KR102332018B1 - High temperature titanium alloy and method for manufacturing the same - Google Patents

Abstract

In weight %, Al: 3.0 to 7.0%, Sn: 2.0 to 5.0%, Hf: 1.0 to 6.0%, Nb: 0.1 to 1.5%, Mo: 0.2 to 0.6%, Si: 0.2 to 0.6%, C: 0.05 to A near alpha titanium alloy comprising 0.2%, other unavoidable impurities, and the balance Ti. According to the present invention, it is possible to provide a new near alpha titanium alloy containing Hf that replaces Zr segregated in precipitates and having superior mechanical strength than Ti-1100 alloy or IMI 834 alloy, which is an existing near alpha titanium alloy, and a method for manufacturing the same There are advantages.

Description

HIGH TEMPERATURE TITANIUM ALLOY AND METHOD FOR MANUFACTURING THE SAME

The present invention relates to a high-temperature titanium alloy having excellent mechanical properties and a method for manufacturing the titanium alloy.

Titanium alloy is used in many industrial fields because it has a low density, high strength, excellent specific strength (specific strength) and biocompatibility (biocompatibility).

Titanium alloys currently widely used include CP (commercially pure) titanium in a pure (or commercially pure) state, a near alpha titanium alloy containing a low temperature alpha phase as a main phase, a high temperature beta phase and a low temperature It is classified into an alpha + beta titanium alloy in which an alpha phase coexists, and a beta titanium alloy containing a lot of beta stabilizing elements.

Conventionally, heat-resistant steel or nickel or cobalt-based superalloys have been mainly used as structural materials used in high-temperature environments. However, although the heat-resistant steel or superalloy has excellent mechanical properties at high temperatures, there is a fundamental problem in that it weighs too much due to its high density.

For this reason, a near-alpha titanium alloy with low density compared to heat-resistant steel or superalloy, which has low weight, high specific strength, and excellent mechanical properties at high temperatures, is attracting attention.

Representative near-alpha titanium alloys currently used include Ti-1100 alloy (Ti-6Al-2.7Sn-4Zr-0.4Mo-0.45Si based on wt%) and IMI 834 alloy (Ti-5.8Al-4Sn-3.5Zr-0.5 Mo-0.35Si) alloy and the like.

However, all of the near alpha titanium alloys have disadvantages in that the usable temperature is limited to 600° C. or less, and the strength at room temperature and high temperature is weak. In particular, all of the near alpha titanium alloys contain about 4% of Zr effective for solid solution strengthening in the titanium alloy. However, the Zr is highly segregated in the precipitates formed during the heat treatment of the near alpha titanium alloy, and as a result, there is a problem in that it is not effective in improving the mechanical strength of the near alpha titanium alloy.

An object of one aspect of the present invention is to provide a near alpha titanium alloy that has excellent mechanical properties at room temperature and high temperature and can be used even at a high temperature of 600° C. or more, and a method for manufacturing the same.

More specifically, in the present invention, it is another object to provide a near alpha titanium alloy excellent in mechanical strength and a method for manufacturing the same, which can have a solid solution strengthening effect equal to or greater than that of Zr and does not excessively escape as precipitates formed during aging treatment. do it with

Another object of the present invention is to provide a near alpha titanium alloy having excellent mechanical properties and a method for manufacturing the same by controlling precipitates in the heat treatment process.

In order to achieve the above object, in one aspect of the present invention

In weight %, Al: 3.0 to 7.0%, Sn: 2.0 to 5.0%, Hf: 1.0 to 6.0%, Nb: 0.1 to 1.5%, Mo: 0.2 to 0.6%, Si: 0.2 to 0.6%, C: 0.05 to A near alpha titanium alloy is provided comprising 0.2%, other unavoidable impurities and the balance Ti.

In addition, in another aspect of the present invention

In weight %, Al: 3.0 to 7.0%, Sn: 2.0 to 5.0%, Hf: 1.0 to 6.0%, Nb: 0.1 to 1.5%, Mo: 0.2 to 0.6%, Si: 0.2 to 0.6%, C: 0.05 to casting a titanium alloy containing 0.2%, other unavoidable impurities, and the remainder Ti;

solution heat treating the cast alloy;

There is provided a method for producing a near alpha titanium alloy comprising; aging the solution heat treated alloy.

According to the present invention, it is possible to provide a new near alpha titanium alloy containing Hf that replaces Zr segregated in precipitates and having superior mechanical strength than Ti-1100 alloy or IMI 834 alloy, which is an existing near alpha titanium alloy, and a method for manufacturing the same There are advantages.

In addition, according to the present invention, by controlling the precipitates during heat treatment, it is possible to provide a new near alpha titanium alloy and a method for manufacturing the same, which have superior mechanical strength as well as elongation than the existing near alpha titanium alloy Ti-1100 alloy or IMI 834 alloy. There is this.

Meanwhile, according to the present invention, there is an advantage in that it is possible to provide a new near alpha titanium alloy having excellent mechanical properties and a method for manufacturing the same by suppressing grain growth during high-temperature heat treatment through the addition of alloying elements.

In addition to the above-described effects, the specific effects of the present invention will be described together while describing specific details for carrying out the invention below.

1 is a photograph of the microstructure of Example 5 alloy of the present invention subjected to solution treatment and aging at 700° C./5 hours under an optical microscope.
2 is a microstructure (a) and SADP (selected area diffraction pattern) (b) obtained by solution heat treatment and 650° C./5 hour aging treatment of the alloy of Example 5 of the present invention, and solution treatment of the alloy of Example 5 and The microstructure (c) and SADP (selected area diffraction pattern) (d) after aging at 700° C./5 hours are shown.
3 is a microstructure (a) of Example 5 alloy of the present invention subjected to solution heat treatment and aging at 650° C./5 hours and microstructure (b) aged at 700° C./5 hours observed with a transmission electron microscope. .
4 shows the tensile strength results measured at room temperature and 650° C. after solution heat treatment and aging treatment of the titanium alloys of Examples 4 and 5 of the present invention.
5 shows the tensile strength results measured at room temperature and 650° C. after solution heat treatment and aging treatment of the titanium alloys of Examples 4 and 5 of the present invention.
6 is a view showing a cross-sectional photograph of a specimen in which the titanium alloy of Example 5 of the present invention was subjected to a tensile test at room temperature and 650° C. after aging treatment at 650° C. and 700° C., respectively.
7 is a graph showing the results of the tensile test at 650 ℃ after aging treatment at 700 ℃ the titanium alloy of Example 7 of the present invention.
8 is a graph showing the results of a tensile test at room temperature after aging the titanium alloy of Example 7 of the present invention at 700 ℃.
9 is a graph showing the results of the tensile test at 650 ℃ after aging treatment at 700 ℃ the titanium alloy of Example 8 of the present invention.
10 is a graph showing the results of a tensile test at room temperature after aging the titanium alloy of Example 8 of the present invention at 700 ℃.

Hereinafter, with reference to the drawings, embodiments of the present invention will be described in detail so that those of ordinary skill in the art to which the present invention pertains can easily implement them. The present invention may be embodied in many different forms and is not limited to the embodiments described herein.

In order to clearly explain the present invention, parts irrelevant to the description are omitted, and the same reference numerals are given to the same or similar elements throughout the specification. Further, some embodiments of the present invention will be described in detail with reference to exemplary drawings. In adding reference numerals to components of each drawing, the same components may have the same reference numerals as much as possible even though they are indicated in different drawings. In addition, in describing the present invention, if it is determined that a detailed description of a related known configuration or function may obscure the gist of the present invention, the detailed description may be omitted.

In describing the components of the present invention, terms such as first, second, A, B, (a), (b), etc. may be used. These terms are only for distinguishing the elements from other elements, and the nature, order, order, or number of the elements are not limited by the terms. When it is described that a component is “connected”, “coupled” or “connected” to another component, the component may be directly connected or connected to the other component, but other components may be interposed between each component. It will be understood that each component may be “interposed” or “connected”, “coupled” or “connected” through another component.

Titanium is a polymorphous element having two crystal structures: a high-temperature beta phase with a body centered cubic (BCC) structure and a low-temperature alpha phase with a hexagonal closed packed (HCP) structure. .

When a transition metal is added to the titanium, a region in which the high-temperature beta phase of the titanium is stabilized is widened. In other words, the beta transus temperature, which is the temperature at which the high-temperature beta phase is transformed into the low-temperature alpha phase, decreases.

At this time, the degree of stabilization of the beta phase varies depending on each added alloying element. The degree of stabilization of the beta phase for each alloying element based on Mo (molybdenum) is referred to as the following Mo equivalency.

[Mo]eq = [Mo] + 0.67 [V] + 0.44 [W] + 0.28 [Nb] + 0.22 [Ta] + 2.9 [Fe] + 1.6 [Cr] + 1.25 [Ni] + 1.7 [Mn] + 1.7 [Co] - 1.0 [Al]

On the other hand, Al is a representative alloying element that stabilizes the low-temperature alpha phase in a titanium alloy.

Therefore, when calculating the Mo equivalent, Al is calculated to have a negative value unlike other beta stabilizing element values.

On the other hand, the degree of stabilization of the alpha phase varies depending on the alloying element added to the titanium alloy. The degree of stabilization of the alpha phase for each alloying element based on Al (aluminum) is a concept distinct from the Mo equivalent. Al equivalency).

[Al] _eq = [Al] + 1/6 [Zr] + 1/3 [Sn] + 10 [O]

The near alpha titanium alloy in the present invention is characterized in that most of the Zr in the existing Ti-1100 alloy or IMI 834 alloy is substituted with other elements or does not contain Zr. For this purpose, Hf (hafnium) is included as an essential component in the present invention.

The reason for including Hf in the present invention is as follows.

In general, titanium alloys are strengthened by interstitial element strengthening, solid solution strengthening, precipitation strengthening, and dislocation density or grain refining strengthening mechanisms.

Among the strengthening mechanisms, strengthening by interstitial elements is not preferable because it may cause brittleness.

Therefore, the strengthening mechanism usable by the titanium alloy of the present invention results in the remaining solid solution strengthening, precipitation strengthening, and dislocation density or grain refining strengthening.

Among them, since solid solution strengthening is maintained at a certain level or more even at high temperatures, as a result, high temperature strength and creep resistance can also be improved.

At this time, Hf is the most representative solid solution strengthening element in a titanium alloy along with Al and the like. In particular, Hf may not only strengthen the alpha phase, which is the matrix of the titanium alloy of the present invention, but also be relatively uniformly distributed in the matrix and precipitates formed during the subsequent heat treatment process. As a result, Hf has the advantage of continuously maintaining the solid solution strengthening effect even after heat treatment.

Also, in the present invention, B (boron) may be included as another compositional feature.

When B is added to the titanium alloy, B reacts with Ti to form Ti boride of _{TiB or TiB 2 composition.}

Since Ti boride is crystallized in a liquid state (primary precipitation), it helps to improve mechanical properties by suppressing coarsening of grains during subsequent heat treatment (especially solution heat treatment) of the titanium alloy of the present invention.

Based on the above technical reasons, in one aspect of the present invention

In weight %, Al: 3.0 to 7.0%, Sn: 2.0 to 5.0%, Hf: 1.0 to 6.0%, Nb: 0.1 to 1.5%, Mo: 0.2 to 0.6%, Si: 0.2 to 0.6%, C: 0.05 to 0.2%, other unavoidable impurities and the balance Ti, providing a near alpha titanium alloy.

In one embodiment,

The near alpha titanium alloy is in weight %, Al: 3.0 to 7.0%, Sn: 2.0 to 5.0%, Zr: 1.5 to 3.0%, Hf: 1.5 to 4.5%, Nb: 0.1 to 1.5%, Mo: 0.2 to 0.6 %, Si: 0.2 to 0.6%, C: 0.05 to 0.2%, and other unavoidable impurities and the remainder may include Ti.

In another embodiment,

The near alpha titanium alloy is weight %, Al: 5.5-7.0%, Sn: 2.5-4.0%, Zr: 1.5-3.0%, Hf: 1.5-4.5%, Nb: 0.1-1.5%, Mo: 0.2-0.6 %, Si: 0.2 to 0.6%, C: 0.05 to 0.2%, and other unavoidable impurities and the remainder may include Ti.

In another embodiment,

The near alpha titanium alloy is

In weight %, Al: 3.0 to 7.0%, Sn: 2.0 to 5.0%, Hf: 1.0 to 6.0%, Nb: 0.1 to 1.5%, Mo: 0.2 to 0.6%, Si: 0.2 to 0.6%, C: 0.05 to 0.2%, other unavoidable impurities and the remainder may contain Ti, and may not contain Zr.

In another embodiment

The near alpha titanium alloy is

In weight %, Al: 5.5-7.0%, Sn: 2.5-4.0%, Hf: 2.5-5.5%, Nb: 0.1-1.5%, Mo: 0.2-0.6%, Si: 0.2-0.6%, B: 0.05- 0.2% C: 0.05 to 0.2%, including other unavoidable impurities and the remainder Ti, and may not contain Zr.

In addition, preferably, the near alpha titanium alloy may further include B: 0.05 to 0.2% in the composition.

If the addition amount of Al is less than 3%, the solid solution strengthening effect according to the Al addition is small, as a result, there is a problem that the strength and hardness of the final titanium alloy is lowered.

On the other hand, when the amount of Al added is more than 7%, there _{is a problem in that α 2} phase (Ti ₃ Al) precipitates are excessively formed in the aging process according to the excessive Al addition.

On the other hand, Sn is an almost neutral element with less alpha phase stabilization ability than Al.

It improves high temperature strength and creep resistance along with solid solution strengthening ability.

When the amount of Sn added is less than 2%, the strengthening effect according to the Sn addition is small, and as a result, there is a problem in that the strength and high temperature characteristics of the final titanium alloy are lowered.

On the other hand, when the amount of Sn added is more than 5%, there _{is a problem in that α 2} phase (Ti ₃ Al) precipitates are excessively formed in the aging process according to the excessive Sn addition. In addition, excessive addition of Sn increases the specific gravity of the titanium alloy, there is a problem of lowering the specific strength.

Zr is an almost neutral element that has less alpha phase stabilization ability than Al and Sn, and forms a constant solid solution with Ti and has a solid solution strengthening effect.

When the amount of Zr added is more than 3%, there is a problem in that excessive Zr is segregated in precipitates formed during aging during subsequent aging treatment, making it difficult to maintain the solid solution strengthening effect of the matrix.

Preferably, it may contain 2.5% or less of Zr, more preferably 2% or less, and may not contain Zr at all.

As described above, since most of Hf remains in the base even after aging along with the solid solution strengthening effect, it has an excellent advantage in the solid solution strengthening and maintaining effect.

When the amount of Hf added is less than 1%, the strengthening effect according to the addition of Hf is small, and as a result, there is a problem in that the strength and high temperature characteristics of the final titanium alloy are lowered.

On the other hand, if the amount of Hf added is more than 6%, there is a problem in that the ductility is reduced according to the excessive addition of Hf and the specific strength is lowered by increasing the density of the titanium alloy.

Nb is a beta-phase stabilizing element and improves creep strength and fatigue strength.

When the amount of Nb added is less than 0.1%, there is a problem in that the effect of improving the high-temperature characteristics according to the addition of Nb is not properly exhibited.

On the other hand, when the amount of Nb added is more than 1.5%, there is a problem in that the high-temperature beta phase ratio is increased according to the excessive addition of Nb, so that the high-temperature strength and creep resistance properties are rather deteriorated.

Mo is also a beta-phase stabilizing element like Nb, and improves strength at room temperature and high temperature.

When the amount of Mo added is less than 0.2%, the strengthening effect according to the addition of Mo is small, and as a result, there is a problem in that the strength and high temperature characteristics of the final titanium alloy are lowered.

On the other hand, when the addition amount of Mo is more than 0.6%, there is a problem in that high temperature strength and creep resistance properties are rather deteriorated because the high temperature beta phase ratio is increased according to the excessive addition of Mo.

Si improves high-temperature creep properties by forming Ti silicide such as _{Ti 5} Si ₃ during aging treatment.

When the amount of Si added is less than 0.2%, the amount of Ti silicide to be precipitated is too small, so that it is difficult to improve the high temperature creep properties.

On the other hand, when the amount of Si added is more than 0.6%, the amount of Ti silicide precipitated by excessive Si addition is too large, and as a result, there is a problem that ductility is greatly reduced.

C is a representative alpha stabilizing element along with O, contributing to strength improvement from room temperature to high temperature and improving high temperature creep strength.

When the amount of C added is 0.05% or more, the effect of the addition of C is expressed, and when the amount of C added is more than 0.2%, there is a problem that brittleness may occur in the titanium alloy.

Lastly, in the case of B, if the addition amount is less than 0.05%, the amount of crystallized Ti boride is too small, so there is a problem in that it is difficult to implement the effect of refining the grains.

On the other hand, if the amount of B added is more than 0.2%, the composition of Ti boride that is crystallized according to the excessive B addition _{changes from TiB to TiB 2,} and the shape of Ti boride also changes blocky. As a result, the ductility of the titanium alloy is greatly reduced. there is a problem.

The alloy may include silicide precipitates present at the alpha phase interface.

The alloy, when further comprising B, may include a silicide precipitate present at the alpha phase interface, and a Ti-boride crystal.

The alloy may not include a _{regular α 2 phase.}

The near alpha titanium alloy provided in one aspect of the present invention having the above composition has an Mo equivalent of 0.4 to 0.7, and an Al equivalent of 7.2 to 8.0.

If the Mo equivalent is less than 0.4 or the Al equivalent is higher than 8.0, the fraction of the alpha phase of the cooled titanium alloy during aging treatment or after solution heat treatment is excessively high, so that there is a problem in that the heat treatment temperature is too high. On the other hand, when the Mo equivalent is greater than 0.7 or the Al equivalent is lower than 7.2, the fraction of the beta phase is excessively high and has a microstructure of an alpha + beta alloy, not a near alpha alloy, and thus there is a problem in that mechanical properties at high temperatures are deteriorated.

In another aspect of the invention

In weight %, Al: 3.0 to 7.0%, Sn: 2.0 to 5.0%, Hf: 1.0 to 6.0%, Nb: 0.1 to 1.5%, Mo: 0.2 to 0.6%, Si: 0.2 to 0.6%, C: 0.05 to casting a titanium alloy containing 0.2%, other unavoidable impurities, and the remainder Ti;

solution heat treating the cast alloy;

There is provided a method for producing a near alpha titanium alloy comprising; aging the solution heat treated alloy.

Each component and content range of the titanium alloy will not be repeatedly described as previously described.

After the alloys are forged in the beta single phase region, the solution heat treatment step of quenching after being maintained for 1 hour in the alpha + beta region may be performed.

Thereafter, the aging treatment of the solution heat-treated alloy may be performed.

It can be aged at a temperature of 600°C or higher.

In this case, the forging process may be omitted. In particular, in the titanium alloy of the present invention to which B is added, the crystal grains of the titanium alloy may be refined and uniform due to the crystallization of Ti-boride. In this case, the forging process may be omitted.

The aging treatment temperature is preferably 600 to 800 ℃. Preferably it may be 700 to 800 ℃.

If the aging treatment temperature is lower than 600°C, the aging time is excessively increased due to the low aging treatment temperature, and furthermore, the difference between the aging temperature and the temperature at which the near alpha titanium is used is small. As a result, aging occurs during the use of the near alpha titanium alloy, which causes changes in mechanical properties along with changes in the microstructure of the titanium alloy.

On the other hand, the aging treatment temperature is theoretically possible up to the beta single-phase region or less, but when it is higher than 800° C., coarsening of the precipitates precipitated with grain growth occurs, which may adversely affect the mechanical properties.

Hereinafter, the present invention will be described in more detail through Examples and Experimental Examples. The scope of the present invention is not limited to specific embodiments, and should be construed by the appended claims. In addition, those skilled in the art will understand that many modifications and variations are possible without departing from the scope of the present invention.

<Example>

First, the titanium alloy of the present invention was prepared using a bar-shaped ingot (vacuum arc remelting) using VAR (vacuum arc remelting). Compositions of the titanium alloy, Ti-1100 alloy, and IMI 834 alloy prepared in the present invention are shown in Table 1 below.

alloy name Al
(wt%) Sn
(wt%) Zr
(wt%) Hf
(wt%) Nb
(wt%) Mo
(wt%) Si
(wt%) B
(wt%) C
(wt%) O
(wt%) Ti Ti-1100 6 2.7 4 0.4 0.45 Bal. IMI834 5.8 4 3.5 0.7 0.5 0.35 0.06 Bal. Example 1 3.66 3 0 4 0.2 0.4 0.4 0.1 0.3 Bal. Example 2 3.13 4 2.5 1.5 One 0.4 0.4 0.1 0.3 Bal. Example 3 6.66 3 0 4 0.2 0.4 0.4 0.1 Bal. Example 4 6.13 4 2.5 1.5 One 0.4 0.4 0.1 Bal. Example 5 6.13 4 2.5 1.5 One 0.4 0.4 0.1 0.1 Bal. Example 6 6.66 3 0 4 0.2 0.4 0.4 0.1 0.1 Bal. Example 7 6.66 3 2 2 0.2 0.4 0.4 0.1 0.1 Bal. Example 8 6.66 3 4 0 0.2 0.4 0.4 0.1 0.1 Bal.

The alloys of Table 1 were then density measured using an MD-300S electronic hydrometer, and the measured density was found to be approximately 46 g/cm 3 .

The beta transus temperature of the alloys was measured to be approximately 1050° C. through metallographic experiments.

After that, the alloys were forged in the beta single-phase region and then solution heat treated for rapid cooling after being maintained for 1 hour in the alpha + beta region, aged at a temperature of 600° C. or higher while changing time, and then air-cooled.

In this case, the forging process may be omitted. In particular, in the titanium alloy of the present invention to which B is added, the crystal grains of the titanium alloy may be refined and uniform due to the crystallization of Ti-boride. In this case, the forging process may be omitted.

Aging was carried out at a temperature of 600 to 800 °C.

<Experimental Example 1> Morphology analysis of titanium alloy

The microstructure of the alloys of the examples was observed through an optical microscope, a scanning electron microscope (JEOL, JSM-6610LV), and a transmission electron microscope (JEOL, JEM 2100).

1 is a photograph of the microstructure obtained by solution heat treatment and aging at 700° C./5 hours of the alloy of Example 6 of the present invention observed under an optical microscope.

Since the alloy of Example 6 of the present invention contains 0.1% B, it can be seen from FIG. 1 that it has relatively fine grains even after solution treatment.

The long precipitate shape indicated by the arrow in FIG. 1 is a TiB crystal, and other blocky phases indicated by the arrows correspond to the alpha phase.

2 is a microstructure (a) and selected area diffraction pattern (SADP) (b) of the alloy of Example 6 of the present invention subjected to solution heat treatment and aging at 650° C./5 hours, and solution heat treatment and 700° C./5 The microstructure (c) and selected area diffraction pattern (SADP) (d) of the alloy of Example 6 subjected to time aging are shown.

As shown in FIGS. 2 (a) and (b), it can be seen that the _{α 2} phase is uniformly distributed in the alloy obtained by aging the titanium alloy of Example 6 at 650° C. of the present invention. In addition, a superlattice is observed in the SADP of FIG. 2(b), which directly proves the existence of the _{regular α 2 phase.}

_{On the other hand, as shown in Fig. 2 (c), the α 2} phase was not observed in the alloy obtained by aging the titanium alloy of Example 6 at 700° C. of the present invention. Also, no superlattice was observed in the SADP of FIG. 2(d), which shows that the α ₂ phase no longer exists in the alloy obtained by aging the titanium alloy of Example 6 of the present invention at 700°C.

The α ₂ phase is a regular phase, and in general, the regular phases become thermodynamically unstable as the temperature increases, and eventually disappear. Therefore, the titanium alloy of the present invention, when the aging temperature is run through the 700 ℃, the aging temperature is higher than the rule change temperature on the α ₂ is determined to not exist α ₂ phase.

In addition, the α ₂ phase is generally known to have a very bad effect on the ductility of the titanium alloy. Therefore, if high ductility is required, it is more preferable to age at a high aging temperature where the _{α 2 phase does not exist.}

Meanwhile, it was confirmed in this experimental example that other precipitates were present in addition to _{the α 2} phase during the aging treatment of the titanium alloy of the present invention.

3 is a microstructure (a) of the Example 6 alloy of the present invention subjected to solution heat treatment and aging at 650 ° C. / 5 hours, and microstructure (b) aged at 700 ° C. / 5 hours observed with a transmission electron microscope. will be.

As the EDS component analysis result of FIG. 3 shows, when the alloy of Example 6 of the present invention is aged, it can be seen that very fine silicide precipitates exist at the interface of the alpha-phase plate regardless of the aging temperature.

In addition, the EDS results show that the silicide precipitates contain Hf to a similar extent to the matrix, regardless of the aging temperature. Therefore, Hf contained in the titanium alloy of the present invention is relatively uniformly present without segregation in precipitates such as silicide, and the uniform existence of the Hf indicates that the titanium alloy of the present invention is stronger than the existing titanium alloys containing only Zr. It can be said that it is more advantageous in terms of aspects.

<Experimental Example 2> Analysis of mechanical properties of titanium alloy

For the alloys of Examples, mechanical properties such as tensile strength and elongation were analyzed.

FIG. 4 shows tensile strength results measured at room temperature and 650° C. after solution heat treatment and aging treatment of the titanium alloys of Examples 4 and 6 of the present invention.

5 shows the tensile strength results measured at room temperature and 650° C. after solution heat treatment and aging treatment of the titanium alloys of Examples 4 and 6 of the present invention.

Aging was performed at 700°C.

As shown in FIG. 4, it can be seen that the titanium alloys of Examples 4 and 6 of the present invention have superior strength at both room temperature and high temperature compared to the Ti-1100 alloy and IMI 834 alloy, which are comparative examples.

In particular, the titanium alloy of Example 6 of the present invention was measured to have a very good (about 17%) elongation at room temperature elongation. The excellent elongation of the titanium alloy of the present invention as described above is considered to be _{because the α 2} phase did not precipitate due to the aging treatment at 700° C. or higher together with the factor that the crystal grains of the titanium alloy were refined by the addition of B.

6 is a view showing a cross-sectional photograph of a specimen in which the titanium alloy of Example 6 of the present invention was subjected to a tensile test at room temperature and 650° C. after aging treatment at 650° C. and 700° C., respectively.

First, it can be seen that the fracture surface (FIG. 6 (a)) of the alloy of Example 6 subjected to room temperature tensile test after aging at 650° C. has almost a quasi-cleavage.

On the other hand, the fracture surface of the alloy of Example 5 (FIG. 6 (b)), which was subjected to a tensile test at room temperature after aging at 700° C., coexists with a fracture surface of a shallow and extended dimple shape and a fracture surface of a cleavage plane crossing the grain boundary. FIG. 6( b ) agrees well with the alloy having a room temperature elongation of about 17% at the room temperature elongation of FIG. 5 .

Changes in the cutting plane, and an elongation as described above, when the titanium alloy of the present invention the aging treatment at 650 ℃ case where, while the α ₂ phase to precipitate, which adversely affect the ductility of the titanium alloy aged at 700 ℃ process α ₂ on the It is considered that this is because precipitation is suppressed.

Therefore, the tensile test results and fracture surface observation results of FIGS. 4 to 6 show that aging treatment at 650° C. or higher is effective to produce a high-strength near alpha titanium alloy, and 700° C. or higher to produce high-strength and high ductility near alpha titanium alloy. It can be said that aging treatment is more preferable.

On the other hand, it can be seen that the fracture surface of the alloy of Example 6 subjected to the tensile test at 650 ° C has numerous small and large dimples when the aging temperature is 650 ° C (Fig. 6 (c)) and 700 ° C (Fig. 6 (d)). . The presence of such dimples indicates a typical ductile failure mode, and the fracture surface results are in good agreement with the high high temperature elongation in FIG. 5 .

<Experimental Example 3> Analysis of the influence of Zr and Hf content in titanium alloy

In order to analyze the influence of the Zr and Hf content in the titanium alloy of the example, a tensile strength test was performed for Examples 7 and 8, and the results are shown in FIGS. 7 to 10 and Table 2 below, which are Zr And the remaining content other than the Hf content was compared with the same Example 6.

7 is a graph showing the results of a tensile test at 650° C. after solution heat treatment and aging treatment at 700° C. of the titanium alloy of Example 7 of the present invention.

8 is a graph showing the results of a tensile test at room temperature after solution heat treatment and aging treatment at 700° C. of the titanium alloy of Example 7 of the present invention.

9 is a graph showing the results of a tensile test at 650° C. after solution heat treatment and aging treatment at 700° C. of the titanium alloy of Example 8 of the present invention.

10 is a graph showing the results of a tensile test at room temperature after solution heat treatment and aging treatment at 700° C. of the titanium alloy of Example 8 of the present invention.

Specific results for this are shown in Table 2 below.

Sample (experiment execution temperature) Tensile strength (MPa) Yield strength (MPa) Elongation (%) speed rate Example 7 (room temperature) 1176.54 1084.48 7.25 0.001 Example 7 (650°C) 541.39 500.15 29 0.001 Example 8 (room temperature) 1191.78 - 0.96 0.001 Example 8 (650° C.) 455.68 428.43 18.4 0.001

Comparing the tensile strengths measured at high temperature of Examples 6 to 8, Example 8 was 454.68 MPa and Example 7 was 541.39 MPa, whereas Example 6 was 662.88 MPa (see FIG. 4), It can be seen that the case has higher tensile strength at high temperature.

Comparing the elongation measured at room temperature of Examples 6 to 8, Example 8 is 0.93%, Example 7 is 7.25%, whereas Example 6 is 16.94% (see FIG. 5), in the case of Example 6 It can be seen that has a higher elongation at room temperature.

Comparing the elongation measured at high temperature of Examples 6 to 8, Example 8 is 18.4%, Example 7 is 29%, whereas Example 6 is 61.96% (see FIG. 5), in the case of Example 6 It can be seen that has a significantly higher elongation at room temperature.

Under the same conditions where the content of other elements is the same, the content of Zr is large in the order of Example 8> Example 7> Example 6 improvement can be seen.

That is, it can be seen that the addition of Hf is effective in improving the room temperature and high temperature tensile properties of the titanium alloy.

Claims (12)

In weight %, Al: 5.5-7.0%, Sn: 2.5-4.0%, Hf: 2.5-5.5%, Nb: 0.1-1.5%, Mo: 0.2-0.6%, Si: 0.2-0.6%, B: 0.05- 0.2% C: 0.05 to 0.2%, other unavoidable impurities and the remainder Ti, near alpha titanium alloy containing no Zr.
delete delete delete delete delete According to claim 1,
wherein the alloy comprises silicide precipitates present at the alpha phase interface.
According to claim 1,
The alloy comprises a silicide precipitate present at the alpha phase interface,
A near alpha titanium alloy comprising a Ti-boride crystal.
According to claim 1,
The alloy is a near alpha titanium alloy, characterized in that it does not contain a _{regular α 2 phase.}
In weight %, Al: 5.5-7.0%, Sn: 2.5-4.0%, Hf: 2.5-5.5%, Nb: 0.1-1.5%, Mo: 0.2-0.6%, Si: 0.2-0.6%, B: 0.05- 0.2% C: 0.05 to 0.2%, including other unavoidable impurities and the remainder Ti, comprising the steps of casting a titanium alloy containing no Zr;
solution heat treating the cast alloy;
A method of producing a near alpha titanium alloy comprising; aging the solution heat treated alloy.
11. The method of claim 10,
The aging treatment is a method of producing a near alpha titanium alloy, characterized in that the aging treatment at 600 ~ 800 ℃.
11. The method of claim 10,
The aging treatment is a method of producing a near alpha titanium alloy, characterized in that the aging treatment at 700 ~ 800 ℃.

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