-6Al-2Sn-4Zr-2Mo may be mentioned.

(28) The present invention can be applied to all products of Ti alloys which are subjected to superplastic forming and to all Ti alloy members which are subjected to superplastic blow molding or diffusion bonding (SPF/DB). For example, the present invention can be applied to Ti alloy members for aircraft (refer to “JOM” L. D. HefTi, 2010, 62-5, pp. 42-45). Furthermore, the present invention can be applied to members that are subjected to superplastic forming, such as chemical plants, energy production plants, general consumer products, and sporting goods. Furthermore, since the α+β type Ti alloys of the present invention cause superplasticity at low temperatures (650° C. or more) with a high strain rate of 10.sup.−2/sec, which is identical to the industrial production rate, and high strength and fine crystal structure can be obtained after superplastic deformation, the invention can be used in primary processing to produce plates, rods, and wires.