JPWO2002070763A1 - Titanium alloy bar and method of manufacturing the same - Google Patents
Titanium alloy bar and method of manufacturing the same Download PDFInfo
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- JPWO2002070763A1 JPWO2002070763A1 JP2002570785A JP2002570785A JPWO2002070763A1 JP WO2002070763 A1 JPWO2002070763 A1 JP WO2002070763A1 JP 2002570785 A JP2002570785 A JP 2002570785A JP 2002570785 A JP2002570785 A JP 2002570785A JP WO2002070763 A1 JPWO2002070763 A1 JP WO2002070763A1
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- 229910001069 Ti alloy Inorganic materials 0.000 title claims abstract description 38
- 238000004519 manufacturing process Methods 0.000 title claims description 14
- 238000005096 rolling process Methods 0.000 claims abstract description 121
- 239000013078 crystal Substances 0.000 claims abstract description 30
- 239000010936 titanium Substances 0.000 claims abstract description 6
- 239000000463 material Substances 0.000 claims description 45
- 229910045601 alloy Inorganic materials 0.000 claims description 19
- 239000000956 alloy Substances 0.000 claims description 19
- 230000009466 transformation Effects 0.000 claims description 15
- 238000005098 hot rolling Methods 0.000 claims description 14
- 238000010438 heat treatment Methods 0.000 claims description 13
- 229910001040 Beta-titanium Inorganic materials 0.000 claims 1
- 230000007423 decrease Effects 0.000 description 9
- 229910000883 Ti6Al4V Inorganic materials 0.000 description 8
- 238000000034 method Methods 0.000 description 8
- 238000010586 diagram Methods 0.000 description 7
- 238000005242 forging Methods 0.000 description 5
- 238000001816 cooling Methods 0.000 description 4
- 238000005336 cracking Methods 0.000 description 4
- 238000007788 roughening Methods 0.000 description 4
- 239000000126 substance Substances 0.000 description 4
- 238000009864 tensile test Methods 0.000 description 4
- 230000020169 heat generation Effects 0.000 description 3
- 239000002344 surface layer Substances 0.000 description 3
- 238000009661 fatigue test Methods 0.000 description 2
- 238000003303 reheating Methods 0.000 description 2
- 230000035945 sensitivity Effects 0.000 description 2
- 230000003746 surface roughness Effects 0.000 description 2
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 2
- 206010013786 Dry skin Diseases 0.000 description 1
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 description 1
- RTAQQCXQSZGOHL-UHFFFAOYSA-N Titanium Chemical compound [Ti] RTAQQCXQSZGOHL-UHFFFAOYSA-N 0.000 description 1
- 238000010521 absorption reaction Methods 0.000 description 1
- 238000000137 annealing Methods 0.000 description 1
- 230000015572 biosynthetic process Effects 0.000 description 1
- 238000005260 corrosion Methods 0.000 description 1
- 230000007797 corrosion Effects 0.000 description 1
- 238000009792 diffusion process Methods 0.000 description 1
- 238000005516 engineering process Methods 0.000 description 1
- 238000009499 grossing Methods 0.000 description 1
- 229910052739 hydrogen Inorganic materials 0.000 description 1
- 239000001257 hydrogen Substances 0.000 description 1
- 238000002844 melting Methods 0.000 description 1
- 230000008018 melting Effects 0.000 description 1
- 239000002245 particle Substances 0.000 description 1
- 238000005204 segregation Methods 0.000 description 1
- 230000000087 stabilizing effect Effects 0.000 description 1
- 230000001629 suppression Effects 0.000 description 1
- 229910052719 titanium Inorganic materials 0.000 description 1
- 238000009721 upset forging Methods 0.000 description 1
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
- C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
- C22F1/16—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of other metals or alloys based thereon
- C22F1/18—High-melting or refractory metals or alloys based thereon
- C22F1/183—High-melting or refractory metals or alloys based thereon of titanium or alloys based thereon
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C14/00—Alloys based on titanium
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
- B21B—ROLLING OF METAL
- B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
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- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Metal Rolling (AREA)
Abstract
本発明は、実質的に、質量％で、Ａｌ：４％以上５％以下、Ｖ：２．５％以上３．５％以下、Ｆｅ：１．５％以上２．５％以下、Ｍｏ：１．５％以上２．５％以下、残部：Ｔｉからなり、かつ初析α相の体積分率が１０％以上９０％以下、初析α相の平均結晶粒径が１０μｍ以下、圧延方向に平行な断面における初析α相の結晶粒のアスペクト比が４以下であるα＋β型チタン合金棒材に関する。本発明のα＋β型チタン合金棒材は、延性、疲労特性および加工性に優れる。In the present invention, Al: 4% or more and 5% or less, V: 2.5% or more and 3.5% or less, Fe: 1.5% or more and 2.5% or less, Mo: 1 by mass%. 0.5% or more and 2.5% or less, balance: Ti, volume fraction of proeutectoid α phase is 10% or more and 90% or less, average crystal grain size of proeutectoid α phase is 10 μm or less, parallel to the rolling direction The present invention relates to an α + β-type titanium alloy rod having an aspect ratio of crystal grains of a pro-eutectoid α phase in a simple cross section of 4 or less. The α + β type titanium alloy bar of the present invention is excellent in ductility, fatigue properties and workability.
Description
技術分野
本発明は、延性、疲労特性および加工性に優れたチタン合金棒材、特に、α＋β型チタン合金棒材およびその製造方法に関する。
背景技術
チタン合金は、高強度で、軽く、その上耐食性にも優れているので、化学プラント、発電機、航空機などの分野で構造用材料として用いられている。なかでも、高強度であり、比較的良好な加工性を備えているα＋β型チタン合金が多用されている。
チタン合金を用いた製品には、薄板、厚板や棒材など様々な形状のものがある。棒材には、そのままの形状で使用される場合もあるが、ボルトのネジ部のように複雑な形状に加工されたり、鍛造素材として使用される場合もあるので、棒材自体の優れた延性や疲労特性の他に、優れた加工性も要求される。
図１に、棒材の代表的な製造方法を示す。
溶解により製造されたインゴットは、鍛造によりビレットに加工され圧延素材となる。図２Ａ、２Ｂに示すように、圧延は、ビレットを加熱炉で加熱後、リバース型圧延機あるいはタンデム型圧延機により行われ、必要に応じて中間加熱炉により圧延可能な温度まで再加熱される。
しかし、チタン合金棒材、特に、α＋β型チタン合金棒材では、加工発熱により熱間圧延中に被圧延材の温度が上昇するため、安定した熱間圧延ができず、延性、疲労特性、加工性ともに優れたチタン合金棒材が得られないのが実状である。例えば、被圧延材の温度がβ変態点以上まで上昇すると、針状のα相が主体のβ組織となるため、優れた延性や疲労特性が得られない。また、β変態点が高いため、加工発熱により被圧延材の温度がβ変態点以上になり難いＴｉ−６Ａｌ−４Ｖ合金の場合でも、加工発熱により圧延温度が上昇すると粒成長が促進され、優れた延性、疲労特性および加工性が得られない。
加工発熱による温度上昇の問題を解決する方法として、特開昭５９−８２１０１号公報には、α域温度およびα＋β域温度での１圧延パス当たりの断面減少率を４０％以下にする圧延方法が開示されている。また特開昭５８−２５４６５号公報には、加工発熱による温度上昇を抑えるために被圧延材を水冷する方法が開示されている。さらに、論文１：「Ｈｏｔ Ｂａｒ Ｒｏｌｌｉｎｇ ｏｆ Ｔｉ−６Ａｌ−４Ｖ ｉｎ ａ Ｃｏｎｔｉｎｕｏｕｓ Ｍｉｌｌ（Ｔｉｔａｎｉｕｍ’９２ Ｓｃｉｅｎｃｅ ａｎｄ Ｔｅｃｈｎｏｌｏｇｙ）」には、加工発熱自体を抑制するために圧延速度を使用する圧延機の性能限界まで低速化することが記載されている。
しかしながら、特開昭５９−８２１０１号公報および特開昭５８−２５４６５号公報の方法では、延性、疲労特性および加工性ともに優れたチタン合金棒材が得られない。
また、特開昭５９−８２１０２号公報の方法により１圧延パス当たり４０％以下の断面減少率としても、合金の種類によっては加工発熱の抑制が不充分の場合もある。さらに、特開昭５８−２５４６５号公報の方法では、水冷しているため水素吸収による材質劣化が生じたり、急速冷却による歪みのために正確な温度制御が困難になるという問題もある。
論文１に記載の方法は、対象がＴｉ−６Ａｌ−４Ｖ合金であって、以下に示すように、加工発熱が大きくより低温域で圧延される合金に対しては必ずしも当てはまらず、優れた延性、疲労特性および加工性が得られない。
図３に、Ｔｉ−６Ａｌ−４Ｖ合金およびＴｉ−４．５Ａｌ−３Ｖ−２Ｆｅ−２Ｍｏ合金の棒材圧延における被圧延材の温度と圧延時間の関係を示す。
ここで、加熱温度は、Ｔｉ−６Ａｌ−４Ｖ合金では９５０℃、β変態点がＴｉ−６Ａｌ−４Ｖ合金より１００℃低いＴｉ−４．５Ａｌ−３Ｖ−２Ｆｅ−２Ｍｏ合金ではβ変態点の差の分だけ下げて８５０℃とした。圧延は、リバース型圧延機およびタンデム型圧延機を用いて行い、圧延速度、圧下率、パススケジュールは両合金とも同じにした。リバース圧延機での圧延速度は両合金とも２．７ｍ／ｓｅｃ、タンデム圧延機での圧延速度は最も速くなる最終圧延パスで両合金とも２．２５ｍ／ｓｅｃとした。なお、この圧延速度は上記の論文１に記載の圧延速度（６ｍ／ｓｅｃ）よりさらに低い速度である。断面減少率は、両合金ともに最大で２６％とした。
Ｔｉ−６Ａｌ−４Ｖ合金の圧延の場合には、この合金のβ変態点である１０００℃より充分に低温域で圧延が行われ、良好な組織が得られる。一方、Ｔｉ−４．５Ａｌ−３Ｖ−２Ｆｅ−２Ｍｏ合金の場合には、β変態点が低い分だけ加熱温度を低下させたにもかかわらず、低温圧延により変形抵抗が増大して加工発熱が大きくなり、β変態点を超える温度域まで被圧延材の温度が上昇して良好な組織が得られない。したがって、優れた延性、疲労特性および加工性が得られないことになる。このことは、圧延速度のみならず、圧延温度、圧下率、圧延パス間時間などの圧延条件を考慮する必要のあることを示唆している。
発明の開示
本発明の目的は、延性、疲労特性および加工性に優れた高強度チタン合金棒材およびその製造方法を提供することにある。
この目的は、実質的に、質量％で、Ａｌ：４％以上５％以下、Ｖ：２．５％以上３．５％以下、Ｆｅ：１．５％以上２．５％以下、Ｍｏ：１．５％以上２．５％以下、残部：Ｔｉからなり、かつ初析α相の体積分率が１０％以上９０％以下、初析α相の平均結晶粒径が１０μｍ以下、圧延方向に平行な断面における初析α相の結晶粒のアスペクト比が４以下であるα＋β型チタン合金棒材によって達成される。
このα＋β型チタン合金棒材は、上記成分のチタン合金を、その表面温度が常にβ変態点以下になるように熱間圧延する工程を有するα＋β型チタン合金棒材の製造方法により製造できる。
発明の実施の形態
まず、本発明者等は、α＋β型チタン合金棒材の組織をどのような組織にすれば、優れた延性、疲労特性および加工性が得られるかを検討した。その結果、以下のことが明らかになった。
α＋β型チタン合金は初析α相と変態β相から成るが、すべり系が少ないＨＣＰ構造を有するα相の体積分率が極めて多くなったり、針状α相を内部に有する変態β相の体積分率が極めて多くなっても加工性や延性が低下するので、初析α相の体積分率は１０％以上９０％以下とする。なお、熱間加工の加熱時にα相とβ相との体積分率が等しいか、あるいは近い場合にはさらに加工性は良好となるので、初析α相の体積分率は５０％以上８０％以下が望ましい。
図４に、初析α相の平均結晶粒径と高温引張試験による全伸びとの関係を示す。
初析α相の平均結晶粒径が１０μｍを超えると、高温引張試験による全伸びが急激に低下する。また、それにともない加工性が低下する。
図５に、初析α相の平均結晶粒径と疲労試験における１０８回での疲労強度との関係を示す。
初析α相の平均結晶粒径が１０μｍを超えると、疲労強度が低下する。また、初析α相の平均結晶粒径が６μｍより小さいと、より高い疲労強度が得られる。
棒材を鍛造する場合、金型に接しない自由変形面では結晶粒の形状、すなわち結晶粒のアスペクト比に起因して肌荒れが発生する。一般に棒材の場合は、結晶粒が圧延方向に展伸する傾向にある。特に、アップセット鍛造のような場合には、展伸した組織が自由変形面となる棒材の側面表面に現れるため、鍛造後の製品の肌荒れを防ぐには、鍛造時のアスペクト比が大きくなり過ぎないように、具体的には棒材の圧延方向に平行な断面での初析α相の結晶粒のアスペクト比を４以下にする必要がある。
以上のことから、初析α相の体積分率を１０％以上９０％以下、より好ましくは５０％以上８０％以下、初析α相の平均結晶粒径を１０μｍ以下、より好ましくは６μｍ以下、さらに圧延方向に平行な断面での初析α相の結晶粒のアスペクト比を４以下にすれば、延性、疲労特性および加工性に優れた高強度チタン合金棒材が得られる。
このような組織を有するα＋β型チタン合金棒材は、実質的に、質量％で、Ａｌ：４％以上５％以下、Ｖ：２．５％以上３．５％以下、Ｆｅ：１．５％以上２．５％以下、Ｍｏ：１．５％以上２．５％以下、残部：Ｔｉからなる成分とする必要がある。各元素の含有量の限定理由を以下に説明する。
Ａｌ：α相を安定化させるのに必須の元素であり、高強度化にも寄与する元素である。４％未満では高強度が十分に達成されず、５％を超えると延性が劣化する。
Ｖ：β相を安定化させる元素であり、高強度化にも寄与する元素である。２．５％未満では高強度が十分に達成されないとともに、β相が安定せず、３．５％を超えるとβ変態点の低下により加工温度領域が狭くなることに加え、高コスト化を招く。
Ｍｏ：β相を安定化させる元素であり、高強度化にも寄与する元素である。１．５％未満では高強度化が十分に達成されないとともに、β相が安定せず、２．５％を超えるとβ変態点の低下により加工温度領域が狭くなることに加え、高コスト化を招く。
Ｆｅ：β相を安定化させる元素であり、高強度化にも寄与する元素である。また、拡散速度が速く加工性を改善する効果を有するが、１．５％未満では高強度化が十分に達成されないとともに、β相が安定せず、優れた加工性が得られない。２．５％を超えるとβ変態点の低下により加工温度領域が狭くなることに加え、偏析による特性の劣化を招く。
本発明のα＋β型チタン合金棒材は、上記の成分のα＋β型チタン合金を、加熱温度、圧延温度域、圧下率、圧延速度、パス間時間などの圧延条件を調整して加工発熱による温度上昇を抑制しながら、その表面温度が常にβ変態点以下になるように熱間圧延する方法により製造できる。例えば、β変態点がＴβ℃であるα＋β型チタン合金を、表面温度が（Ｔβ−１５０）℃以上Ｔβ℃以下の範囲になるように加熱する工程と、加熱されたα＋β型チタン合金を、圧延中の被圧延材の表面温度が（Ｔβ−３００）℃以上（Ｔβ−５０）℃以下の範囲に、仕上表面温度が（Ｔβ−３００）℃以上（Ｔβ−１００）℃以下の範囲になるように熱間圧延する工程を有する方法である。
ここで、圧延前に表面温度を（Ｔβ−１５０）℃以上Ｔβ℃以下の範囲に加熱する理由は、（Ｔβ−１５０）℃未満だと圧延最終段階における被圧延材の温度低下が大きくなり割れ感受性や変形抵抗の上昇を引き起こし、Ｔβ℃を超えると被圧延材の組織が針状αを主体とするβ組織となり延性や加工性が劣化するためである。また、圧延中の被圧延材の表面温度を（Ｔβ−３００）℃以上（Ｔβ−５０）℃以下の範囲にする理由は、（Ｔβ−３００）℃未満だと熱間加工性が低下し割れなどの問題が生じ、（Ｔβ−５０）℃を超えると加工発熱による温度上昇で結晶粒の粗大化や針状組織の形成を招くためである。さらに、被圧延材が最終圧延パスを終えた直後の表面温度である仕上表面温度を（Ｔβ−３００）℃以上（Ｔβ−１００）℃以下の範囲にする理由は、（Ｔβ−３００）℃未満だと割れ感受性や変形抵抗が上昇し、（Ｔβ−１００）℃を超えると結晶粒の粗大化を招くためである。
熱間圧延は複数回の圧延パスにより行われるが、加工発熱による温度上昇を防ぐために、１圧延パス当りの圧下率を４０％以下にすることが好ましい。
熱間圧延をリバース型圧延機を用いて行うとき、加工発熱による温度上昇を防ぐために、圧延速度を６ｍ／ｓｅｃ以下にすることが好ましい。また、タンデム型圧延機を用いて行うときは、圧延速度を１．５ｍ／ｓｅｃ以下にすることが好ましい。
各圧延パス後、被圧延材は表面から冷却されるので、加工発熱による温度上昇があっても、被圧延材表層部では次の圧延パスまでにある程度の温度低下がある。しかし、図６に示すように、被圧延材の径が大きい場合は（直径１０６ｍｍの場合）被圧延材の中心部での温度低下が小さいため、被圧延材の表層部と中心部で大きな温度差が現れる。中心部の温度低下が小さい場合は、中心部の温度が低下する前に次の圧延パスを受け、加工発熱によってさらに温度が上昇する。この現象が続けば、中心部は初期温度より高温で圧延されることになる。このため、被圧延材の径が大きい場合は、十分な圧延パス間時間を取って中心部を冷却する必要がある。
そこで、本発明者らが表層部と中心部の温度差について詳細に検討したところ、図７に示すように、温度差は被圧延材の圧延方向と直角方向の断面積が３５００ｍｍ２以上で顕著に大きくなり、そのような大きな断面積を有する被圧延材を断面積がＳｍｍ２となるように圧延したとき、次の圧延を開始するまでの時間を０．１６７×Ｓ１／２ｓｅｃ以上にすることが温度差を小さくでき、均一な特性の棒材を製造する上で好ましいことを見出した。
なお、本発明の製造方法においては、被圧延材の表面温度が常にβ変態点以下になるように圧延するため、圧延パス間時間や被圧延材の径によっては被圧延材の表面温度が圧延中に適切な圧延温度域より低温側に外れる可能性がある。そのような場合は、高周波加熱設備などにより再加熱することも可能である。
実施例１
表１に示す化学成分のα＋β型チタン合金Ａ０１（本発明範囲内）およびＡ０２（本発明範囲外）より１２５ｍｍ角の圧延素材を切り出し、カリバー圧延機を用いて表２に示す圧延条件Ｂ０１−Ｂ１８で直径２０ｍｍまたは５０ｍｍの棒材を製造した。表２における圧延パス間時間は、各圧延条件において、全ての圧延パスで圧延パス間時間が０．１６７×Ｓ１／２ｓｅｃ以上になっている場合を○で、なっていない場合を×で示してある。また、各圧延条件におけるそれぞれの圧延パスでの被圧延材の断面積Ｓ、圧下率、０．１６７×Ｓ１／２、圧延パス間時間、表面温度および圧延速度を表３−表２０に示した。表中の圧延設備欄のＲはリバース圧延機、Ｔはタンデム圧延機を表す。
製造した棒材を７００℃以上７２０℃以下で焼鈍後引張試験を行い、降伏強度（０．２％ＰＳ）、引張強度（ＵＴＳ）、のび（Ｅｌ）、絞り（ＲＡ）を測定した。また、平滑試験（条件：Ｋｔ＝１）および切欠試験（条件：Ｋｔ＝３）を行い、疲労強度を測定した。さらに、微細組織を調べるために、棒材の中心部および直径の１／４位置（１／４Ｄ）のミクロ組織の観察を行い、初析α相の結晶粒径、その体積分率および圧延方向に平行な断面での結晶粒のアスペクト比を測定した。
結果を表２１に示す。表のミクロ組織で結晶粒径の記載がないところは、その部位が針状αを主体とするβ組織のみからなり、等軸の初析α相を観測できなかったためである。
表面の加熱温度が（Ｔβ−１５０）℃未満の場合は、被圧延材の表面温度が低過ぎ、圧延荷重が超過して圧延できなかった。また、加熱温度がＴβ℃より高い場合は、圧延条件Ｂ０２およびＢ１１のように、圧延パス間時間が本発明範囲内であっても被圧延材の表面温度が高過ぎるため、加工発熱によって被圧延材の表面温度がＴβを超え、中心部の組織が針状αを主体とするβ組織となり、延性や疲労特性が劣化する。
仕上表面温度が（Ｔβ−３００）℃未満の場合は、被圧延材の温度が低過ぎ、加工性が低下して圧延中に割れが生じる。また、仕上表面温度が（Ｔβ−１００）℃より高い場合は、圧延条件Ｂ０４、Ｂ０５およびＢ０７のように、組織の微細化が図れず延性や疲労特性が劣化する。
圧延中の被圧延材の表面温度が（Ｔβ−３００）℃未満の場合は、被圧延材の表面温度が低過ぎ、圧延中に割れが生じる。また、被圧延材の表面温度が（Ｔβ−５０）℃より高い場合は、圧延条件Ｂ０２−Ｂ０５、Ｂ０７およびＢ１１のように、圧延後の中心部や１／４Ｄ部の組織が針状αを主体とするβ組織となり延性や疲労特性が劣化する。
１圧延パスの圧下率が４０％を超えると、加工発熱が大きくなり、被圧延材の温度がＴβを超え組織の微細化が図れなかった。
圧延条件Ｂ１４のようにリバース型圧延機を用い、圧延速度が６ｍ／ｓｅｃを超えると、また圧延条件Ｂ１５のようにタンデム型圧延機を用い、圧延速度が１．５ｍ／ｓｅｃを超えると、加工発熱が大きくなり、被圧延材の表面温度がＴβを超え、組織の微細化が図れなかった。
圧延パス間時間が本発明範囲外の場合は、加工発熱による被圧延材の表面温度上昇が放冷による温度低下に勝り、被圧延材の表面温度がＴβを超え、組織の微細化が図れなかった。
化学成分が本発明範囲内のＡ０１を用い、圧延条件Ｂ０１、Ｂ０６、Ｂ０８、Ｂ０９、Ｂ１６、Ｂ１７およびＢ１８によって製造した棒材では、初析α相の結晶粒径が１０μｍ以下の均一な微細組織が観察され、優れた延性や疲労特性が得られる。これらは、伸びが１５％以上、絞りが４０％以上、平滑の疲労強度が５００ＭＰａ以上でかつ切欠（Ｋｔ＝３）の疲労強度が２００ＭＰａというより優れた延性や疲労特性が得られる。さらに、圧延条件Ｂ０１、Ｂ０６、Ｂ０８およびＢ０９のように初析α相の体積分率が５０％以上８０％以下、初析α相の平均結晶粒径が６μｍ以下であるα＋β型チタン合金棒材によって、伸びが２０％以上、絞りが５０％以上、平滑の疲労強度が５５０ＭＰａ以上でかつ切欠（Ｋｔ＝３）の疲労強度が２００ＭＰａというさらに優れた延性や疲労特性が得られる。
一方、化学成分が本発明範囲外のＡ０２を用い、圧延条件Ｂ１０およびＢ１２によって製造した棒材では、圧延条件が本発明範囲内であるため加工発熱は抑制されたが、初析α相の結晶粒径が１０μｍを超えたため、十分な延性や疲労強度が得られない。
実施例２
実施例１で圧延条件Ｂ０１−Ｂ１８により製造した棒材の径方向中心部よりφ８ｍｍ×高さ１２ｍｍの円柱試験片を採取し、８００℃に加熱して７０％圧縮し、圧縮後の表面における割れや肌荒れの発生の有無を調べて、熱間鍛造性の評価を行った。
結果を表２１に示す。
ミクロ組織が本発明範囲内に入る圧延条件Ｂ０１、Ｂ０６、Ｂ０８、Ｂ０９、Ｂ１６、Ｂ１７、Ｂ１８で製造された棒材では、割れや肌荒れが発生せず良好な熱間鍛造性が得られた。
一方、初析α相の結晶粒径が１０μｍを超える圧延条件Ｂ１０、Ｂ１２で製造された棒材では、割れは発生しなかったが肌荒れが発生した。また、中心部や１／４Ｄ部がβ相のみからなる圧延条件Ｂ０２、Ｂ０３、Ｂ０４、Ｂ０５、Ｂ０７、Ｂ１１、Ｂ１４、Ｂ１５で製造された棒材では、割れと肌荒れの両方が発生した。さらに、初析α相の結晶粒径および体積分率は本発明範囲内であるが、圧延方向に平行な断面での結晶粒のアスペクト比が４を超える圧延条件Ｂ１３では、やはり肌荒れが生じた。
【図面の簡単な説明】
図１は、棒材の代表的な製造方法を示す図である。
図２は、棒材圧延の工程を示す図である。
図３は、Ｔｉ−６Ａｌ−４Ｖ合金およびＴｉ−４．５Ａｌ−３Ｖ−２Ｆｅ−２Ｍｏ合金の棒材圧延における被圧延材の温度と圧延時間の関係を示す図である。
図４は、初析α相の平均結晶粒径と高温引張試験による全伸びとの関係を示す図である。
図５は、初析α相の平均結晶粒径と疲労試験における１０８回での疲労強度との関係を示す図である。
図６は、表層部と中心部の温度の経時変化を示す図である。
図７は、被圧延材の断面積と、表層部と中心部の温度差との関係を示す図である。TECHNICAL FIELD The present invention relates to a titanium alloy bar having excellent ductility, fatigue properties and workability, particularly to an α + β type titanium alloy bar and a method for producing the same.
BACKGROUND ART Titanium alloys are used as structural materials in fields such as chemical plants, generators, and aircraft because of their high strength, lightness, and excellent corrosion resistance. Among them, α + β type titanium alloys having high strength and relatively good workability are frequently used.
Products using a titanium alloy include various shapes such as a thin plate, a thick plate and a bar. The rod material may be used as it is, but it may be processed into a complicated shape like a screw part of a bolt or used as a forged material, so the excellent ductility of the rod material itself In addition to fatigue properties and fatigue properties, excellent workability is also required.
FIG. 1 shows a typical method of manufacturing a bar.
The ingot produced by melting is processed into a billet by forging and becomes a rolled material. As shown in FIGS. 2A and 2B, rolling is performed by a reverse type rolling mill or a tandem type rolling mill after heating a billet in a heating furnace, and if necessary, reheating to a temperature at which rolling can be performed by an intermediate heating furnace. .
However, in the case of titanium alloy rods, particularly α + β type titanium alloy rods, the temperature of the material to be rolled increases during hot rolling due to the heat generated during processing, so that stable hot rolling cannot be performed, and ductility, fatigue characteristics, and processing In fact, it is impossible to obtain a titanium alloy bar having excellent properties. For example, when the temperature of the material to be rolled rises to the β transformation point or higher, a β-structure mainly composed of acicular α-phase is obtained, so that excellent ductility and fatigue characteristics cannot be obtained. Further, since the β transformation point is high, even in the case of a Ti-6Al-4V alloy in which the temperature of the material to be rolled hardly becomes equal to or higher than the β transformation point due to the processing heat, the grain growth is promoted when the rolling temperature increases due to the processing heat. Ductility, fatigue properties and workability cannot be obtained.
As a method for solving the problem of the temperature rise due to the heat generated during processing, Japanese Patent Application Laid-Open No. Sho 59-82101 discloses a rolling method in which the cross-sectional reduction rate per rolling pass at the α-region temperature and α + β-region temperature is 40% or less. It has been disclosed. Japanese Patent Application Laid-Open No. 58-25465 discloses a method of cooling a material to be rolled with water in order to suppress a rise in temperature due to heat generated during processing. In addition, the paper 1: "Hot Bar Rolling of Ti-6Al-4V in a Continuous Mill (Titanium '92 Science and Technology)" includes the performance limit of a rolling mill that uses a rolling speed to suppress processing heat itself. It is stated that the speed is reduced.
However, according to the methods disclosed in JP-A-59-82101 and JP-A-58-25465, a titanium alloy rod excellent in ductility, fatigue properties and workability cannot be obtained.
Further, even if the cross-sectional reduction rate is 40% or less per rolling pass according to the method of JP-A-59-82102, suppression of heat generation during processing may be insufficient depending on the type of alloy. Further, the method disclosed in Japanese Patent Application Laid-Open No. 58-25465 has problems that the material is deteriorated due to hydrogen absorption due to water cooling, and that accurate temperature control becomes difficult due to distortion due to rapid cooling.
The method described in the
FIG. 3 shows the relationship between the temperature of the material to be rolled and the rolling time in bar rolling of the Ti-6Al-4V alloy and the Ti-4.5Al-3V-2Fe-2Mo alloy.
Here, the heating temperature is 950 ° C. for the Ti-6Al-4V alloy, and the β-transformation point difference is 100 ° C. lower than that of the Ti-6Al-4V alloy for the Ti-4.5Al-3V-2Fe-2Mo alloy. The temperature was lowered to 850 ° C. Rolling was performed using a reverse type rolling mill and a tandem type rolling mill, and the rolling speed, rolling reduction, and pass schedule were the same for both alloys. The rolling speed in the reverse rolling mill was 2.7 m / sec for both alloys, and the rolling speed in the tandem rolling mill was 2.25 m / sec for both alloys in the final rolling pass at which the rolling speed was the highest. This rolling speed is a lower speed than the rolling speed (6 m / sec) described in the above-mentioned
In the case of rolling of a Ti-6Al-4V alloy, rolling is performed sufficiently in a low temperature range from 1000 ° C., which is the β transformation point of the alloy, and a good structure is obtained. On the other hand, in the case of the Ti-4.5Al-3V-2Fe-2Mo alloy, the deformation resistance increases due to the low-temperature rolling and the processing heat is increased despite the lowering of the heating temperature by the lower β transformation point. As a result, the temperature of the material to be rolled rises to a temperature range exceeding the β transformation point, and a good structure cannot be obtained. Therefore, excellent ductility, fatigue characteristics and workability cannot be obtained. This suggests that it is necessary to consider not only rolling speed but also rolling conditions such as rolling temperature, rolling reduction, and time between rolling passes.
DISCLOSURE OF THE INVENTION An object of the present invention is to provide a high-strength titanium alloy bar excellent in ductility, fatigue properties and workability, and a method for producing the same.
The purpose is substantially in terms of mass% Al: 4% to 5%, V: 2.5% to 3.5%, Fe: 1.5% to 2.5%, Mo: 1. 0.5% or more and 2.5% or less, balance: Ti, volume fraction of proeutectoid α phase is 10% or more and 90% or less, average crystal grain size of proeutectoid α phase is 10 μm or less, parallel to the rolling direction This is achieved by an α + β-type titanium alloy bar having an aspect ratio of crystal grains of the primary α phase in a complicated cross section of 4 or less.
This α + β-type titanium alloy bar can be manufactured by a method of manufacturing an α + β-type titanium alloy bar including a step of hot rolling the titanium alloy of the above component so that the surface temperature thereof is always equal to or lower than the β transformation point.
DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS First, the present inventors studied what kind of structure of an α + β type titanium alloy rod should be formed to obtain excellent ductility, fatigue characteristics and workability. As a result, the following became clear.
The α + β type titanium alloy is composed of a proeutectoid α phase and a transformed β phase. The volume fraction of the α phase having an HCP structure with few slip systems becomes extremely large, and the volume of the transformed β phase having an acicular α phase inside. Since the workability and ductility are reduced even when the fraction is extremely large, the volume fraction of the pro-eutectoid α phase is set to 10% or more and 90% or less. When the volume fractions of the α phase and the β phase are equal to or close to each other during heating in hot working, the workability is further improved. Therefore, the volume fraction of the proeutectoid α phase is 50% or more and 80% or more. The following is desirable.
FIG. 4 shows the relationship between the average crystal grain size of the pro-eutectoid α phase and the total elongation by a high temperature tensile test.
When the average crystal grain size of the pro-eutectoid α phase exceeds 10 μm, the total elongation by a high-temperature tensile test rapidly decreases. In addition, the workability decreases accordingly.
Figure 5 shows the relationship between the fatigue strength at 10 8 times in the fatigue test and the average crystal grain size of the pro-eutectoid α phase.
When the average crystal grain size of the pro-eutectoid α phase exceeds 10 μm, the fatigue strength decreases. Further, when the average crystal grain size of the pro-eutectoid α phase is smaller than 6 μm, higher fatigue strength can be obtained.
In the case of forging a bar, rough surface occurs on a free deformation surface not in contact with a mold due to the shape of crystal grains, that is, the aspect ratio of the crystal grains. Generally, in the case of a rod, crystal grains tend to expand in the rolling direction. In particular, in the case of upset forging, the expanded structure appears on the side surface of the bar, which becomes the free-form surface, so the aspect ratio during forging increases to prevent roughening of the product after forging. Specifically, the aspect ratio of the crystal grains of the pro-eutectoid α phase in a cross section parallel to the rolling direction of the bar needs to be 4 or less.
From the above, the volume fraction of the pro-eutectoid α phase is 10% or more and 90% or less, more preferably 50% or more and 80% or less, and the average crystal grain size of the pro-eutectoid α phase is 10 μm or less, more preferably 6 μm or less. Further, when the aspect ratio of the crystal grains of the pro-eutectoid α phase in the cross section parallel to the rolling direction is set to 4 or less, a high-strength titanium alloy bar excellent in ductility, fatigue properties and workability can be obtained.
The α + β type titanium alloy bar having such a structure is substantially 4% to 5% Al, 2.5% to 3.5% V, and 1.5% Fe by mass%. Not less than 2.5%, Mo: not less than 1.5% and not more than 2.5%, the balance: it is necessary to be a component composed of Ti. The reasons for limiting the content of each element will be described below.
Al: An element indispensable for stabilizing the α phase, and also an element contributing to high strength. If it is less than 4%, high strength is not sufficiently achieved, and if it exceeds 5%, ductility is deteriorated.
V: An element that stabilizes the β phase and also contributes to high strength. If it is less than 2.5%, high strength is not sufficiently achieved, and the β phase is not stable. If it exceeds 3.5%, the processing temperature range becomes narrow due to a decrease in β transformation point, and the cost is increased. .
Mo: an element that stabilizes the β phase and also contributes to high strength. If it is less than 1.5%, the strength cannot be sufficiently increased, and the β phase is not stable. If it exceeds 2.5%, the processing temperature range becomes narrow due to the decrease in β transformation point, and the cost increases. Invite.
Fe: an element that stabilizes the β phase and also contributes to increasing the strength. In addition, although the diffusion speed is high and the processability is improved, if the content is less than 1.5%, the high strength cannot be sufficiently achieved, and the β phase is not stable, so that excellent processability cannot be obtained. If it exceeds 2.5%, the processing temperature region is narrowed due to the decrease in the β transformation point, and the characteristics are deteriorated due to segregation.
The α + β-type titanium alloy rod of the present invention is obtained by adjusting the rolling conditions such as the heating temperature, the rolling temperature range, the rolling reduction, the rolling speed, and the time between passes by adjusting the temperature of the α + β-type titanium alloy of the above-mentioned components by the heat generated during processing. And hot rolling so that the surface temperature is always equal to or lower than the β transformation point. For example, a step of heating an α + β type titanium alloy having a β transformation point of Tβ ° C. so that the surface temperature is in a range of (Tβ−150) ° C. or more and Tβ ° C. or less, and rolling the heated α + β type titanium alloy The surface temperature of the material to be rolled is in the range of (Tβ-300) ° C or more and (Tβ-50) ° C or less, and the finished surface temperature is in the range of (Tβ-300) ° C or more and (Tβ-100) ° C or less. This method has a step of hot rolling.
Here, the reason for heating the surface temperature to a range of (Tβ-150) ° C. or more and Tβ ° C. or less before rolling is that if the surface temperature is lower than (Tβ-150) ° C., the temperature of the material to be rolled at the final stage of rolling becomes large and cracks occur. This causes an increase in sensitivity and deformation resistance, and when the temperature exceeds Tβ ° C., the structure of the material to be rolled becomes a β structure mainly composed of needle-like α and ductility and workability deteriorate. The reason why the surface temperature of the material to be rolled during rolling is in the range of (Tβ-300) ° C or more and (Tβ-50) ° C or less is that if the temperature is less than (Tβ-300) ° C, hot workability decreases and cracking occurs. If the temperature exceeds (Tβ−50) ° C., the temperature rises due to the heat generated during processing, which causes coarsening of crystal grains and formation of a needle-like structure. Further, the reason why the finished surface temperature, which is the surface temperature immediately after the material to be rolled after the final rolling pass is in the range of (Tβ-300) ° C or more and (Tβ-100) ° C or less, is less than (Tβ-300) ° C. If so, cracking sensitivity and deformation resistance increase, and if it exceeds (Tβ-100) ° C., crystal grains become coarse.
Hot rolling is performed in a plurality of rolling passes, but it is preferable to reduce the rolling reduction per rolling pass to 40% or less in order to prevent a temperature rise due to processing heat.
When hot rolling is performed using a reverse type rolling mill, the rolling speed is preferably set to 6 m / sec or less in order to prevent a temperature rise due to processing heat. When using a tandem type rolling mill, the rolling speed is preferably 1.5 m / sec or less.
After each rolling pass, the material to be rolled is cooled from the surface. Therefore, even if the temperature rises due to the heat generated during processing, the temperature of the surface layer of the material to be rolled will decrease to some extent before the next rolling pass. However, as shown in FIG. 6, when the diameter of the material to be rolled is large (in the case of a diameter of 106 mm), the temperature drop at the center of the material to be rolled is small, so that a large A difference appears. If the temperature drop at the center is small, the next rolling pass is performed before the temperature at the center drops, and the temperature further rises due to the heat generated during processing. If this phenomenon continues, the central part will be rolled at a temperature higher than the initial temperature. For this reason, when the diameter of the material to be rolled is large, it is necessary to cool the central portion with sufficient time between rolling passes.
Then, when the present inventors examined in detail the temperature difference between the surface layer portion and the center portion, as shown in FIG. 7, the temperature difference was remarkable when the cross-sectional area in the direction perpendicular to the rolling direction of the material to be rolled was 3500 mm 2 or more. When a material to be rolled having such a large cross-sectional area is rolled so that the cross-sectional area becomes Smm 2 , the time until the start of the next rolling is 0.167 × S 1/2 sec or more. It has been found that it is possible to reduce the temperature difference and to produce a rod having uniform characteristics.
In the production method of the present invention, since the rolling is performed so that the surface temperature of the material to be rolled is always equal to or lower than the β transformation point, the surface temperature of the material to be rolled may be reduced depending on the time between rolling passes or the diameter of the material to be rolled. There is a possibility that the temperature may fall below the appropriate rolling temperature range. In such a case, reheating can be performed by a high-frequency heating facility or the like.
Example 1
Rolled material of 125 mm square was cut out from α + β type titanium alloys A01 (within the scope of the present invention) and A02 (within the scope of the present invention) having the chemical components shown in Table 1 and rolling conditions B01-B18 shown in Table 2 using a caliber rolling mill. Produced a rod having a diameter of 20 mm or 50 mm. The time between rolling passes in Table 2 is indicated by ○ when the time between rolling passes is 0.167 × S 1/2 sec or more in all rolling passes under each rolling condition, and × when not. Is shown. Tables 3 to 20 show the cross-sectional area S, rolling reduction, 0.167 × S 1/2 , time between rolling passes, surface temperature, and rolling speed of the material to be rolled in each rolling pass under each rolling condition. Was. In the rolling equipment column in the table, R represents a reverse rolling mill, and T represents a tandem rolling mill.
Tensile tests were performed on the manufactured rods after annealing at 700 ° C. or more and 720 ° C. or less, and the yield strength (0.2% PS), tensile strength (UTS), extension (El), and drawing (RA) were measured. Further, a smoothing test (condition: Kt = 1) and a notch test (condition: Kt = 3) were performed to measure the fatigue strength. Further, in order to examine the microstructure, the microstructure at the center of the bar and at a quarter position (直径 D) of the diameter was observed, and the crystal grain size of the proeutectoid α phase, its volume fraction, and the rolling direction The aspect ratio of the crystal grain in the cross section parallel to was measured.
The results are shown in Table 21. The absence of the crystal grain size in the microstructures in the table is because the site consisted only of a β structure mainly composed of acicular α, and no equiaxed primary α phase was observed.
When the heating temperature of the surface was lower than (Tβ-150) ° C., the surface temperature of the material to be rolled was too low, and the rolling load was too large to perform rolling. When the heating temperature is higher than Tβ ° C., the surface temperature of the material to be rolled is too high even if the time between rolling passes is within the range of the present invention as in the rolling conditions B02 and B11. The surface temperature of the material exceeds Tβ, and the structure at the center becomes a β structure mainly composed of acicular α, and the ductility and fatigue characteristics are deteriorated.
If the finished surface temperature is lower than (Tβ-300) ° C., the temperature of the material to be rolled is too low, the workability is reduced, and cracks occur during rolling. When the finish surface temperature is higher than (Tβ-100) ° C., the microstructure cannot be refined as in the rolling conditions B04, B05, and B07, and the ductility and fatigue characteristics deteriorate.
When the surface temperature of the material to be rolled is lower than (Tβ-300) ° C., the surface temperature of the material to be rolled is too low, and cracks occur during rolling. Further, when the surface temperature of the material to be rolled is higher than (Tβ-50) ° C., as in the rolling conditions B02-B05, B07 and B11, the structure of the central part and the 1 / 4D part after rolling shows needle-like α. Becomes the main β-structure, and the ductility and fatigue properties deteriorate.
When the rolling reduction in one rolling pass exceeded 40%, the heat generated during processing increased, and the temperature of the material to be rolled exceeded Tβ, making it impossible to refine the structure.
When a reverse type rolling mill is used as in rolling condition B14 and the rolling speed exceeds 6 m / sec, or when a tandem type rolling mill is used as in rolling condition B15 and the rolling speed exceeds 1.5 m / sec, processing is performed. Heat generation increased, the surface temperature of the material to be rolled exceeded Tβ, and the structure could not be refined.
When the time between rolling passes is out of the range of the present invention, the surface temperature of the material to be rolled due to the heat generated by processing exceeds the temperature decrease due to cooling, the surface temperature of the material to be rolled exceeds Tβ, and the structure cannot be refined. Was.
In the bar manufactured using A01 whose chemical component is within the range of the present invention and under rolling conditions B01, B06, B08, B09, B16, B17 and B18, a uniform microstructure in which the crystal grain size of the proeutectoid α phase is 10 μm or less. Are observed, and excellent ductility and fatigue properties are obtained. These have more excellent ductility and fatigue properties of elongation of 15% or more, reduction of 40% or more, smooth fatigue strength of 500 MPa or more, and notch (Kt = 3) fatigue strength of 200 MPa. Further, α + β-type titanium alloy bars having a volume fraction of pro-eutectoid α phase of 50% or more and 80% or less and an average crystal grain size of pro-eutectoid α phase of 6 μm or less as in rolling conditions B01, B06, B08 and B09. Thereby, more excellent ductility and fatigue properties such that the elongation is 20% or more, the drawing is 50% or more, the smooth fatigue strength is 550 MPa or more, and the notch (Kt = 3) fatigue strength is 200 MPa are obtained.
On the other hand, in the bar manufactured using A02 whose chemical component is outside the range of the present invention and manufactured under the rolling conditions B10 and B12, since the rolling condition was within the range of the present invention, the heat generation during processing was suppressed. Since the particle size exceeds 10 μm, sufficient ductility and fatigue strength cannot be obtained.
Example 2
A cylindrical test piece having a diameter of 8 mm and a height of 12 mm was collected from the center in the radial direction of the bar manufactured under the rolling conditions B01-B18 in Example 1, heated to 800 ° C., compressed by 70%, and cracked on the compressed surface. The hot forging property was evaluated by examining the presence or absence of the occurrence of roughening and skin roughness.
The results are shown in Table 21.
Bars manufactured under rolling conditions B01, B06, B08, B09, B16, B17, and B18 whose microstructure falls within the range of the present invention did not cause cracking or roughening, and obtained good hot forgeability.
On the other hand, in the rods manufactured under the rolling conditions B10 and B12 in which the crystal grain diameter of the pro-eutectoid α phase exceeds 10 μm, cracks did not occur, but surface roughness occurred. Further, in the bars manufactured under the rolling conditions B02, B03, B04, B05, B07, B11, B14, and B15 in which the central portion and the 1 / 4D portion consist only of the β phase, both cracking and roughening occurred. Further, the crystal grain size and the volume fraction of the pro-eutectoid α phase are within the range of the present invention, but under the rolling condition B13 in which the aspect ratio of the crystal grains in the cross section parallel to the rolling direction exceeds 4, the surface roughness still occurs. .
[Brief description of the drawings]
FIG. 1 is a diagram showing a typical method of manufacturing a bar.
FIG. 2 is a diagram showing a bar rolling process.
FIG. 3 is a diagram showing the relationship between the temperature of the material to be rolled and the rolling time in bar rolling of the Ti-6Al-4V alloy and the Ti-4.5Al-3V-2Fe-2Mo alloy.
FIG. 4 is a diagram showing the relationship between the average crystal grain size of the pro-eutectoid α phase and the total elongation by a high-temperature tensile test.
Figure 5 is a diagram showing a relationship between fatigue strength at 10 8 times in the fatigue test and the average crystal grain size of the pro-eutectoid α phase.
FIG. 6 is a diagram showing changes over time in the temperature of the surface layer portion and the central portion.
FIG. 7 is a diagram showing the relationship between the cross-sectional area of the material to be rolled and the temperature difference between the surface portion and the central portion.
Claims (9)
前記加熱されたα＋β型チタン合金を、圧延中の被圧延材の表面温度が（Ｔβ−３００）℃以上（Ｔβ−５０）℃以下の範囲に、最終圧延パスを終えた直後の表面温度である仕上表面温度が（Ｔβ−３００）℃以上（Ｔβ−１００）℃以下の範囲になるように熱間圧延する工程と、
を有する請求の範囲３のα＋β型チタン合金棒材の製造方法。heating an α + β type titanium alloy having a β transformation point of Tβ ° C. so that the surface temperature falls within a range of (Tβ−150) ° C. or more and Tβ ° C. or less;
The surface temperature of the heated α + β-type titanium alloy is a surface temperature immediately after finishing the final rolling pass so that the surface temperature of the material to be rolled during rolling is in a range of (Tβ-300) ° C or more and (Tβ-50) ° C or less. Hot rolling so that the finish surface temperature is in the range of (Tβ-300) ° C. or more and (Tβ-100) ° C. or less;
4. The method for producing an α + β type titanium alloy bar according to claim 3, comprising:
Applications Claiming Priority (3)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2001054809 | 2001-02-28 | ||
| JP2001054809 | 2001-02-28 | ||
| PCT/JP2002/001710 WO2002070763A1 (en) | 2001-02-28 | 2002-02-26 | Titanium alloy bar and method for production thereof |
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| JPWO2002070763A1 true JPWO2002070763A1 (en) | 2004-07-02 |
| JP4013761B2 JP4013761B2 (en) | 2007-11-28 |
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| US (2) | US20030223902A1 (en) |
| EP (1) | EP1382695A4 (en) |
| JP (1) | JP4013761B2 (en) |
| RU (1) | RU2259413C2 (en) |
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| US20040221929A1 (en) | 2003-05-09 | 2004-11-11 | Hebda John J. | Processing of titanium-aluminum-vanadium alloys and products made thereby |
| JP4264411B2 (en) * | 2004-04-09 | 2009-05-20 | 新日本製鐵株式会社 | High strength α + β type titanium alloy |
| US7837812B2 (en) | 2004-05-21 | 2010-11-23 | Ati Properties, Inc. | Metastable beta-titanium alloys and methods of processing the same by direct aging |
| RU2269584C1 (en) * | 2004-07-30 | 2006-02-10 | Открытое Акционерное Общество "Корпорация Всмпо-Ависма" | Titanium-base alloy |
| JP4655666B2 (en) * | 2005-02-23 | 2011-03-23 | Ｊｆｅスチール株式会社 | Golf club head |
| RU2383654C1 (en) * | 2008-10-22 | 2010-03-10 | Государственное образовательное учреждение высшего профессионального образования "Уфимский государственный авиационный технический университет" | Nano-structural technically pure titanium for bio-medicine and method of producing wire out of it |
| 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 |
| 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 |
| JP5196083B2 (en) | 2011-02-24 | 2013-05-15 | 新日鐵住金株式会社 | High-strength α + β-type titanium alloy hot-rolled sheet excellent in cold coil handling and manufacturing method thereof |
| WO2012115242A1 (en) | 2011-02-24 | 2012-08-30 | 新日本製鐵株式会社 | α+β TYPE TITANIUM ALLOY SHEET WITH EXCELLENT COLD ROLLING PROPERTIES AND COLD HANDLING PROPERTIES, AND PRODUCTION METHOD THEREFOR |
| US8652400B2 (en) | 2011-06-01 | 2014-02-18 | Ati Properties, Inc. | Thermo-mechanical processing of nickel-base alloys |
| ES2620310T3 (en) | 2011-06-17 | 2017-06-28 | Titanium Metals Corporation | Method for manufacturing alpha-beta alloy plates from Ti-Al-V-Mo-Fe |
| CN102586639A (en) * | 2012-03-16 | 2012-07-18 | 广州有色金属研究院 | Method for preparing titanium alloy through high-speed pressing formation |
| 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 |
| JP6230885B2 (en) * | 2013-11-22 | 2017-11-15 | 東邦チタニウム株式会社 | α + β type titanium alloy and method for producing the same |
| CN104532057B (en) * | 2014-12-11 | 2017-01-04 | 西部超导材料科技股份有限公司 | A kind of Ti6242 titanium alloy and the preparation method of small-sized bar thereof |
| US10094003B2 (en) | 2015-01-12 | 2018-10-09 | Ati Properties Llc | Titanium alloy |
| CN105251804B (en) * | 2015-10-28 | 2018-05-08 | 西部超导材料科技股份有限公司 | A kind of milling method of six square rod of TC6 titanium alloys |
| US10502252B2 (en) | 2015-11-23 | 2019-12-10 | Ati Properties Llc | Processing of alpha-beta titanium alloys |
| CA3009962C (en) * | 2015-12-22 | 2021-11-09 | Stock Company "Chepetsky Mechanical Plant" (Sc Cmp) | A method of manufacturing rods from titanium alloys |
| CN107138523B (en) * | 2017-06-29 | 2019-07-02 | 西部超导材料科技股份有限公司 | A kind of TB9 titanium alloy wire bar and its milling method |
| CN109283205B (en) * | 2018-10-19 | 2021-03-26 | 中国航发北京航空材料研究院 | Method for measuring volume fraction of primary alpha phase in titanium alloy structure |
| JP7024861B2 (en) * | 2018-11-15 | 2022-02-24 | 日本製鉄株式会社 | Titanium alloy wire rod and titanium alloy wire rod manufacturing method |
| CN111545574A (en) * | 2020-05-20 | 2020-08-18 | 攀钢集团攀枝花钛材有限公司江油分公司 | TA15 hot rolling plate structure control method |
| CN114535343B (en) * | 2022-04-26 | 2022-08-30 | 西部宝德科技股份有限公司 | Titanium fiber preparation method |
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| JPS5825465A (en) * | 1981-08-05 | 1983-02-15 | Sumitomo Metal Ind Ltd | Manufacture of rolled titanium alloy material having excellent structure |
| JPS5982101A (en) * | 1982-11-01 | 1984-05-12 | Sumitomo Metal Ind Ltd | Production of titanium alloy bar |
| US5362441A (en) * | 1989-07-10 | 1994-11-08 | Nkk Corporation | Ti-Al-V-Mo-O alloys with an iron group element |
| DE69024418T2 (en) * | 1989-07-10 | 1996-05-15 | Nippon Kokan Kk | Titanium-based alloy and process for its superplastic shaping |
| US5346217A (en) * | 1991-02-08 | 1994-09-13 | Yamaha Corporation | Hollow metal alloy wood-type golf head |
| JP2884913B2 (en) * | 1992-04-21 | 1999-04-19 | 日本鋼管株式会社 | Manufacturing method of α + β type titanium alloy sheet for superplastic working |
| JP3083225B2 (en) * | 1993-12-01 | 2000-09-04 | オリエント時計株式会社 | Manufacturing method of titanium alloy decorative article and watch exterior part |
| JP3114503B2 (en) * | 1994-07-14 | 2000-12-04 | 日本鋼管株式会社 | Method for producing (α + β) type titanium alloy having locally excellent wear resistance |
| JPH08103831A (en) * | 1994-10-05 | 1996-04-23 | Nkk Corp | Punching method for titanium alloy plate material |
| JP3319195B2 (en) * | 1994-12-05 | 2002-08-26 | 日本鋼管株式会社 | Toughening method of α + β type titanium alloy |
| JPH10306335A (en) * | 1997-04-30 | 1998-11-17 | Nkk Corp | Alpha plus beta titanium alloy bar and wire rod, and its production |
| JP4655666B2 (en) * | 2005-02-23 | 2011-03-23 | Ｊｆｅスチール株式会社 | Golf club head |
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| JP4013761B2 (en) | 2007-11-28 |
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| RU2259413C2 (en) | 2005-08-27 |
| WO2002070763A1 (en) | 2002-09-12 |
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