JP2968822B2 - Manufacturing method of high strength and high ductility β-type Ti alloy material - Google Patents
Manufacturing method of high strength and high ductility β-type Ti alloy materialInfo
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- JP2968822B2 JP2968822B2 JP19003290A JP19003290A JP2968822B2 JP 2968822 B2 JP2968822 B2 JP 2968822B2 JP 19003290 A JP19003290 A JP 19003290A JP 19003290 A JP19003290 A JP 19003290A JP 2968822 B2 JP2968822 B2 JP 2968822B2
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- phase temperature
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Description
【発明の詳細な説明】 ［産業上の利用分野］ 本発明は、微細な金属組織を有し、強度および延性の
優れたβ型Ti合金材を製造する方法に関するものであ
る。Description: TECHNICAL FIELD The present invention relates to a method for producing a β-type Ti alloy material having a fine metal structure and excellent strength and ductility.
［従来の技術］ β型Ti合金は、冷間加工が容易であり、且つ溶体化処
理後に時効処理を施してα相を析出させることにより高
強度を示すものとなる、といった特徴を有しているとこ
ろから、自動車用構造材料や航空機の機体用構造材料等
として需要な次第に増大してきている。[Prior art] β-type Ti alloys are characterized in that they are easily cold-worked and exhibit high strength by precipitating an α-phase by aging after solution treatment. Therefore, the demand for structural materials for automobiles, structural materials for aircraft bodies, and the like has been gradually increasing.
現在実用化されているβ型Ti合金としては、たとえば
Ti−13V−11Cr−3Al、Ti−15V−3Cr−3Sn−3Al、Ti−15
Mo−5Zr−3Al、Ti−3Al−8V−6Cr−4Mo−4Zr等が挙げら
れるが、これらのβ型Ti合金は前述の如く熱間加工後β
相温度域まで昇温して溶体化処理を行ない、その後時効
処理によって25％程度のα相を析出させて強化する方法
が採用されている。また高強度化を更に増進するための
手段として、β相温度域で溶体化処理した後冷間加工を
施して結晶内部に転位を導入し、次いで時効処理するこ
とにより微細なα相を析出させる方法も提案されてい
る。ところがこの様な方法で十分な強度を得るには、冷
間加工時に90％程度以上の強圧下を加えなければなら
ず、冷間加工によってその様な圧下率を与えることは工
業的に非常に困難であり、しかも得られる強圧下・時効
処理材は延性の乏しいものになるという問題があった。Currently available β-type Ti alloys include, for example,
Ti-13V-11Cr-3Al, Ti-15V-3Cr-3Sn-3Al, Ti-15
Mo-5Zr-3Al, Ti-3Al-8V-6Cr-4Mo-4Zr, etc., and these β-type Ti alloys after hot working as described above β
A method is adopted in which the temperature is raised to the phase temperature range, a solution treatment is performed, and then about 25% of the α phase is precipitated by aging treatment and strengthened. Further, as a means for further enhancing the strength, a solution treatment is performed in a β phase temperature range, then cold working is performed to introduce dislocations into the crystal, and then a fine α phase is precipitated by aging treatment. Methods have also been proposed. However, in order to obtain sufficient strength by such a method, it is necessary to apply a strong reduction of about 90% or more at the time of cold working, and it is extremely industrially difficult to provide such a reduction rate by cold working. There is a problem that it is difficult, and the obtained material under high pressure and aging becomes poor in ductility.
［発明が解決しようとする課題］ 本発明は上記の様な事情に着目してなされたものであ
って、その目的は、液体化処理後の強圧下およびそれに
伴なう延性低下の問題を解消し、高強度で且つ高延性の
β型Ti合金材を得ることのできる方法を確立しようとす
るものである。[Problems to be Solved by the Invention] The present invention has been made in view of the above-mentioned circumstances, and an object of the present invention is to solve the problem of high pressure after liquefaction treatment and the accompanying decrease in ductility. However, it is an object of the present invention to establish a method capable of obtaining a high-strength and high-ductility β-type Ti alloy material.
［課題を解決するための手段］ 上記課題を解決することのできた本発明の構成は、β
型Ti合金材に対し、熱間加工、溶体化処理、冷間加工、
時効処理を順次施すに当たり、上記溶体化処理を（α＋
β）２相温度域で行うところに要旨を有するものであ
る。この場合、溶体化処理前に行なわれる熱間圧延も
（α＋β）２相温度域で行なえば、時効処理後のβ型Ti
合金の結晶組織は更に微細なものとなり、強度および延
性の一段と良好なものを得ることができる。更に上記方
法において、Ti合金材が線材であるものは本発明の好ま
しい態様である。[Means for Solving the Problems] The configuration of the present invention that can solve the above-mentioned problems includes β
Hot working, solution treatment, cold working,
When sequentially performing the aging treatment, the solution treatment is performed by (α +
β) The gist is that the process is performed in the two-phase temperature range. In this case, if the hot rolling performed before the solution treatment is also performed in the (α + β) two-phase temperature range, the β-type Ti
The crystal structure of the alloy becomes finer, and more excellent strength and ductility can be obtained. Further, in the above method, the Ti alloy material being a wire is a preferred embodiment of the present invention.
［作用］ 前述の如く従来から実施されているβ型Ti合金材の加
工例では、冷間加工に先立って行われる溶体化処理を高
温のβ単相温度域で実施するのが常識とされていた。こ
れは、（α＋β）２相温度域で溶体化処理を行うと、初
折α相の存在によって延性が低下し冷間加工が困難にな
ると考えられていたからである。ところが、本発明者ら
が確認したところでは、（α＋β）２相温度域で溶体化
処理を行なった材料では、β相温度域溶体化材に比べて
僅かに強度が高く延性が低くなるものの、冷間加工性に
ついてはほとんど差がなく、従来と同様な強加工が可能
なことが分かった。[Operation] As described above, in the conventional processing example of the β-type Ti alloy material, it is common sense that the solution treatment performed prior to the cold working is performed in a high β single phase temperature range. Was. This is because it has been considered that when the solution treatment is performed in the (α + β) two-phase temperature range, the ductility is reduced due to the presence of the α phase, and the cold working becomes difficult. However, the present inventors have confirmed that the material subjected to the solution treatment in the (α + β) two-phase temperature range has a slightly higher strength and lower ductility than the β-phase temperature range solution-treated material. There was almost no difference in cold workability, and it was found that the same strong work as in the past could be performed.
さらに、本発明者らが研究したところでは、β型Ti合
金を高温のβ相温度域で溶体化処理すると結晶粒が粗大
化し易く、結晶粒を一旦粗大化させると、冷間加工で強
圧下を加えなければ時効処理後に析出α相を均一で微細
なものに分散させることが出来ず、強度および延性を満
足し得る程度まで高めることができなかった。しかし、
後記実施例で具体的に示す如く、溶体化処理を（α＋
β）２相温度域で行って少量の初析α相混入状態で冷間
加工し、次いで時効処理を行うと、ミクロ組織は極めて
均一でかつ微細なものとなり、強度および延性の非常に
優れたものとなることを知った。しかも上記の溶体化処
理に先立って行われる熱間加工についても、（α＋β）
２相温度域で実施すると、時効処理後のミクロ組織は一
段と均一かつ微細なものとなり、物性は更に改善される
ことをつきとめた。Furthermore, the present inventors have studied that, when the β-type Ti alloy is subjected to a solution treatment in a high β-phase temperature range, the crystal grains are likely to be coarsened, and once the crystal grains are coarsened, they are subjected to high pressure under cold working. Without adding, the precipitated α phase could not be dispersed uniformly and finely after the aging treatment, and the strength and ductility could not be increased to a satisfactory degree. But,
As specifically shown in Examples described later, the solution treatment was carried out by (α +
β) Cold working in a two-phase temperature range with a small amount of pro-eutectoid α phase mixed and then aging treatment, the microstructure becomes extremely uniform and fine, and the strength and ductility are extremely excellent. I knew it would be something. In addition, the hot working performed prior to the solution treatment is also (α + β)
It was found that, when the treatment was carried out in a two-phase temperature range, the microstructure after the aging treatment became more uniform and fine, and the physical properties were further improved.
こうした傾向が得られる理由は次のように考えること
ができる。即ち溶体化処理をβ相温度域よりも低温の
（α＋β）２相温度域で行うと、その後の冷間加工時
に、β相中に少量混在する初析α相との界面にも歪が生
じて転位が結晶全体に均一に導入され、それが析出サイ
トとなり、これを時効処理すると、無数に分布して形成
された上記析出サイトからα相の析出が起こり、均一か
つ微細なミクロ組織が得られるものと考えられる。The reason why this tendency is obtained can be considered as follows. That is, if the solution treatment is performed in the (α + β) two-phase temperature range lower than the β-phase temperature range, during the subsequent cold working, a strain also occurs at the interface with the pro-eutectoid α-phase mixed in a small amount in the β-phase. The dislocations are uniformly introduced into the entire crystal and become precipitation sites, and when this is aged, the α phase precipitates from the above-mentioned precipitation sites formed in innumerable distribution, and a uniform and fine microstructure is obtained. It is thought that it is possible.
さらに、上記溶体化処理に先立って行われる熱間加工
を（α＋β）２相温度域で行い、かつ溶体化処理も（α
＋β）２相温度域で行った場合には、熱間加工工程でも
少量の初析α相が生成し、更に溶体化処理工程でも初析
α相が生成してくるため、これら少量の初析α相が一層
均一に分布した溶体化処理材が得られ、その後の冷間加
工工程でより均一多数の析出サイトが形成されるととも
に、時効処理によりα相が一段と均一な分散状態で析出
してくるものと考えられる。また、（α＋β）２相温度
域で熱間加工し、かつ溶体化処理を施した材料では結晶
の成長が抑えられて結晶粒（β粒）が非常に小さいた
め、これも最終時効材のミクロ組織の均一微細化に好ま
しい影響を与えているものと思われる。Further, the hot working performed prior to the solution treatment is performed in the (α + β) two-phase temperature range, and the solution treatment is also performed at (α
+ Β) When performed in the two-phase temperature range, a small amount of pro-eutectoid α phase is generated even in the hot working step, and a pro-eutectoid α phase is also generated in the solution treatment step. A solution-treated material in which the α phase is more uniformly distributed is obtained, and in the subsequent cold working step, more and more precipitation sites are formed, and the aging treatment causes the α phase to precipitate in a more uniform dispersion state. It is thought to come. In the case of a material subjected to hot working and solution treatment in the (α + β) two-phase temperature range, crystal growth is suppressed and crystal grains (β grains) are very small. This seems to have a favorable effect on the uniform fineness of the structure.
いずれにしても本発明によれば、時効処理後のミクロ
組織を非常に微細なものとすることができ、高強度でし
かも高延性のものを得ることができる。In any case, according to the present invention, the microstructure after the aging treatment can be made very fine, and a material having high strength and high ductility can be obtained.
尚本発明を実施する際に採用される溶体化処理条件
は、前述の如くβ型Ti合金の種類に応じて（α＋β）２
相温度域の任意の温度に設定されるが、より好ましいの
はβトランザスよりも５〜150℃低温側の温度域であ
る。また熱間加工時の温度についても同様の温度範囲を
採用するのがよい。また溶体化処理後の冷間加工条件は
特に限定されないが、通常は30〜95％程度、必要な強度
に応じて加工率を増加させる。冷間加工後の時効処理は
言うまでもなく微細なα相を析出させて高強度化を果た
すために行われるものであり、通常は400〜600℃で10〜
1200分程度、より一般的には400〜500℃で60〜600分程
度の範囲が採用される。Note that the solution treatment conditions employed in carrying out the present invention are (α + β) 2 according to the type of β-type Ti alloy as described above.
The temperature is set to an arbitrary temperature in the phase temperature range, and more preferably a temperature range 5 to 150 ° C. lower than β Transus. It is preferable to adopt the same temperature range as the temperature during hot working. The cold working conditions after the solution treatment are not particularly limited, but are usually about 30 to 95%, and the working rate is increased according to the required strength. The aging treatment after cold working is, of course, performed in order to precipitate a fine α phase to achieve high strength, and is usually performed at 400 to 600 ° C. for 10 to 10 hours.
A range of about 1200 minutes, more generally about 400 to 500 ° C. and about 60 to 600 minutes is employed.
次に実施例を挙げて本発明を具体的に説明するが、本
発明はもとより下記実施例によって限定されるものでは
なく、前・後記の趣旨に適合し得る範囲で適当に変更し
て実施することも可能であり、それらはいずれも本発明
の技術的範囲に含まれる。Next, the present invention will be specifically described with reference to examples. However, the present invention is not limited to the following examples, and the present invention is appropriately modified and implemented within a range that can conform to the purpose of the preceding and the following. It is also possible and they are all included in the technical scope of the present invention.
［実施例］ Ti−15V−3Cr−3Sn−3AlおよびTi−15Mo−5Zr−3Alよ
りなるβ型Ti合金を、真空アーク溶解後鍛造および熱間
圧延して得た9.5mmφの線材を供試材として使用し、夫
々を第１表に示す条件で熱間加工（熱間スウェ−ジ）→
溶体化処理→冷間伸線→時効処理を順次行ない、約2mm
φの線材を得た。得られた各β型Ti合金線材の引張強度
および絞りを第1,2図に示す。尚これらの図には、溶体
化処理ののち圧下率80％で冷間加工し、次いで各合金に
つき２種類の温度（図中に示す）で８時間の時効処理を
行なったものの物性を示しており、比較のため溶体化処
理ままの物性も併記した。[Example] A 9.5 mmφ wire rod obtained by subjecting a β-type Ti alloy composed of Ti-15V-3Cr-3Sn-3Al and Ti-15Mo-5Zr-3Al to vacuum arc melting, forging and hot rolling was used. And hot working (hot swaging) under the conditions shown in Table 1
Solution treatment → cold drawing → aging treatment sequentially, about 2 mm
φ wire was obtained. FIGS. 1 and 2 show the tensile strength and drawing of each of the obtained β-type Ti alloy wires. These figures show the physical properties of the alloys, which were subjected to cold working at a rolling reduction of 80% after solution treatment, and then to aging treatment for two hours for each alloy at two different temperatures (shown in the figures). For comparison, the physical properties of the as-solution-treated solution are also shown.
尚第１表および第1,2図における各符号の意味は次の
通りである。The meanings of the respective symbols in Table 1 and FIGS. 1 and 2 are as follows.
△：β相温度域で熱間加工した後、β相温度域で溶体
化処理したもの（従来例）。Δ: hot-worked in the β-phase temperature range and then solution-treated in the β-phase temperature range (conventional example).
○：β相温度域で熱間加工した後、（α＋β）２相温
度域で溶体化処理したもの（本発明）。：: hot-worked in the β-phase temperature range and then solution-treated in the (α + β) two-phase temperature range (the present invention).
□：（α＋β）２相温度域で熱間加工した後、β相温
度域で溶体化処理したもの（比較例）。□: Hot-worked in the (α + β) two-phase temperature range and then solution-treated in the β-phase temperature range (comparative example).
◎：（α＋β）２相温度域で熱間加工した後、（α＋
β）２相温度域で溶体化処理したもの（本発明例）。：: After hot working in the (α + β) two-phase temperature range, (α + β)
β) Solution-treated in a two-phase temperature range (Example of the present invention).
第１表および第1,2図からも明らかである様に、溶体
化処理ままの物性を見ると、溶体化処理を（α＋β）２
相温度域で行なったもの（○，◎）の絞りは、β相温度
域で溶体化処理を行なったもの（□，△）に比べて悪
く、延性は前者の方が悪い。ところがこれらを圧下率80
％で冷間加工したのち時効処理したものの絞りを比較す
ると、上記の傾向は逆転し、溶体化処理を（α＋β）２
相温度域で行なったもの（○，◎）の方が明らかに高い
絞り率を示しており、高延性を示すことが分かる。尚、
引張強度については両者の間に殆んど差は認められな
い。 As is clear from Table 1 and FIGS. 1 and 2, when the physical properties of the solution treatment were observed, the solution treatment was performed using (α + β) 2
The squeezing performed in the phase temperature range (相, ◎) is worse than that performed in the solution treatment in the β phase temperature range (□, △), and the former has poorer ductility. However, these were reduced by 80
%, The above tendency is reversed, and the solution treatment is carried out by (α + β) 2
The test performed in the phase temperature range (，, ◎) clearly shows a higher drawing ratio, which indicates that the test shows higher ductility. still,
There is almost no difference in tensile strength between the two.
特に第１図の結果を見ると、Ti−15V−3Cr−3Sn−3Al
合金を使用し、溶体化処理を（α＋β）２相温度域で行
なったものは、その後の例間加工および時効処理でα相
を析出させた場合でも延性は殆んど低下せず、高強度で
高延性を示すものになることが分かる。また第２図のTi
−15Mo−5Zr−3Al合金を用いた実験例では、溶体化処理
をβ相温度域あるいは（α＋β）２相温度域のどちらで
行なった場合でも、冷間加工および時効処理後の絞りは
かなり低下するが、その低下傾向は（α＋β）２相温度
域で溶体化処理を行なったもの（○，◎）の方が緩やか
であり、冷間加工および時効処理後の絞り（％）はβ相
温度域で溶体化処理を行なったものより高い値を示して
いる。In particular, looking at the results in FIG. 1, it can be seen that Ti-15V-3Cr-3Sn-3Al
In the case of using an alloy and performing the solution treatment in the (α + β) two-phase temperature range, the ductility hardly decreases even when the α phase is precipitated by inter-working and aging treatment, and high strength It can be seen that the sample shows high ductility. In addition, in FIG.
In the experimental example using -15Mo-5Zr-3Al alloy, the reduction after cold working and aging treatment is considerably reduced regardless of whether solution treatment is performed in β phase temperature range or (α + β) two phase temperature range. However, the tendency of the decrease is slower in the case where the solution treatment is performed in the (α + β) two-phase temperature range (○, ◎), and the reduction (%) after the cold working and the aging treatment is the β phase temperature. The values are higher than those obtained by solution treatment in the region.
また第3,4,5図は、Ti−15Mo−5Zr−3Alを従来法およ
び本発明法により処理して得た金属組織を示す図面代用
写真であり、夫々下記の条件で処理したものである。FIGS. 3, 4, and 5 are drawing substitute photographs showing the metal structures obtained by treating Ti-15Mo-5Zr-3Al by the conventional method and the method of the present invention, respectively, which were processed under the following conditions. .
第３図（従来冷）：熱間圧延線材（9.5mmφ）→スウ
ェージ（75％−850℃：β相温度域）→溶体化処理（835
℃×15分：β相温度域）→冷間伸線（80％）→時効処理
（500℃×８時間） 第４図（本発明例１）：熱間圧延線材（9.5mmφ）→
熱間スウェージ（75％−850℃：β相温度域）→溶体化
処理（735℃×１時間：（α＋β）２相温度域）→冷間
伸線（80％）→時効処理（500℃×８時間） 第５図（本発明例２）：熱間圧延線材（9.5mmφ）→
熱間スウェージ（75％−700℃：（α＋β）２相温度
域）→溶体化処理（735℃×１時間：（α＋β）２相温
度域）→冷間伸線（80％）→時効処理（500℃×８時
間） 第３〜５図からも明らかである様に、溶体化処理をβ
相温度域で行なった従来例（第３図）では結晶粒が粗大
であるが、溶体化処理を（α＋β）２相温度域で行なっ
た本発明例１（第４図）では結晶粒が著しく微細化して
おり、更に熱間スウェージと溶体化処理をいずれも（α
＋β）２相温度域で行なった本発明例２（第５図）で
は、結晶粒は一段と微細になると共に組織が極めて均質
になっていることが分かる。Fig. 3 (conventionally cold): hot-rolled wire (9.5 mmφ) → swage (75% -850 ° C: β-phase temperature range) → solution treatment (835)
° C × 15 minutes: β phase temperature range) → cold drawing (80%) → aging treatment (500 ° C × 8 hours) Fig. 4 (Example 1 of the present invention): hot-rolled wire (9.5mmφ) →
Hot swaging (75% -850 ° C: β phase temperature range) → solution treatment (735 ° C x 1 hour: (α + β) two phase temperature range) → cold drawing (80%) → aging treatment (500 ° C x 8 hours) FIG. 5 (Example 2 of the present invention): hot-rolled wire (9.5 mmφ) →
Hot swaging (75% -700 ° C: (α + β) two-phase temperature range) → solution treatment (735 ° C × 1 hour: (α + β) two-phase temperature range) → cold drawing (80%) → aging treatment ( (500 ° C. × 8 hours) As is clear from FIGS.
In the conventional example (FIG. 3) performed in the phase temperature range, the crystal grains were coarse, but in the present invention example 1 (FIG. 4) in which the solution treatment was performed in the (α + β) two phase temperature range, the crystal grains were remarkably large. It has been miniaturized, and both hot swaging and solution treatment (α
+ Β) In Example 2 of the present invention (FIG. 5) performed in the two-phase temperature range, it can be seen that the crystal grains are further refined and the structure is extremely homogeneous.
この様に本発明によれば時効処理後の結晶粒を著しく
微細化し得ると共に組織を均質化することができ、それ
により高強度化と高延性化が達成されたものと考えられ
る。As described above, according to the present invention, it is considered that the crystal grains after the aging treatment can be remarkably refined and the structure can be homogenized, thereby achieving high strength and high ductility.
［発明の効果］ 本発明は以上の様に構成されており、溶体化処理を
（α＋β）２相温度域で行ない、あるいは溶体化処理と
その前の熱間加工を共に（α＋β）２相温度域で実施す
ることにより、その後冷間加工および時効処理を行なっ
た後の結晶粒を著しく微細化すると共に組織を均質化す
ることができ、高強度で高延性のβ型Ti合金材を提供し
得ることになった。[Effects of the Invention] The present invention is configured as described above, and the solution treatment is performed in the (α + β) two-phase temperature range, or both the solution treatment and the hot working before the solution treatment are performed in the (α + β) two-phase temperature range. By performing it in the region, it is possible to remarkably refine the crystal grains after cold working and aging treatment and homogenize the structure, and to provide a high-strength, high-ductility β-type Ti alloy material. I got it.
第1,2図は実施例で得たβ型Ti合金材の引張強さと絞り
を示すグラフ、第３〜５図は従来例および本発明例で得
たβ型Ti合金材の金属組織を示す図面代用写真である。FIGS. 1 and 2 are graphs showing the tensile strength and drawing of the β-type Ti alloy material obtained in the examples, and FIGS. It is a drawing substitute photograph.
───────────────────────────────────────────────────── フロントページの続き (51)Int.Cl.6 識別記号 ＦＩ Ｃ２２Ｆ 1/00 ６８３ Ｃ２２Ｆ 1/00 ６８３ ６８４ ６８４Ｃ ６８５ ６８５Ｚ ６８６ ６８６Ａ ６９１ ６９１Ｂ ６９４ ６９４Ｂ (56)参考文献 特開 平１−28348（ＪＰ，Ａ) 特開 平１−279736（ＪＰ，Ａ) 特開 昭61−204358（ＪＰ，Ａ) 東京大学工学部総合試験所年報 Ｖｏ ｌ，46 ｐｐ．197−202（1987) (58)調査した分野(Int.Cl.6，ＤＢ名) C22F 1/18 C22C 14/00 ────────────────────────────────────────────────── ─── front page continued (51) Int.Cl. 6 identifications FI C22F 1/00 683 C22F 1/00 683 684 684C 685 685Z 686 686A 691 691B 694 694B (56) references Patent Rights 1-28348 ( JP, A) JP-A-1-279736 (JP, A) JP-A-61-204358 (JP, A) The University of Tokyo Faculty of Engineering Annual Report Vol, 46 pp. 197-202 (1987) (58) Fields investigated (Int. Cl. 6 , DB name) C22F 1/18 C22C 14/00
Claims (3)
理、冷間加工、時効処理を順次施すに当たり、 上記溶体化処理を（α＋β）２相温度域で行うことを特
徴とする高強度・高延性β型Ti合金材の製法。The present invention is characterized in that, when sequentially performing hot working, solution treatment, cold working, and aging treatment on a β-type Ti alloy material, the solution treatment is performed in an (α + β) two-phase temperature range. Of high strength and high ductility β-type Ti alloy material.
求項（１）記載の製法。2. The method according to claim 1, wherein the hot working is performed in a (α + β) two-phase temperature range.
たは（２）記載の製法。3. The method according to claim 1, wherein the Ti alloy material is a wire.
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| JP19003290A JP2968822B2 (en) | 1990-07-17 | 1990-07-17 | Manufacturing method of high strength and high ductility β-type Ti alloy material |
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|---|---|---|---|
| JP19003290A JP2968822B2 (en) | 1990-07-17 | 1990-07-17 | Manufacturing method of high strength and high ductility β-type Ti alloy material |
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| Publication Number | Publication Date |
|---|---|
| JPH0474856A JPH0474856A (en) | 1992-03-10 |
| JP2968822B2 true JP2968822B2 (en) | 1999-11-02 |
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| JP4562830B2 (en) * | 1999-09-10 | 2010-10-13 | トクセン工業株式会社 | Manufacturing method of β titanium alloy fine wire |
| JP3934372B2 (en) * | 2001-08-15 | 2007-06-20 | 株式会社神戸製鋼所 | High strength and low Young's modulus β-type Ti alloy and method for producing the same |
| US20040221929A1 (en) | 2003-05-09 | 2004-11-11 | Hebda John J. | Processing of titanium-aluminum-vanadium alloys and products made thereby |
| US7837812B2 (en) | 2004-05-21 | 2010-11-23 | Ati Properties, Inc. | Metastable beta-titanium alloys and methods of processing the same by direct aging |
| US10053758B2 (en) | 2010-01-22 | 2018-08-21 | Ati Properties Llc | Production of high strength titanium |
| US9255316B2 (en) | 2010-07-19 | 2016-02-09 | Ati Properties, Inc. | Processing of α+β titanium alloys |
| US8499605B2 (en) | 2010-07-28 | 2013-08-06 | Ati Properties, Inc. | Hot stretch straightening of high strength α/β processed titanium |
| US9206497B2 (en) | 2010-09-15 | 2015-12-08 | Ati Properties, Inc. | Methods for processing titanium alloys |
| US8613818B2 (en) | 2010-09-15 | 2013-12-24 | Ati Properties, Inc. | Processing routes for titanium and titanium alloys |
| US10513755B2 (en) | 2010-09-23 | 2019-12-24 | Ati Properties Llc | High strength alpha/beta titanium alloy fasteners and fastener stock |
| US8652400B2 (en) | 2011-06-01 | 2014-02-18 | Ati Properties, Inc. | Thermo-mechanical processing of nickel-base alloys |
| US9050647B2 (en) | 2013-03-15 | 2015-06-09 | Ati Properties, Inc. | Split-pass open-die forging for hard-to-forge, strain-path sensitive titanium-base and nickel-base alloys |
| US9869003B2 (en) | 2013-02-26 | 2018-01-16 | Ati Properties Llc | Methods for processing alloys |
| US9192981B2 (en) | 2013-03-11 | 2015-11-24 | Ati Properties, Inc. | Thermomechanical processing of high strength non-magnetic corrosion resistant material |
| US9777361B2 (en) | 2013-03-15 | 2017-10-03 | Ati Properties Llc | Thermomechanical processing of alpha-beta titanium alloys |
| US11111552B2 (en) | 2013-11-12 | 2021-09-07 | Ati Properties Llc | Methods for processing metal alloys |
| US10094003B2 (en) | 2015-01-12 | 2018-10-09 | Ati Properties Llc | Titanium alloy |
| US10502252B2 (en) | 2015-11-23 | 2019-12-10 | Ati Properties Llc | Processing of alpha-beta titanium alloys |
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1990
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Non-Patent Citations (1)
| Title |
|---|
| 東京大学工学部総合試験所年報 Ｖｏｌ，４６ ｐｐ．１９７−２０２（１９８７） |
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