JP2020152971A - Titanium alloy bar and its manufacturing method - Google Patents
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- 229910001069 Ti alloy Inorganic materials 0.000 title claims abstract description 143
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 14
- 239000013078 crystal Substances 0.000 claims abstract description 118
- 238000005242 forging Methods 0.000 claims description 77
- 238000010438 heat treatment Methods 0.000 claims description 55
- 229910052751 metal Inorganic materials 0.000 claims description 33
- 239000002184 metal Substances 0.000 claims description 33
- 238000000034 method Methods 0.000 claims description 22
- 238000001816 cooling Methods 0.000 claims description 19
- 238000012545 processing Methods 0.000 claims description 18
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- 239000012535 impurity Substances 0.000 claims description 8
- 239000010936 titanium Substances 0.000 claims description 8
- 238000010791 quenching Methods 0.000 claims description 6
- 239000000126 substance Substances 0.000 claims description 6
- 230000000171 quenching effect Effects 0.000 claims description 5
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- 239000000203 mixture Substances 0.000 claims description 3
- 238000005096 rolling process Methods 0.000 claims description 3
- 230000009466 transformation Effects 0.000 description 28
- 230000009467 reduction Effects 0.000 description 21
- 238000005259 measurement Methods 0.000 description 15
- 230000010354 integration Effects 0.000 description 13
- 239000000463 material Substances 0.000 description 13
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- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 8
- 238000001887 electron backscatter diffraction Methods 0.000 description 7
- PCHJSUWPFVWCPO-UHFFFAOYSA-N gold Chemical compound [Au] PCHJSUWPFVWCPO-UHFFFAOYSA-N 0.000 description 6
- 239000010931 gold Substances 0.000 description 6
- 229910052737 gold Inorganic materials 0.000 description 6
- 238000012360 testing method Methods 0.000 description 6
- 238000009825 accumulation Methods 0.000 description 5
- 229910045601 alloy Inorganic materials 0.000 description 5
- 239000000956 alloy Substances 0.000 description 5
- 238000010586 diagram Methods 0.000 description 4
- 238000009661 fatigue test Methods 0.000 description 4
- 230000007246 mechanism Effects 0.000 description 3
- RTAQQCXQSZGOHL-UHFFFAOYSA-N Titanium Chemical compound [Ti] RTAQQCXQSZGOHL-UHFFFAOYSA-N 0.000 description 2
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- 229910052719 titanium Inorganic materials 0.000 description 2
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- 229910000883 Ti6Al4V Inorganic materials 0.000 description 1
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Abstract
Description
本発明は、チタン合金棒材及びその製造方法に関する。 The present invention relates to a titanium alloy rod and a method for producing the same.
チタン合金は軽量高強度の材料として、航空機、自動車、ゴルフクラブ等の民生品などの分野で使用されている。チタン合金の中で汎用的に使われる合金は、主としてα相とβ相から構成され、Ｔｉ−６Ａｌ−４Ｖ、Ｔｉ−６Ａｌ−２Ｓｎ−４Ｚｒ−２Ｍｏ、Ｔｉ−５Ａｌ−１Ｆｅ合金などが知られている。 Titanium alloy is used as a lightweight and high-strength material in the fields of consumer products such as aircraft, automobiles, and golf clubs. Among titanium alloys, alloys generally used are mainly composed of α phase and β phase, and Ti-6Al-4V, Ti-6Al-2Sn-4Zr-2Mo, Ti-5Al-1Fe alloy and the like are known. There is.
稠密六方晶構造からなるチタンのα相は、高い応力が加わると室温などの低温においてもクリープ変形しやすく、α相を含むチタン合金においても室温でクリープ変形を生じることが知られている。さらに、α相を含むチタン合金におけるクリープ変形しやすい特性は、台形波型の負荷サイクルに代表される高負荷状態が一定時間継続する疲労（Ｄｗｅｌｌ疲労）において、寿命低下を招くことが知られている。（非特許文献１〜３）
It is known that the α phase of titanium having a dense hexagonal structure is easily creep-deformed even at a low temperature such as room temperature when a high stress is applied, and that a titanium alloy containing an α phase also undergoes creep deformation at room temperature. Further, it is known that the creep deformation property of a titanium alloy containing an α phase causes a decrease in life in fatigue (Dwell fatigue) in which a high load state such as a trapezoidal wave type load cycle continues for a certain period of time. There is. (Non-Patent
Ｄｗｅｌｌ疲労では、高負荷状態が継続することがない三角波あるいは正弦波の負荷サイクルの場合と比較して、少ないサイクル数で破断に至るため、特に、航空機のジェットエンジン部品として使用される場合に問題になることがある。 Dwell fatigue causes rupture in a smaller number of cycles compared to triangular or sinusoidal load cycles where high load conditions do not continue, which is a problem, especially when used as an aircraft jet engine component. May become.
特許文献１（特開２０１６−１９９７９６号公報）では、優れた疲労特性を有するチタン合金棒材およびその製造方法が開示されている。特許文献１では、初析α粒のうち、稠密六方構造のｃ軸方向とチタン合金棒材の長さ方向とのなす角度（ｃ軸の傾き）が２５°以上５５°以下で、かつ円相当直径が２０μｍ以上である初析α粒の金属組織中の面積率が２．０％以下であることが述べられている。これは特許文献１の段落００２０に記載の、「稠密六方晶の底面すべりは、結晶方位（図２においては符号「θ」で示す。）が４５°に近いほど生じやすく、結晶方位が２５°以上５５°以下であると活発になる。また、金属組織に含まれる等軸状の初析α粒の大きさが大きいほど、試験片に付与される応力が集中しやすく、円相当直径が２０μｍ以上であると応力の集中が顕著となる。したがって、ｃ軸の傾きが２５°以上５５°以下で、かつ円相当直径が２０μｍ以上の初析α粒は、稠密六方晶の底面すべりが生じやすく、しかも応力が集中しやすいため、疲労寿命が短くなったと考えられる。」との技術思想に基づくものであり、通常の疲労破壊の機構として妥当なものである。
Patent Document 1 (Japanese Unexamined Patent Publication No. 2016-199796) discloses a titanium alloy rod having excellent fatigue characteristics and a method for producing the same. In
一方、非特許文献１〜３に説明されているように、Ｄｗｅｌｌ疲労では、異なる破壊機構が知られている。これらの文献によると、ｃ軸の傾きが４５°付近のα粒（Ｓ）と、ｃ軸が応力方向に対し垂直に近い方位のα粒（Ｈ）が隣接する場合、Ｈ粒に応力が集中して応力軸に垂直なファセット状破面が生じるとされる。また、このファセットは稠密六方晶の底面とほぼ平行であることが、別の研究により知られている。
On the other hand, as described in
特許文献１には、Ｄｗｅｌｌ疲労について何の言及もされていない。
特許文献２（特表２００９−５３１５４６号公報）には、Ｄｗｅｌｌ疲労に対する抵抗力を改善する技術が開示されている。ここでは、ＴＡ６Ｚｒ４ＤＥ（Ｔｉ−６Ａｌ−２Ｓｎ−４Ｚｒ−２Ｍｏ）合金において、β変態点−２０〜−１５℃の温度で４〜８時間の熱処理を施すことで、破断寿命が５５００回から１００００回に向上した。しかし、熱処理以前の工程はβ域におけるスタンピングのみであり、それ以前の加工熱処理工程は不明確であり、充分に微細なミクロ組織を形成することができず、通常の疲労寿命に対するＤｗｅｌｌ疲労寿命の低下代を縮小する効果は不確実である。 Patent Document 2 (Japanese Patent Laid-Open No. 2009-531546) discloses a technique for improving resistance to Dwell fatigue. Here, the TA6Zr4DE (Ti-6Al-2Sn-4Zr-2Mo) alloy is heat-treated at a β transformation point-20 to -15 ° C. for 4 to 8 hours to increase the fracture life from 5500 to 10000. Improved. However, the process before the heat treatment is only stamping in the β region, and the processing heat treatment process before that is unclear, and it is not possible to form a sufficiently fine microstructure, and the Dwell fatigue life is different from the normal fatigue life. The effect of reducing the reduction allowance is uncertain.
特許文献３（特開２０１２−２２４９３５号公報）には、α相のｃ軸の特定方向に対する集積度が規定されたチタン合金ビレットが開示されている。しかし、疲労破壊の起点となるα相の粒径については言及されておらず、単に集積度を高めただけで疲労特性が改善されるものではない。 Patent Document 3 (Japanese Unexamined Patent Publication No. 2012-224935) discloses a titanium alloy billet in which the degree of integration of the α phase in a specific direction is defined. However, the particle size of the α phase, which is the starting point of fatigue fracture, is not mentioned, and the fatigue characteristics are not improved simply by increasing the degree of integration.
特許文献４（特開２０１４−６５９６７号公報）には、α相のｃ軸の特定方向に対する集積度が規定されたチタン合金ビレットが開示されている。しかし、同特許文献は、疲労強度の向上を意図したものではなく、また、ｃ軸の集積方法は、本発明の方向とは異なっている。 Patent Document 4 (Japanese Unexamined Patent Publication No. 2014-60967) discloses a titanium alloy billet in which the degree of integration of the α phase with respect to the c-axis is defined. However, the patent document is not intended to improve fatigue strength, and the c-axis integration method is different from the direction of the present invention.
本発明は上記事情に鑑みてなされたものであり、Ｄｗｅｌｌ疲労特性の良好なチタン合金棒材及びその製造方法を提供することを課題とする。 The present invention has been made in view of the above circumstances, and an object of the present invention is to provide a titanium alloy rod having good Dwell fatigue characteristics and a method for producing the same.
上記課題を解決する手段は下記の通りである。なお、本発明において良好なＤｗｅｌｌ疲労特性とは、通常の正弦波あるいは三角波の疲労寿命に対するＤｗｅｌｌ疲労寿命の低下代が小さいことを意味する。 The means for solving the above problems are as follows. The good Dwell fatigue characteristic in the present invention means that the reduction margin of the Dwell fatigue life is small with respect to the fatigue life of a normal sine wave or a triangular wave.
［１］ ２５℃においていてα相を主相としβ相を第２相とする金属組織を有するチタン合金棒材であって、
α結晶粒を構成する稠密六方結晶の（０００１）面の法線方向と、前記チタン合金棒材の長軸方向とのなす角度θ１が０°以上２５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が５．０％以下であるとともに、
前記（０００１）面の法線方向と、前記長軸方向とのなす角度θ１が２５°以上５５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が２．０％以下であり、
かつ、α結晶粒を構成する稠密六方結晶の（１０−１０）面の法線方向のうちのひとつの方向と、前記長軸方向とのなす角度θ２が０°以上３０°以下の範囲にあるα結晶粒の面積率が４０％以上であることを特徴とする、チタン合金棒材。
［２］ 化学成分が、Ａｌ：５．５０〜６．７５質量％、Ｖ：３．５〜４．５質量％、Ｆｅ：０．０５〜０．４０質量％、Ｏ：０．０５〜０．２５質量％を含有し、残部がＴｉおよび不純物からなる［１］に記載のチタン合金棒材。
［３］ 化学成分が、Ａｌ：５．５０〜６．５０質量％、Ｓｎ：１．７５〜２．２５質量％、Ｚｒ：３．５〜４．５質量％、Ｍｏ：１．８〜２．２質量％、Ｆｅ：０．０２〜０．２５質量％、Ｏ：０．０２〜０．１５質量％を含有し、残部がＴｉおよび不純物からなる［１］に記載のチタン合金棒材。
［４］ 鋳塊を熱間加工して得られた、２５℃においてα相を主相としβ相を第２相とする金属組織を有するチタン合金ビレットをβ単相域の温度に加熱した後に急冷する第１の工程と、
前記チタン合金ビレットをα＋β二相域の温度に加熱し、前記チタン合金ビレットを鍛造した後に冷却する第２の工程と、
前記チタン合金ビレットを、α＋β二相域の温度であって前記第２の工程の加熱温度以下の温度に加熱し、前記チタン合金ビレットを鍛造する処理を１回以上行い、少なくとも最後に３００℃以下まで冷却する処理を行う第３の工程と、
をこの順で行う際に、
前記第２の工程における前記鍛造は、前記チタン合金ビレットを送り量Ｌiniで長軸方向に送りつつ金敷で圧下する加工であって、鍛造前の前記チタン合金ビレットの幅をＷiniとしたときにＬini／Ｗiniが０．８０以下を満たし、鍛造後の前記チタン合金ビレットの高さＨafterと幅Ｗafterとの比Ｈafter／Ｗafterが０．６７以上１．５以下となるように、かつ、前記Ｗiniと前記Ｗafterとの比ΔＷ（ΔＷ＝Ｗafter／Ｗini）が１．０５以上１．１５以下になるように圧下する鍛造であり、この鍛造を少なくとも２回以上行い、また、前記チタン合金ビレットを長軸周りに回転させて前記チタン合金ビレットに対する圧下方向を各回毎に変更させることとし、
前記第２の工程における鍛錬比を１．５以上とし、前記第３の工程の鍛錬比を３．０以上とする、
ことを特徴とする［１］〜［３］のいずれか一項に記載のチタン合金棒材の製造方法。
［５］ 前記第１の工程が、前記チタン合金ビレットをβ単相域の温度に加熱した後に、加工してから急冷する工程である、［４］に記載のチタン合金棒材の製造方法。
[1] A titanium alloy rod having a metal structure having an α phase as a main phase and a β phase as a second phase at 25 ° C.
The circle-equivalent diameter in which the angle θ 1 between the normal direction of the (0001) plane of the dense hexagonal crystal constituting the α crystal grain and the long axis direction of the titanium alloy bar is in the range of 0 ° or more and 25 ° or less The area ratio of α crystal grains over 20 μm is 5.0% or less, and
The area ratio of α crystal grains having a circle-equivalent diameter of more than 20 μm in the range where the angle θ 1 formed by the normal direction of the (0001) plane and the semimajor axis is 25 ° or more and 55 ° or less is 2.0%. Is below
In addition, the angle θ 2 formed by one of the normal directions of the (10-10) planes of the dense hexagonal crystals constituting the α crystal grain and the major axis direction is in the range of 0 ° or more and 30 ° or less. A titanium alloy rod having an area ratio of a certain α crystal grain of 40% or more.
[2] The chemical components are Al: 5.50 to 6.75% by mass, V: 3.5 to 4.5% by mass, Fe: 0.05 to 0.40% by mass, O: 0.05 to 0. The titanium alloy rod according to [1], which contains .25% by mass and the balance is Ti and impurities.
[3] The chemical components are Al: 5.50 to 6.50% by mass, Sn: 1.75 to 2.25% by mass, Zr: 3.5 to 4.5% by mass, Mo: 1.8 to 2 The titanium alloy rod according to [1], which contains 2% by mass, Fe: 0.02 to 0.25% by mass, O: 0.02 to 0.15% by mass, and the balance is Ti and impurities.
[4] After heating a titanium alloy billet having a metal structure having an α phase as a main phase and a β phase as a second phase at 25 ° C., which is obtained by hot working an ingot, to a temperature in the β single phase region. The first step of quenching and
A second step of heating the titanium alloy billet to a temperature in the α + β two-phase region, forging the titanium alloy billet, and then cooling the billet.
The titanium alloy billet is heated to a temperature in the α + β two-phase region and equal to or lower than the heating temperature of the second step, and the titanium alloy billet is forged at least once, and at least finally at 300 ° C. or lower. The third step of cooling to
When doing in this order
The forging in the second step is a process in which the titanium alloy billet is fed in the major axis direction with a feed amount of Lini and pressed down with a metal pad, and when the width of the titanium alloy billet before forging is Wini, the Lini. / Wini satisfies 0.80 or less, the ratio Hafter / Wafter of the height Hafter and the width Wafter of the forged titanium alloy billet is 0.67 or more and 1.5 or less, and the Wini and the said Forging is forging so that the ratio ΔW (ΔW = Wafter / Wini) with Wafter is 1.05 or more and 1.15 or less. This forging is performed at least twice, and the titanium alloy billet is placed around the long axis. It was decided to change the rolling direction with respect to the titanium alloy billet each time.
The forging ratio in the second step is 1.5 or more, and the forging ratio in the third step is 3.0 or more.
The method for producing a titanium alloy bar according to any one of [1] to [3].
[5] The method for producing a titanium alloy rod according to [4], wherein the first step is a step of heating the titanium alloy billet to a temperature in the β single-phase region, processing it, and then quenching it.
本発明によれば、Ｄｗｅｌｌ疲労特性の良好なチタン合金棒材及びその製造方法を提供できる。 According to the present invention, it is possible to provide a titanium alloy rod having good Dwell fatigue characteristics and a method for producing the same.
チタン合金の引張特性には、集合組織によって異方性があることが知られている。応力方向に（０００１）面に垂直な方位が集積した場合は０．２％耐力や引張強度が高くなるが、応力方向に（１０−１０）面に垂直な方位が集積した場合は０．２％耐力や引張強度が低くなる。通常の三角波あるいは正弦波による疲労特性も同様である。例えば、疲労寿命を横軸に、最大応力（σＭＡＸ）を０．２％耐力（σ０．２）で規格化した”σＭＡＸ／σ０．２”を縦軸にとってグラフ化（規格化されたＳ−Ｎ線図）した場合、集合組織によらずほぼ同一の線上に表される。
なお、本明細書において、「（１０−１０）面」と表記する場合の「−１」は、「１」の上に線を引いたことを意味する。
It is known that the tensile properties of titanium alloys are anisotropic depending on the texture. When the orientations perpendicular to the (0001) plane in the stress direction are accumulated, the yield strength and tensile strength are increased by 0.2%, but when the orientations perpendicular to the (10-10) plane in the stress direction are accumulated, 0.2%. % The yield strength and tensile strength are low. The same applies to the fatigue characteristics of a normal triangular wave or sine wave. For example, the horizontal axis is fatigue life, and the vertical axis is "σMAX / σ0.2", which is the standardized maximum stress (σMAX) with 0.2% proof stress (σ0.2). In the case of a diagram), it is represented on almost the same line regardless of the texture.
In this specification, "-1" in the case of "(10-10) plane" means that a line is drawn on "1".
しかし、Ｄｗｅｌｌ疲労特性は、規格化されたＳ−Ｎ線図で表される挙動が異なっていることがわかった。すなわち、稠密六方晶の底面（以下、（０００１）面という場合がある）の法線方向が応力軸に平行に集積した集合組織の場合、異なる方位に集積した集合組織と比較して寿命が大幅に低下する。 However, it was found that the Dwell fatigue characteristics differed in the behavior represented by the standardized SN diagram. That is, in the case of an aggregate in which the normal direction of the bottom surface of the dense hexagonal crystal (hereinafter, may be referred to as (0001) plane) is parallel to the stress axis, the lifetime is significantly longer than that of the aggregate in different directions. Decreases to.
航空機エンジン部品の素材として使用されるチタン合金棒材の長軸方向においては、通常の疲労寿命に対するＤｗｅｌｌ疲労寿命の低下代が小さいことが好ましい。 In the long axis direction of the titanium alloy rod used as a material for aircraft engine parts, it is preferable that the reduction allowance of the Dwell fatigue life with respect to the normal fatigue life is small.
Ｄｗｅｌｌ疲労における通常疲労に対する寿命低下は以下の機構によるものと考えられる。Ｄｗｅｌｌ疲労では、ひずみ蓄積によりき裂発生が促進され、また、稠密六方晶の底面（（０００１）面）にほぼ平行なファセット破面の形成によりき裂進展が促進されることから、寿命低下に至る。Ｄｗｅｌｌ疲労寿命は、応力軸方向に対する特定の結晶方位を有する粗大なα相の面積率が大きいほどき裂発生が促進され、低下する。 It is considered that the decrease in life with respect to normal fatigue in Dwell fatigue is due to the following mechanism. In Dwell fatigue, strain accumulation promotes crack formation, and the formation of faceted fracture surfaces almost parallel to the bottom surface ((0001) plane) of dense hexagonal crystals promotes crack growth, resulting in a decrease in life. To reach. The Dwell fatigue life decreases as the area ratio of the coarse α phase having a specific crystal orientation with respect to the stress axis direction increases, the crack generation is promoted.
そこで、特定の結晶方位を有する粗大なα粒が少なく、かつ、稠密六方晶の底面が負荷方向に対して垂直になる比率（面積率）が少ないことが、Ｄｗｅｌｌ疲労寿命向上に有利である。 Therefore, it is advantageous for improving the Dwell fatigue life that the number of coarse α grains having a specific crystal orientation is small and the ratio (area ratio) at which the bottom surface of the dense hexagonal crystal is perpendicular to the load direction is small.
また、通常の疲労においてき裂発生の起点となりやすい特定方位を有する粗大なα結晶粒は、Ｄｗｅｌｌ疲労においてもき裂発生の起点になりやすい。そのため、α結晶粒の（０００１）面の法線方向と、応力軸方向とのなす角度が２５°以上５５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が小さいことが好ましい。 Further, coarse α crystal grains having a specific orientation that are likely to be the starting point of crack generation in normal fatigue are likely to be the starting point of crack generation in Dwell fatigue. Therefore, the area ratio of the α crystal grains having a circle-equivalent diameter of more than 20 μm in the range of 25 ° or more and 55 ° or less between the normal direction of the (0001) plane of the α crystal grains and the stress axis direction is small. Is preferable.
上記のようにα結晶粒の大きさや結晶方位を制御するには、チタン合金の熱間加工中の金属組織変化挙動を把握することが重要である。一般に、チタン合金の鍛造工程において、β単相域に加熱することで、それ以前に存在するα相の結晶方位の偏りを軽減してランダム化する工程が組み込まれる。しかし、その後にα＋β域で加工することにより、新たにα相の集合組織が形成される。特に、β単相域から冷却した後の最初のα＋β域での加工によって形成されるα相の集合組織を、その後のα＋β域での加工によって消滅させることは困難である。そのため、β単相域から冷却した後の最初のα＋β域での加工方法を制御することが必要である。 In order to control the size and crystal orientation of α crystal grains as described above, it is important to understand the metal structure change behavior during hot working of the titanium alloy. Generally, in the forging process of a titanium alloy, a step of reducing the bias of the crystal orientation of the α phase existing before that and randomizing it is incorporated by heating to the β single phase region. However, by subsequent processing in the α + β region, a new α-phase texture is formed. In particular, it is difficult to eliminate the α-phase texture formed by the first processing in the α + β region after cooling from the β single-phase region by the subsequent processing in the α + β region. Therefore, it is necessary to control the processing method in the first α + β region after cooling from the β single-phase region.
本発明では、α結晶粒の（０００１）面の法線方向がチタン合金棒材の長軸方向に集積することを低減することを狙いとした。すなわち、α結晶粒の集合組織が形成されるβ水冷後のα＋β鍛造で、ビレット軸方向への延伸を促進させる加工を行い、チタン合金棒材の長軸方向に（１０−１０）面法線方向が集積することを促進させた。これにより、航空機エンジン部品に使用される素材に適したチタン合金棒材になる。 The present invention aims to reduce the accumulation of the normal direction of the (0001) plane of the α crystal grains in the semimajor direction of the titanium alloy bar. That is, by α + β forging after β-water cooling to form an aggregated structure of α crystal grains, processing is performed to promote stretching in the billet axial direction, and the (10-10) plane normal is performed in the long axis direction of the titanium alloy bar. Promoted the accumulation of directions. This makes the titanium alloy rod suitable for the material used for aircraft engine parts.
以下、本実施形態のチタン合金棒材について説明する。
本実施形態のチタン合金棒材は、例えば、２５℃においてα相を主相としβ相を第２相とする金属組織を有するものがよい。すなわち、ＡＭＳ４９２８で規定される成分で形成されていてもよい。つまり、Ａｌ：５．５０〜６．７５質量％、Ｖ：３．５〜４．５質量％、Ｆｅ：０．０５〜０．４０質量％、Ｏ：０．０５〜０．２５質量％を含有し、残部がＴｉおよび不純物であってもよい。不純物としては、例えば、Ｎ：０．０８質量％以下、Ｃ：０．０８質量％以下、Ｈ：０．０１５質量％以下を含有してもよい。
Hereinafter, the titanium alloy rod material of the present embodiment will be described.
The titanium alloy bar of the present embodiment may have, for example, a metal structure having an α phase as a main phase and a β phase as a second phase at 25 ° C. That is, it may be formed of the components defined by AMS4928. That is, Al: 5.50 to 6.75% by mass, V: 3.5 to 4.5% by mass, Fe: 0.05 to 0.40% by mass, O: 0.05 to 0.25% by mass. It may be contained and the balance may be Ti and impurities. As the impurities, for example, N: 0.08% by mass or less, C: 0.08% by mass or less, and H: 0.015% by mass or less may be contained.
また、本実施形態のチタン合金棒材は、例えば、ＡＭＳ４９７５で規定される成分で形成されていてもよい。つまり、Ａｌ：５．５０〜６．５０質量％、Ｓｎ：１．７５〜２．２５質量％、Ｚｒ：３．５〜４．５質量％、Ｍｏ：１．８〜２．２質量％、Ｆｅ：０．０２〜０．２５質量％、Ｏ：０．０２〜０．１５質量％を含有し、残部がＴｉおよび不純物であってもよい。不純物としては、例えば、Ｓｉ：０．１０質量％以下、Ｎ：０．０８質量％以下、Ｃ：０．０８質量％以下、Ｈ：０．０１５質量％以下を含有していてもよい。 Further, the titanium alloy bar of the present embodiment may be formed of, for example, the components specified by AMS4975. That is, Al: 5.50 to 6.50% by mass, Sn: 1.75 to 2.25% by mass, Zr: 3.5 to 4.5% by mass, Mo: 1.8 to 2.2% by mass, Fe: 0.02 to 0.25% by mass and O: 0.02 to 0.15% by mass may be contained, and the balance may be Ti and impurities. As the impurities, for example, Si: 0.10% by mass or less, N: 0.08% by mass or less, C: 0.08% by mass or less, and H: 0.015% by mass or less may be contained.
本実施形態のチタン合金棒材の形状は、円柱状の棒材でもよく、多角形状の棒材でもよい。チタン合金棒材の長軸方向に直交する断面は円の場合、真円であってもよいが、真円である必要はなく、おおよそ円形状であれば良い。多角形状の場合もおおよそ多角形であればよい。 The shape of the titanium alloy bar of the present embodiment may be a columnar bar or a polygonal bar. In the case of a circle, the cross section of the titanium alloy rod material orthogonal to the long axis direction may be a perfect circle, but it does not have to be a perfect circle and may be approximately a circular shape. In the case of a polygon, it may be approximately a polygon.
一方で、鋳塊から棒材に製造されるまでの中間形態の形状については、長軸方向に直交する断面形状は円形状に限定されず、四角形や八角形の多角形や、角が丸い多角形であってもよい。 On the other hand, regarding the shape of the intermediate form from the ingot to the production of the bar, the cross-sectional shape orthogonal to the long axis direction is not limited to the circular shape, but is a quadrangular or octagonal polygon, or many with rounded corners. It may be polygonal.
次に、本実施形態のチタン合金棒材の結晶組織について図１を参照しながら説明する。
本実施形態のチタン合金棒材は、長軸方向の断面において、α結晶粒を構成する稠密六方結晶の（０００１）面の法線方向と、長軸方向とのなす角度θ１が０°以上２５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が５．０％以下であることが好ましい。すなわち、チタン合金棒材の長軸方向に対して稠密六方結晶のｃ軸が０〜２５°の範囲で傾斜し、かつ、円相当直径が２０μｍ超であるα結晶粒が、長軸方向の断面において５．０面積％の割合であることが好ましい。
Next, the crystal structure of the titanium alloy bar of the present embodiment will be described with reference to FIG.
In the cross section of the titanium alloy bar of the present embodiment, the angle θ 1 formed by the normal direction of the (0001) plane of the dense hexagonal crystal constituting the α crystal grain and the long axis direction is 0 ° or more. It is preferable that the area ratio of α crystal grains having a circle-equivalent diameter of more than 20 μm in the range of 25 ° or less is 5.0% or less. That is, α crystal grains in which the c-axis of the dense hexagonal crystal is inclined in the range of 0 to 25 ° with respect to the long-axis direction of the titanium alloy rod and the equivalent circle diameter is more than 20 μm are cross-sections in the long-axis direction. The ratio is preferably 5.0 area%.
Ｄｗｅｌｌ疲労では、稠密六方晶の底面（（０００１）面）にほぼ平行なファセット破面が形成され、き裂発生および進展が促進され、寿命低下に至る。このため、稠密六方結晶のｃ軸の方向（（０００１）面の法線方向）が棒材の長軸方向に対して大きく傾斜していることが好ましい。本実施形態では、ｃ軸の傾斜角度θ１が０〜２５°の範囲にあり、かつ円相当直径が２０μｍ超のα結晶粒が５．０％以下であれば、ファセット破面が形成される確率やファセット破面のサイズが減少し、Ｄｗｅｌｌ疲労を改善することができる。ｃ軸の傾斜角度θ１が０〜２５°の範囲にある円相当直径が２０μｍ超のα結晶粒が５．０％を超えると、Ｄｗｅｌｌ疲労が大幅に悪化するので好ましくない。 In Dwell fatigue, faceted fracture surfaces that are substantially parallel to the bottom surface ((0001) plane) of the dense hexagonal crystal are formed, which promotes crack generation and growth, leading to a decrease in life. Therefore, it is preferable that the c-axis direction (normal direction of the (0001) plane) of the dense hexagonal crystal is greatly inclined with respect to the long axis direction of the bar. In the present embodiment, if the inclination angle θ 1 of the c-axis is in the range of 0 to 25 ° and the number of α crystal grains having a circle-equivalent diameter of more than 20 μm is 5.0% or less, a facet fracture surface is formed. Probability and facet fracture size are reduced, and Dwell fatigue can be improved. If the amount of α crystal grains having a circle-equivalent diameter of more than 20 μm in the c-axis inclination angle θ 1 in the range of 0 to 25 ° exceeds 5.0%, Dwell fatigue is significantly deteriorated, which is not preferable.
また、通常の疲労破壊の起点になりうる粗大なα結晶粒が多く存在するとＤｗｅｌｌ疲労も悪化するので、本実施形態のチタン合金棒材では、稠密六方結晶の（０００１）面の法線方向と、長軸方向とのなす角度θ1が２５°以上５５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が２．０％以下であることが好ましい。これにより、Ｄｗｅｌｌ疲労をより改善できる。 Further, if there are many coarse α crystal grains that can be the starting point of normal fatigue fracture, Dwell fatigue also deteriorates. Therefore, in the titanium alloy rod material of the present embodiment, the normal direction of the (0001) plane of the dense hexagonal crystal It is preferable that the area ratio of α crystal grains having a circle-equivalent diameter of more than 20 μm in the range where the angle θ 1 formed with the major axis direction is 25 ° or more and 55 ° or less is 2.0% or less. Thereby, Dwell fatigue can be further improved.
更に、本実施形態のチタン合金棒材は、長軸方向の断面において、α結晶粒を構成する稠密六方結晶の（１０−１０）面の法線方向と、長軸方向とのなす角度θ２が０°以上３０°以下の範囲にあるα結晶粒の面積率が４０％以上であることが好ましい。稠密六方結晶の（１０−１０）面は、図１に示すように、六角柱形状の単位結晶格子の側面に当たる面であり、この面の面方向と棒材の長軸とのなす角度θ２が０〜３０°の範囲にあるα結晶粒が、長軸方向の直交断面において４０面積％以上あるとよい。４０面積％未満になると、Ｄｗｅｌｌ疲労が大幅に悪化するので好ましくない。 Further, the titanium alloy bar of the present embodiment has an angle θ 2 between the normal direction of the (10-10) plane of the dense hexagonal crystal constituting the α crystal grain and the major axis direction in the cross section in the major axis direction. It is preferable that the area ratio of α crystal grains in the range of 0 ° or more and 30 ° or less is 40% or more. As shown in FIG. 1, the (10-10) plane of the dense hexagonal crystal is a plane corresponding to the side surface of the hexagonal columnar unit crystal lattice, and the angle θ 2 between the plane direction of this plane and the long axis of the bar is formed. It is preferable that the α crystal grains in the range of 0 to 30 ° are 40 area% or more in the orthogonal cross section in the long axis direction. If it is less than 40 area%, Dwell fatigue is significantly deteriorated, which is not preferable.
本実施形態のチタン合金棒材の結晶組織は、ＥＢＳＤ（電子線後方散乱回折；Electron Backscatter Diffraction）を用いて測定することができる。 The crystal structure of the titanium alloy bar of the present embodiment can be measured using EBSD (Electron Backscatter Diffraction).
まず、チタン合金棒材の長さ方向中心部より、長さ方向断面を観察面とする試験片を採取する。観察面における測定箇所は、断面が半径ｒの円形の試料については表面からｒ／２の深さの位置とし、断面の辺長がｄの矩形の試料についてはその辺長がなす表面からｄ／４の深さの位置とする。次に、試験片の観察面の測定箇所における、縦３ｍｍ横３ｍｍの矩形の領域を視野とし、測定間隔は２．０μｍ、加速電圧１５ｋＶで、ＥＢＳＤを用いて測定する。 First, a test piece having a cross section in the length direction as an observation surface is collected from the central portion in the length direction of the titanium alloy rod. The measurement point on the observation surface is at a depth of r / 2 from the surface for a circular sample with a radius r, and d / from the surface formed by the side length for a rectangular sample with a cross section of d. The position is at a depth of 4. Next, the measurement point on the observation surface of the test piece is measured using an EBSD with a rectangular region of 3 mm in length and 3 mm in width as a visual field, a measurement interval of 2.0 μm, and an acceleration voltage of 15 kV.
得られた測定結果を、ＯＩＭ（株式会社 ＴＳＬソリューションズ製の結晶方位解析ソフト）を用いて解析する。まず、α相のみを対象とするＰａｒｔｉｔｏｎを作成し、解析の対象とする。 The obtained measurement results are analyzed using OIM (crystal orientation analysis software manufactured by TSL Solutions Co., Ltd.). First, a partion that targets only the α phase is created and analyzed.
次に、隣り合うＥＢＳＤ測定点の結晶方位の角度差（ミスオリエンテーション角）を５°以下としてα結晶粒を決定し、そのα結晶粒の測定点数から各α結晶粒の面積を求め、各α結晶粒の円相当直径を算出する。 Next, the α crystal grains are determined with the angle difference (misorientation angle) of the crystal orientations of the adjacent EBSD measurement points being 5 ° or less, the area of each α crystal grain is obtained from the number of measurement points of the α crystal grains, and each α Calculate the equivalent circle diameter of the crystal grains.
また、各α結晶粒内のＥＢＳＤ測定点におけるｃ軸方向の平均値を算出し、それを用いて各α結晶粒について、α結晶粒の（０００１）面の法線方向及び（１０−１０）面の法線方向と、チタン合金棒材の長軸方向とのなす角度θ１、θ２を算出する。 In addition, the average value in the c-axis direction at the EBSD measurement point in each α crystal grain is calculated, and the average value in the c-axis direction is used for each α crystal grain in the normal direction of the (0001) plane of the α crystal grain and (10-10). The angles θ 1 and θ 2 formed by the normal direction of the surface and the long axis direction of the titanium alloy rod are calculated.
そして、α結晶粒のうち、角度θ１が０°〜２５°のα結晶粒の面積率と、角度θ２が０〜３０°のα結晶粒の面積率とをそれぞれ求める。また、円相当直径が２０μｍ超のα結晶粒の面積率を求める。 Then, among the α crystal grains, the area ratio of the α crystal grains having an angle θ 1 of 0 ° to 25 ° and the area ratio of the α crystal grains having an angle θ 2 of 0 to 30 ° are obtained. In addition, the area ratio of α crystal grains having a circle-equivalent diameter of more than 20 μm is determined.
あるいは、ＰａｒｔａｔｉｏｎでＣｒｙｓｔａｌ Ｄｉｒｅｃｔｉｏｎ Ｍａｐを作成し、α結晶粒の（１０−１０）面の法線方向と、チタン合金棒材の長軸方向とのなす角度θ2が０°以上３０°以下の範囲にあるα結晶粒の面積率（Ｔｏｔａｌ Ｆｒａｃｔｉｏｎ）を求める。 Alternatively, a Crystal Direction Map is created by partitioning, and the angle θ 2 formed by the normal direction of the (10-10) plane of the α crystal grains and the long axis direction of the titanium alloy bar is in the range of 0 ° or more and 30 ° or less. The area ratio (Total Fraction) of α crystal grains in the above is obtained.
また、Ｐａｒｔａｔｉｏｎ ＰｒｏｐｅｒｔｉｅｓでＧｒａｉｎ Ｓｉｚｅを２０μｍ超とした後、Ｃｒｙｓｔａｌ Ｄｉｒｅｃｔｉｏｎ Ｍａｐを作成し、α結晶粒の（０００１）面の法線方向と、チタン合金棒材の長軸方向とのなす角度θ１が０°以上２５°以下の範囲にあるα結晶粒の面積率（Ｔｏｔａｌ Ｆｒａｃｔｉｏｎ）を求める。 Further, after setting the Grain Size to more than 20 μm in Partition Properties, a Crystal Direction Map was created, and the angle θ 1 formed by the normal direction of the (0001) plane of the α crystal grain and the long axis direction of the titanium alloy rod was set. The area ratio (Total Fraction) of α crystal grains in the range of 0 ° or more and 25 ° or less is obtained.
また、Ｐａｒｔａｔｉｏｎ ＰｒｏｐｅｒｔｉｅｓでＧｒａｉｎ Ｓｉｚｅを２０μｍ超とした後、Ｃｒｙｓｔａｌ Ｄｉｒｅｃｔｉｏｎ Ｍａｐを作成し、α結晶粒の（０００１）面の法線方向と、チタン合金棒材の長軸方向とのなす角度θ１が２５°以上５５°以下の範囲にあるα結晶粒の面積率（Ｔｏｔａｌ Ｆｒａｃｔｉｏｎ）を求める。 Further, after setting the Grain Size to more than 20 μm in Partition Properties, a Crystal Direction Map was created, and the angle θ 1 formed by the normal direction of the (0001) plane of the α crystal grain and the long axis direction of the titanium alloy rod was set. The area ratio (Total Fraction) of α crystal grains in the range of 25 ° or more and 55 ° or less is obtained.
次に、本実施形態のチタン合金棒材の製造方法について説明する。
α相とβ相の２相域で加工を行うと、α相およびβ相それぞれの集合組織が形成される。その後の冷過程でβ相の一部がα相に変態するが、そのα相はβ相の結晶方位に依存した方位関係（Ｂｕｒｇｅｒｓの関係）を有する。特にβ相の面積率が５０％程度を占める温度域においては、β相の集合組織の影響が冷却後も強く残存する。また、加工を加えた後のβ相からα相への変態では、生じうるα相の結晶方位のなかで特定の方位が高頻度で出現するバリアント選択を生じる。
Next, a method for manufacturing the titanium alloy bar of the present embodiment will be described.
When processing is performed in the two-phase region of α phase and β phase, textures of α phase and β phase are formed. A part of the β phase is transformed into the α phase in the subsequent cooling process, and the α phase has an orientation relationship (Burgers relationship) depending on the crystal orientation of the β phase. Especially in the temperature range where the area ratio of β phase occupies about 50%, the influence of the texture of β phase remains strongly even after cooling. In addition, the transformation from the β phase to the α phase after processing results in variant selection in which a specific orientation appears frequently among the possible α phase crystal orientations.
本実施形態で狙いとする棒材の長軸方向への（０００１）面方位の集積を抑制するためには、長軸方向への延伸を促進するように、加工することが望ましい。 In order to suppress the accumulation of the (0001) plane orientation in the long axis direction of the bar material targeted in the present embodiment, it is desirable to process the bar so as to promote the stretching in the long axis direction.
本実施形態のチタン合金棒材は、所定の化学成分に調整された原料を溶解して鋳塊を得た後、得られた鋳塊をβ単相域に加熱し加工するβ鍛造と、α＋β二相域に加熱して加工するα＋β鍛造とを経て得られたチタン合金ビレットを、以下の工程に供することで得られる。 The titanium alloy bar of the present embodiment is formed by β-forging, in which a raw material adjusted to a predetermined chemical composition is melted to obtain an ingot, and then the obtained ingot is heated to a β single-phase region for processing, and α + β. It is obtained by subjecting a titanium alloy billet obtained through α + β forging, which is processed by heating to a two-phase region, to the following steps.
本実施形態のチタン合金棒材は、所定の化学成分を有する上記チタン合金ビレットを、β単相域の温度に加熱した後に急冷する第１の工程と、チタン合金ビレットをα＋β二相域の温度に加熱し、鍛造した後に冷却する第２の工程と、チタン合金ビレットを、α＋β二相域の温度であって第２の工程の加熱温度以下の温度に加熱し、鍛造する第３の工程と、をこの順で行うことにより製造される。
以下、各工程について説明する。
The titanium alloy rod of the present embodiment has a first step of heating the titanium alloy billet having a predetermined chemical component to a temperature in the β single phase region and then quenching the titanium alloy billet, and a temperature of the titanium alloy billet in the α + β two phase region. The second step of heating and forging and then cooling, and the third step of heating the titanium alloy billet to a temperature in the α + β two-phase region, which is lower than the heating temperature of the second step, and forging. , Are manufactured in this order.
Hereinafter, each step will be described.
第１の工程では、チタン合金ビレットを加熱炉内でβ単相温度域に加熱し、その後、急冷することで、金属組織の均質化させ、結晶粒の粗大化を抑制する。β単相温度領域の加熱は、加熱炉内の温度をβ変態点温度より３０℃高い温度以上、β変態点温度より１００℃高い温度以下（β変態点温度＋３０℃〜β変態点温度＋１００℃の温度範囲）とすることが好ましい。加熱炉内の温度が、β変態点温度より３０℃高い温度であると、加熱炉内に温度が不均一な部分があったり、チタン合金ビレットの大きさが大きいものであったりしても、鋳塊全体がβ変態点温度以上に加熱されるため好ましい。また、加熱炉内の温度が、β変態点温度より１００℃高い温度以下であると、チタン合金ビレットの表層の酸化が抑制されるとともに、チタン合金ビレット中の金属組織の粗大化が抑制されるため、高品質のチタン合金棒材が得られる。 In the first step, the titanium alloy billet is heated to the β single-phase temperature range in the heating furnace and then rapidly cooled to homogenize the metal structure and suppress the coarsening of crystal grains. For heating in the β single-phase temperature region, the temperature inside the heating furnace is 30 ° C higher than the β transformation point temperature and 100 ° C higher than the β transformation point temperature (β transformation point temperature + 30 ° C to β transformation point temperature + 100 ° C. Temperature range). If the temperature inside the heating furnace is 30 ° C higher than the β transformation point temperature, even if there is a non-uniform temperature part in the heating furnace or the size of the titanium alloy billet is large, This is preferable because the entire ingot is heated to the β transformation point temperature or higher. Further, when the temperature in the heating furnace is 100 ° C. higher than the β transformation point temperature, the oxidation of the surface layer of the titanium alloy billet is suppressed and the coarsening of the metal structure in the titanium alloy billet is suppressed. Therefore, a high quality titanium alloy bar can be obtained.
第１の工程では、β単相温度域に加熱後、チタン合金ビレットを加熱炉から取り出して速やかに冷却するか、加工を終えた後に急冷することが好ましい。急冷は充分な冷却速度を得るために、十分な量の水にチタン合金ビレットを浸漬することで行う水冷が一般的であるが、水冷相当以上の冷却速度が得られる他の手段を用いても良い。急冷はチタン合金ビレットの表面温度が３００℃以下になるまで続けることが好ましい。 In the first step, it is preferable that the titanium alloy billet is taken out from the heating furnace and cooled immediately after heating to the β single-phase temperature range, or rapidly cooled after the processing is completed. Quench cooling is generally performed by immersing a titanium alloy billet in a sufficient amount of water in order to obtain a sufficient cooling rate, but water cooling that can obtain a cooling rate equivalent to or higher than that of water cooling can also be used. good. Quenching is preferably continued until the surface temperature of the titanium alloy billet is 300 ° C. or lower.
第１の工程では、β単相温度域に加熱して加熱炉から取り出した後に加工を行うことで、チタン合金ビレットに歪みを与えてもよい。歪みを与えることで再結晶を生じ、金属組織の結晶粒の粗大化が抑制される。 In the first step, the titanium alloy billet may be distorted by heating it in the β single-phase temperature range, taking it out of the heating furnace, and then processing it. Recrystallization occurs by giving strain, and coarsening of crystal grains of the metal structure is suppressed.
次に、第の２工程では、第１の工程後のチタン合金ビレットを、α＋β二相域の温度に加熱し、鍛造した後に冷却する。第２の工程では、被加工材料であるチタン合金ビレットがα相およびβ相の状態で加工される。特に、β相が組織中に５０％程度の割合で存在する温度域で加工することが好ましい。 Next, in the second step, the titanium alloy billet after the first step is heated to a temperature in the α + β two-phase region, forged, and then cooled. In the second step, the titanium alloy billet, which is the material to be processed, is processed in the α phase and β phase. In particular, it is preferable to process in a temperature range in which the β phase is present in the structure at a ratio of about 50%.
第２の工程において、チタン合金ビレットを加熱する加熱炉内の温度は、β変態点温度より６０℃低い温度以上、β変態点温度未満（β変態点温度−６０℃〜β変態点温度未満の温度範囲）とすることが好ましい。加工発熱による温度上昇を加味すると、加熱温度の上限はβ変態点温度より２０℃低い温度未満（β変態点温度−２０℃未満）であることが好ましい。 In the second step, the temperature in the heating furnace for heating the titanium alloy billet is 60 ° C. lower than the β transformation point temperature and lower than the β transformation point temperature (β transformation point temperature -60 ° C to less than β transformation point temperature). Temperature range) is preferable. Considering the temperature rise due to processing heat generation, the upper limit of the heating temperature is preferably less than 20 ° C. lower than the β transformation point temperature (β transformation point temperature less than −20 ° C.).
加熱炉内の温度が、β変態点温度より６０℃低い温度以上であると、熱間加工を施す際のチタン合金ビレットの変形抵抗が大きくなりすぎることを防止でき、容易に効率よく熱間加工を行うことができる。また、加熱炉内の温度が、β変態点温度未満であると、チタン合金ビレットの金属組織中にα結晶粒が十分に析出するため、粒成長が抑制されるとともに、α＋β二相温度域で熱間加工を施すことによる効果が十分に得られる。 When the temperature inside the heating furnace is 60 ° C. lower than the β transformation point temperature, it is possible to prevent the deformation resistance of the titanium alloy billet during hot working from becoming too large, and hot working can be done easily and efficiently. It can be performed. Further, when the temperature in the heating furnace is lower than the β transformation point temperature, α crystal grains are sufficiently precipitated in the metal structure of the titanium alloy billet, so that grain growth is suppressed and in the α + β two-phase temperature range. The effect of hot working can be fully obtained.
チタン合金ビレットの表面温度は鍛造中に徐々に低下するため、表面性状が悪化したり表面割れが生じやすくなったりする場合には、第２の工程の終了前に、鍛造を一旦中断し、再度、チタン合金ビレットを加熱してから鍛造を行うことが好ましい。 Since the surface temperature of the titanium alloy billet gradually decreases during forging, if the surface texture deteriorates or surface cracking is likely to occur, forging is temporarily interrupted and then again before the end of the second step. , It is preferable to heat the titanium alloy billet before forging.
第２の工程について、図２を参照して説明する。図２は、チタン合金ビレットと金敷とを示す図であり、図２（ａ）は圧下前の側面図であり、図２（ｂ）は圧下前の平面図であり、図２（ｃ）は圧下後の側面図であり、図２（ｄ）は圧下後の平面図である。図２において、符号１が金敷であり、符号２がビレットである。
The second step will be described with reference to FIG. 2A and 2B are views showing a titanium alloy billet and a metal floor, FIG. 2A is a side view before reduction, FIG. 2B is a plan view before reduction, and FIG. 2C is a plan view before reduction. It is a side view after reduction, and FIG. 2D is a plan view after reduction. In FIG. 2,
第２の工程では、ビレットの長軸方向とほぼ直交する方向から一対の金敷による圧下を加えて、ビレットを長軸方向に伸ばす鍛造、すなわち、鍛伸加工を行う。第２の工程によって、チタン合金中のα結晶粒の（０００１）面方位が、棒材の長軸方向に集積することを抑制する。 In the second step, forging, that is, forging is performed, in which a pair of metal fittings is applied from a direction substantially orthogonal to the long axis direction of the billet to extend the billet in the long axis direction. By the second step, the (0001) plane orientation of the α crystal grains in the titanium alloy is suppressed from accumulating in the long axis direction of the bar.
具体的には、ビレットの外周面の一部である被加工部位を金敷によって圧下した後、ビレットを長軸方向に所定の送り量だけ相対移動させ、金敷に新たな被加工部位を対向させ、この新たな被加工部位に対して圧下を行う。この動作を、ビレットの長手方向一端から他端に向けて順次行い、必要に応じて掴み替えを行い、ビレット全体に対して鍛造を行う。この間、ビレットは長軸方向に沿って金敷に対して相対的に送り出すのみであり、長軸中心に回転させることはしない。これにより、ビレットの外周面の一部に対して圧下が行われる。この操作を、１回の鍛造という。 Specifically, after the part to be processed, which is a part of the outer peripheral surface of the billet, is pressed down by the metal bed, the billet is relatively moved in the long axis direction by a predetermined feed amount, and the new machined part faces the metal bed. The reduction is performed on this new work piece. This operation is sequentially performed from one end to the other end in the longitudinal direction of the billet, and if necessary, the billet is re-grasped and the entire billet is forged. During this time, the billet is only sent out relative to the gold sill along the long axis direction, and is not rotated about the center of the long axis. As a result, reduction is performed on a part of the outer peripheral surface of the billet. This operation is called one forging.
１回目の鍛造が終了したら、ビレットをその長軸を中心にして回転させる。これにより、ビレットの外周面のうち、１回目の被加工部位とは別の被加工部位を金敷に向けさせる。次いで、２回目の鍛造を行う。たとえば、矩形断面の場合には９０°の異なる方向から圧下し、八角形断面の場合には４５°毎の方向から圧下を加えるとよい。 After the first forging is complete, the billet is rotated about its long axis. As a result, of the outer peripheral surface of the billet, a portion to be processed different from the portion to be processed for the first time is directed to the metal floor. Then, the second forging is performed. For example, in the case of a rectangular cross section, the reduction may be applied from different directions of 90 °, and in the case of an octagonal cross section, the reduction may be applied from the direction of every 45 °.
２回目の鍛伸加工が終了したら、３回目、４回目の鍛造を順次行う。鍛造の回数の上限は第２の工程前後での鍛錬比で制限する。第２の工程前後での鍛錬比が１．５以上になるまで鍛造を繰り返す。 After the second forging process is completed, the third and fourth forgings are sequentially performed. The upper limit of the number of forgings is limited by the forging ratio before and after the second step. Forging is repeated until the forging ratio before and after the second step becomes 1.5 or more.
鍛造において、金敷で被加工部位を順次加工する際のチタン合金ビレットの送り量Ｌiniは、鍛伸加工前のチタン合金ビレットの幅Ｗiniとの関係で、Ｌini／Ｗiniが０．８０以下になるように制限する。ここで、鍛伸加工前のチタン合金ビレットの幅Ｗiniとは、図２（ｂ）に示すように、金敷の圧下方向からチタン合金ビレットを見た場合のチタン合金ビレットの最大投影幅である。チタン合金ビレットの送り量Ｌiniは、金敷によって圧下を受ける被加工部位の長さに相当する。チタン合金ビレットの送り量Ｌiniが大きくなると、金敷によって圧下を受ける被加工部位の長さが増し、これにより、金敷によって拘束を受ける領域が増大する。金敷による拘束領域が増大すると、ビレットの長軸方向への延びが抑制され、α結晶粒の方位を適切な方向に向けさせることができなくなる。このため、チタン合金ビレットの送り量Ｌiniを、幅Ｗiniとの関係でＬini／Ｗiniが０．８０以下になるように制限する必要がある。 In forging, the feed amount Lini of the titanium alloy billet when the parts to be processed are sequentially machined with the metal pad is 0.80 or less in Lini / Wini in relation to the width Wini of the titanium alloy billet before the forging process. Limit to. Here, the width Wini of the titanium alloy billet before the forging process is the maximum projected width of the titanium alloy billet when the titanium alloy billet is viewed from the rolling direction of the metal fitting, as shown in FIG. 2 (b). The feed amount Lini of the titanium alloy billet corresponds to the length of the work portion to be pressed by the metal pad. As the feed amount Lini of the titanium alloy billet increases, the length of the work portion to be pressed by the metal pad increases, which increases the area constrained by the metal pad. When the restraint region by the gold sill is increased, the extension of the billet in the semimajor axis direction is suppressed, and the orientation of the α crystal grains cannot be oriented in an appropriate direction. Therefore, it is necessary to limit the feed amount Lini of the titanium alloy billet so that Lini / Wini is 0.80 or less in relation to the width Wini.
更に、鍛造では、１回毎に、鍛造後のチタン合金ビレットの高さＨafterと幅Ｗafterとの比Ｈafter／Ｗafterが０．６７以上１．５以下となるように、かつ、ＷiniとＷafterとの比ΔＷ（ΔＷ＝Ｗafter／Ｗini）が１．０５以上１．１５以下になるように圧下することが好ましい。なお、鍛伸加工後のチタン合金ビレットの幅Ｗafterは、圧下終了時に一方の金敷と接触しているビレットの周長である。これらの条件はいずれも、ビレットを長軸方向に伸ばすための条件である。Ｈafter／Ｗafterは小さい方が好ましく、１．５以下の範囲がよいが、Ｈafter／Ｗafterを小さくし過ぎると、次回の鍛伸加工において加工前のＨini／Ｗiniが過大になってしまい、Ｈafter／Ｗafterを小さくすることができないので、Ｈafter／Ｗafterは０．６７以上とする。また、ΔＷが１．０５以上１．１５以下の範囲から外れると、ビレットを長軸方向ではなく幅方向に拡げるように加工してしまうので好ましくない。 Further, in the forging, the ratio Hafter / Wafter of the height Hafter and the width Wafter of the titanium alloy billet after forging is 0.67 or more and 1.5 or less, and the Wini and Wafter are combined. It is preferable to reduce the ratio so that the ratio ΔW (ΔW = Wafter / Wini) is 1.05 or more and 1.15 or less. The width Wafter of the titanium alloy billet after forging is the circumference of the billet that is in contact with one of the metal mats at the end of reduction. All of these conditions are conditions for extending the billet in the long axis direction. It is preferable that Hafter / Wafter is small, and the range of 1.5 or less is preferable. However, if Hafter / Wafter is made too small, Hini / Wini before processing becomes excessive in the next forging process, and Hafter / Wafter becomes excessive. Since it is not possible to reduce the value, Hafter / Wafter is set to 0.67 or more. Further, if ΔW deviates from the range of 1.05 or more and 1.15 or less, the billet is processed so as to expand in the width direction instead of the long axis direction, which is not preferable.
圧縮方向に加工する鍛造では、ビレットと金敷との接触によってビレットの変形が拘束され、変形の仕方が影響を受ける。拘束が強い方向には伸びにくく、拘束が弱い方向には伸びやすい。そこで、Ｌini／Ｗiniを０．８０以下とすることで、長軸方向の延伸を促進し、ビレットを圧下方向から見た場合の幅方向への（０００１）面方位の集積度を上昇させ、長軸方向への（０００１）面方位の集積度を低下させる。同時に、（１０−１０）面方位の長軸方向への集積度を上昇させる。また、この集積度の向上に伴い、ビレットの長軸方向に平行な（０００１）面方位を有するα結晶粒の面積率が低下し、長軸方向に平行な（１０−１０）面方位を有するα結晶粒の面積率が上昇する。 In forging that is processed in the compression direction, the deformation of the billet is restrained by the contact between the billet and the metal fitting, and the deformation method is affected. It is difficult to stretch in the direction of strong restraint, and it is easy to stretch in the direction of weak restraint. Therefore, by setting Lini / Wini to 0.80 or less, stretching in the semimajor direction is promoted, and the degree of integration of the (0001) plane orientation in the width direction when the billet is viewed from the reduction direction is increased, and the length is increased. It reduces the degree of integration of the (0001) plane orientation in the axial direction. At the same time, the degree of integration of the (10-10) plane orientation in the semimajor direction is increased. Further, as the degree of integration is improved, the area ratio of α crystal grains having a (0001) plane orientation parallel to the major axis direction of the billet decreases, and the area ratio has a (10-10) plane orientation parallel to the major axis direction. The area ratio of α crystal grains increases.
このような集積度あるいは面積率の変化を効率的に行うためには、第１の工程においてβ熱処理後に冷却することで結晶方位がランダム化されたチタン合金ビレットに対して、最初に行う加工を制御することが重要である。 In order to efficiently change the degree of integration or area ratio, the first processing is performed on the titanium alloy billet whose crystal orientation is randomized by cooling after β heat treatment in the first step. It is important to control.
つまり、β相の面積率が５０％程度となる温度域において、Ｌini／Ｗiniが０．８０以下になるようにチタン合金ビレットの送り量を制限しつつ、鍛伸加工後のチタン合金ビレットの高さＨafterと幅Ｗafterとの比Ｈafter／Ｗafterが０．６７以上１．５以下となるように、かつ、ＷiniとＷafterとの比ΔＷ（ΔＷ＝Ｗafter／Ｗini）が１．０５以上１．１５以下になるように圧下する鍛造を２回以上繰り返し、鍛錬比１．５以上となるまで行う。鍛錬比が１．５未満では、α結晶粒の集積度を向上させることができなくなる。 That is, in the temperature range where the area ratio of the β phase is about 50%, the height of the titanium alloy billet after forging is limited while limiting the feed amount of the titanium alloy billet so that Lini / Wini is 0.80 or less. The ratio Hafter / Wafter between Hafter and width Wafter is 0.67 or more and 1.5 or less, and the ratio ΔW (ΔW = Wafter / Wini) between Wini and Wafter is 1.05 or more and 1.15 or less. The forging is repeated twice or more so that the forging ratio becomes 1.5 or more. If the forging ratio is less than 1.5, the degree of accumulation of α crystal grains cannot be improved.
次に、第３の工程では、α＋β二相域の温度であって第２の工程の加熱温度以下の温度に加熱し、鍛錬比が３．０以上になるまで鍛造を行う。第３の工程では、第２の工程の加熱温度以下の温度で鍛造を行うことで、ビレットを圧下方向から見た場合の幅方向への（０００１）面方位の集積度をより高めさせ、長軸方向への（０００１）面方位の集積度を低下させ、同時に、（１０−１０）面方位の長軸方向への集積度を上昇させる。 Next, in the third step, the temperature is heated to a temperature in the α + β two-phase region and equal to or lower than the heating temperature in the second step, and forging is performed until the forging ratio becomes 3.0 or more. In the third step, forging is performed at a temperature equal to or lower than the heating temperature of the second step, so that the degree of integration of the (0001) plane orientation in the width direction when the billet is viewed from the reduction direction is further increased, and the length is increased. The degree of integration of the (0001) plane orientation in the axial direction is reduced, and at the same time, the degree of integration of the (10-10) plane orientation in the major axis direction is increased.
第３の工程において、チタン合金ビレットを加熱する加熱炉内の温度は、β変態点温度より８０℃低い温度以上、第２の工程の加熱温度以下とすることが好ましい。加工発熱による温度上昇を加味すると、加熱温度の上限はβ変態点温度より２０℃低い温度未満（β変態点温度−２０℃未満）であることが好ましい。 In the third step, the temperature in the heating furnace for heating the titanium alloy billet is preferably 80 ° C. lower than the β transformation point temperature and lower than the heating temperature in the second step. Considering the temperature rise due to processing heat generation, the upper limit of the heating temperature is preferably less than 20 ° C. lower than the β transformation point temperature (β transformation point temperature less than −20 ° C.).
加熱炉内の温度が、β変態点温度より８０℃低い温度以上であると、熱間加工を施す際のチタン合金ビレットの変形抵抗が大きくなりすぎることを防止でき、容易に効率よく熱間加工を行うことができる。また、加熱炉内の温度が、第２の工程の温度以上の温度になると、（０００１）面方位及び（１０−１０）面方位の集積度が低下してしまうので好ましくない。 When the temperature inside the heating furnace is 80 ° C. lower than the β transformation point temperature, it is possible to prevent the deformation resistance of the titanium alloy billet during hot working from becoming too large, and hot working easily and efficiently. It can be performed. Further, when the temperature in the heating furnace becomes higher than the temperature of the second step, the degree of integration of the (0001) plane orientation and the (10-10) plane orientation decreases, which is not preferable.
第３の工程においても、チタン合金ビレットの温度が鍛造中に徐々に低下するため、表面性状が悪化したり表面割れが生じやすくなったりする場合には、第３の工程の終了前に、鍛造を一旦中断し、再度、チタン合金ビレットを加熱してから鍛造することが好ましい。 Also in the third step, the temperature of the titanium alloy billet gradually decreases during forging, so if the surface texture deteriorates or surface cracks are likely to occur, forging is performed before the end of the third step. It is preferable to interrupt the process once, heat the titanium alloy billet again, and then forge.
第３の工程では、第２の工程の鍛造の場合と同様に、ビレットの外周面の一部である被加工部位を金敷によって圧下した後、ビレットを長軸方向に所定の送り量だけ相対移動させ、金敷に新たな被加工部位を対向させ、この新たな被加工部位に対して圧下を行う。この動作を、ビレットの長手方向一端から他端に向けて順次行い、ビレット全体に対して鍛造を行う。この間、ビレットは長軸方向に沿って金敷に対して相対的に送り出すのみであり、長軸中心に回転させることはしない。これにより、ビレットの外周面の一部に対して圧下が行われる。この操作を、１回の鍛造という。 In the third step, as in the case of forging in the second step, the part to be processed, which is a part of the outer peripheral surface of the billet, is pressed down by the metal pad, and then the billet is relatively moved in the long axis direction by a predetermined feed amount. Then, the new workpiece is opposed to the forging, and the new workpiece is pressed down. This operation is sequentially performed from one end to the other end in the longitudinal direction of the billet, and forging is performed on the entire billet. During this time, the billet is only sent out relative to the gold sill along the long axis direction, and is not rotated about the center of the long axis. As a result, reduction is performed on a part of the outer peripheral surface of the billet. This operation is called one forging.
１回目の鍛造が終了したら、ビレットをその長軸を中心にして回転させる。これにより、ビレットの外周面のうち、１回目の被加工部位とは別の被加工部位を金敷に向けさせる。次いで、２回目の鍛伸加工を行う。たとえば、矩形断面の場合には９０°の異なる方向から圧下し、八角形断面の場合には４５°毎の方向から圧下を加えるとよい。 After the first forging is complete, the billet is rotated about its long axis. As a result, of the outer peripheral surface of the billet, a portion to be processed different from the portion to be processed for the first time is directed to the metal floor. Then, the second forging process is performed. For example, in the case of a rectangular cross section, the reduction may be applied from different directions of 90 °, and in the case of an octagonal cross section, the reduction may be applied from the direction of every 45 °.
第３の工程では、第２の工程前後での鍛錬比が３．０以上になるまで鍛造を繰り返す。鍛錬比が３．０未満では、α結晶粒の大きさを微細化することができなくなり、疲労寿命が悪化する。 In the third step, forging is repeated until the forging ratio before and after the second step becomes 3.0 or more. If the forging ratio is less than 3.0, the size of α crystal grains cannot be refined, and the fatigue life deteriorates.
鍛造が終了したら、第３の工程の最後に、チタン合金ビレットを３００℃以下まで冷却する。３００℃以下まで冷却することにより、切断加工、品質検査、疵の手入れ等の精整作業を行うことができる。 After the forging is completed, at the end of the third step, the titanium alloy billet is cooled to 300 ° C. or lower. By cooling to 300 ° C. or lower, it is possible to perform rectification work such as cutting, quality inspection, and flaw maintenance.
以上説明したように、本実施形態のチタン合金棒材によれば、（０００１）面の法線方向と長軸方向とのなす角度θ１が０°以上２５°以下の範囲の円相当直径が２０μｍ超のα結晶粒の面積率が５．０％以下であり、（０００１）面の法線方向と長軸方向とのなす角度θ１が２５°以上５５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が２．０％以下であり、（１０−１０）面の法線方向と長軸方向とのなす角度θ２が０°以上３０°以下の範囲のα結晶粒の面積率が４０％以上であるので、Ｄｗｅｌｌ疲労特性を向上させることができる。 As described above, according to the titanium alloy bar of the present embodiment, the circle-equivalent diameter in the range of 0 ° or more and 25 ° or less in the angle θ 1 formed by the normal direction and the major axis direction of the (0001) plane is The area ratio of α crystal grains over 20 μm is 5.0% or less, and the angle θ 1 between the normal direction and the major axis direction of the (0001) plane is in the range of 25 ° or more and 55 ° or less. The area ratio of α crystal grains exceeding 20 μm is 2.0% or less, and the angle θ 2 formed by the normal direction and the major axis direction of the (10-10) plane is α in the range of 0 ° or more and 30 ° or less. Since the area ratio of the crystal grains is 40% or more, the Dwell fatigue characteristic can be improved.
また、本実施形態のチタン合金棒材の製造方法によれば、第１の工程、第２の工程及び第３の工程を順次行うことで、Ｄｗｅｌｌ疲労特性に優れたチタン合金棒材を製造できる。 Further, according to the method for producing a titanium alloy rod of the present embodiment, a titanium alloy rod having excellent Dwell fatigue characteristics can be produced by sequentially performing the first step, the second step, and the third step. ..
本実施形態のチタン合金棒材は、例えば、航空機エンジンのタービンブレードの素材として好適に用いることができる。すなわち、本実施形態のチタン合金棒材に対して更に加工を施してタービンブレードとすることで、Ｄｗｅｌｌ疲労特性に優れたタービンブレードとすることができる。 The titanium alloy bar of the present embodiment can be suitably used, for example, as a material for a turbine blade of an aircraft engine. That is, by further processing the titanium alloy rod material of the present embodiment to obtain a turbine blade, a turbine blade having excellent Dwell fatigue characteristics can be obtained.
次に、本発明の実施例について説明する。
以下に示す方法によりチタン合金棒材を製造し、評価した。
Next, examples of the present invention will be described.
A titanium alloy bar was produced and evaluated by the method shown below.
（事前工程）
溶解して得られた、表１に示す組成を有する直径約７５０ｍｍの円柱状の鋳塊を、β変態温度以上の１０２０℃以上１２００℃以下に加熱した加熱炉内でβ単相温度域に加熱した後、加熱炉から取り出して鍛造するβ鍛造と、β変態温度以下の９００℃以上９８０℃以下のα＋βの二相域に加熱した後、加熱炉から取り出して鍛造するα＋β鍛造を、それぞれ１回または複数回繰り返して、長手方向に直交する断面形状が表１に示す断面形状の棒状のビレットを得た。前記棒状のビレットを中間ビレット（チタン合金ビレット）とした。表１に示すチタン合金ビレットのβ変態点温度は９９０℃〜１０１０℃の範囲であった。
(Preliminary process)
A columnar ingot having the composition shown in Table 1 and having a diameter of about 750 mm obtained by melting is heated to a β single-phase temperature range in a heating furnace heated to 1020 ° C. or higher and 1200 ° C. or lower, which is higher than the β transformation temperature. After that, β forging, which is taken out from the heating furnace and forged, and α + β forging, which is heated to the two-phase region of α + β of 900 ° C. or higher and 980 ° C. or lower below the β transformation temperature, and then taken out from the heating furnace and forged, are performed once. Alternatively, it was repeated a plurality of times to obtain a rod-shaped billet having a cross-sectional shape shown in Table 1 having a cross-sectional shape orthogonal to the longitudinal direction. The rod-shaped billet was used as an intermediate billet (titanium alloy billet). The β transformation point temperature of the titanium alloy billet shown in Table 1 was in the range of 990 ° C to 1010 ° C.
なお、表１の中間ビレットの形状の欄において、例えば「ψ３６０」は、断面形状が直径３６０ｍｍの円形状であることを意味し、「４００＊４００」は断面形状が一辺長さ４００ｍｍの四角形であることを意味し、「６００＊３００」は断面形状が縦６００ｍｍ、横３００ｍｍの四角形であることを意味する。 In the column of the shape of the intermediate billet in Table 1, for example, "ψ360" means that the cross-sectional shape is a circular shape with a diameter of 360 mm, and "400 * 400" is a quadrangle with a cross-sectional shape of 400 mm on a side. "600 * 300" means that the cross-sectional shape is a quadrangle having a length of 600 mm and a width of 300 mm.
（第１の工程）
事前工程で得た中間ビレットを、表２に示す加熱温度の加熱炉内で加熱した後、加熱炉から取り出して、表２に示す条件のように、鍛造（加工）後に水冷、あるいは、鍛造（加工）を行わないで水冷した。水冷は、十分な量の水を入れた水槽に浸漬することで行った。また、水冷は、インゴット表面温度が少なくとも３００℃を下回る温度になるまで行った。第１の工程の加熱温度は、β変態点＋３０℃〜β変態点＋１００℃の範囲とした。
(First step)
The intermediate billet obtained in the pre-process is heated in the heating furnace at the heating temperature shown in Table 2, then taken out from the heating furnace, and water-cooled or forged (processed) and then forged (processed) according to the conditions shown in Table 2. It was water-cooled without processing). Water cooling was performed by immersing in a water tank filled with a sufficient amount of water. Further, water cooling was performed until the surface temperature of the ingot became at least 300 ° C. or lower. The heating temperature in the first step was in the range of β transformation point + 30 ° C. to β transformation point + 100 ° C.
（第２の工程）
第１の工程で得たチタン合金ビレットを、表２に示す加熱温度の加熱炉内で加熱した後、表２に示す鍛錬比になるまで鍛造した。第２の工程での加熱温度は、いずれの試料においても、β変態点温度−６０℃〜β変態点未満の範囲（α＋β二相域の温度）だった。鍛造は、ビレットの外周面の一部である被加工部位を金敷によって圧下した後、ビレットを長軸方向に所定の送り量だけ相対移動させ、金敷に新たな被加工部位を対向させ、この新たな被加工部位に対して圧下を行った。この動作を、ビレットの長手方向一端から他端に向けて順次行い、必要に応じて掴み替えを行い、ビレット全体に対して鍛造を行った。この間、ビレットは長軸方向に沿って金敷に対して相対的に送り出すのみであり、長軸中心に回転させることはしなかった。以上の操作を１回の鍛造とし、鍛造を１回行う毎にビレットを長軸回りに回転させることで鍛造時の圧下方向を各回毎に変更させた。このようにして、第２工程において表２に示す鍛錬比になるまで、２回以上の鍛造を行った。
(Second step)
The titanium alloy billet obtained in the first step was heated in a heating furnace having a heating temperature shown in Table 2 and then forged until the forging ratio shown in Table 2 was reached. The heating temperature in the second step was in the range of β transformation point temperature −60 ° C. to less than β transformation point (temperature in the α + β two-phase region) in all the samples. In forging, after the part to be processed, which is a part of the outer peripheral surface of the billet, is pressed down by the metal bed, the billet is relatively moved in the long axis direction by a predetermined feed amount, and the new machined part faces the metal bed. The compaction was applied to the part to be processed. This operation was sequentially performed from one end to the other end in the longitudinal direction of the billet, and the entire billet was forged by re-grabbing as necessary. During this time, the billet was only sent out relative to the gold sill along the long axis direction, and was not rotated around the long axis. The above operation was regarded as one forging, and the reduction direction at the time of forging was changed each time by rotating the billet around the long axis each time the forging was performed. In this way, forging was performed two or more times in the second step until the forging ratio shown in Table 2 was reached.
第２工程の後は、インゴット表面温度が少なくとも３００℃を下回る温度になるまで空冷（放冷）した。 After the second step, the ingot was air-cooled (released) until the surface temperature of the ingot was at least 300 ° C. or lower.
（第３の工程）
第２の工程後のビレットを、表２に示す加熱温度の加熱炉内で加熱した後、加熱炉から取り出して鍛造した。鍛造は、ビレットの外周面の一部である被加工部位を金敷によって圧下した後、ビレットを長軸方向に所定の送り量だけ相対移動させ、金敷に新たな被加工部位を対向させ、この新たな被加工部位に対して圧下を行った。この動作を、ビレットの長手方向一端から他端に向けて順次行い、必要に応じて掴み替えを行い、ビレット全体に対して鍛造を行った。この間、ビレットは長軸方向に沿って金敷に対して相対的に送り出すのみであり、長軸中心に回転させることはしなかった。その後、表２に示す鍛錬比になるまで、加熱炉での加熱と鍛造とを複数回繰り返して、断面形状が円形または多角形であるビレットを得た。また、鍛造を１回行う毎にビレットを長軸回りに回転させることで鍛造時の圧下方向を各回毎に変更させた。実施例のチタン合金ビレットの第３の工程での加熱温度は、いずれの試料においても、α＋β二相域の温度だった。
(Third step)
The billet after the second step was heated in a heating furnace having a heating temperature shown in Table 2, and then taken out from the heating furnace and forged. In forging, after the part to be processed, which is a part of the outer peripheral surface of the billet, is pressed down by the metal bed, the billet is relatively moved in the long axis direction by a predetermined feed amount, and the new machined part faces the metal bed. The compaction was applied to the part to be processed. This operation was sequentially performed from one end to the other end in the longitudinal direction of the billet, and the entire billet was forged by re-grabbing as necessary. During this time, the billet was only sent out relative to the gold sill along the long axis direction, and was not rotated around the long axis. Then, heating and forging in a heating furnace were repeated a plurality of times until the forging ratio shown in Table 2 was obtained to obtain a billet having a circular or polygonal cross section. Further, the billet was rotated around the long axis each time the forging was performed, so that the reduction direction at the time of forging was changed each time. The heating temperature of the titanium alloy billet of the example in the third step was the temperature in the α + β two-phase region in all the samples.
第３工程の後は、インゴット表面温度が少なくとも３００℃を下回る温度になるまで空冷（放冷）した。 After the third step, the ingot was air-cooled (released) until the surface temperature of the ingot became at least 300 ° C.
得られたチタン合金棒材について、結晶組織の測定を行った。
まず、チタン合金棒材の長さ方向中心部より、長さ方向断面を観察面とする試験片を採取した。観察面における測定箇所は、断面が半径ｒの円形の試料については表面からｒ／２の深さの位置とし、断面の辺長がｄの矩形の試料についてはその辺長がなす表面からｄ／４の深さの位置とした。次に、試験片の観察面の測定箇所における、縦３ｍｍ横３ｍｍの矩形の領域を視野とし、測定間隔は２．０μｍ、加速電圧１５ｋＶで、ＥＢＳＤを用いて測定した。
The crystal structure of the obtained titanium alloy rod was measured.
First, a test piece having a cross section in the length direction as an observation surface was collected from the central portion in the length direction of the titanium alloy rod. The measurement point on the observation surface is at a depth of r / 2 from the surface for a circular sample with a radius r, and d / from the surface formed by the side length for a rectangular sample with a cross section of d. The position was set to a depth of 4. Next, a rectangular region of 3 mm in length and 3 mm in width was used as a visual field at the measurement point on the observation surface of the test piece, the measurement interval was 2.0 μm, the acceleration voltage was 15 kV, and the measurement was performed using EBSD.
得られた測定結果を、ＯＩＭ（株式会社 ＴＳＬソリューションズ製の結晶方位解析ソフト）を用いて解析した。まず、α相のみを対象とするＰａｒｔｉｔｏｎを作成し、解析の対象とした。隣り合うＥＢＳＤ測定点の方位（ｃ軸方向）の角度差（ミスオリエンテーション角）を５°以下としてα結晶粒を決定した。 The obtained measurement results were analyzed using OIM (crystal orientation analysis software manufactured by TSL Solutions Co., Ltd.). First, a partion targeting only the α phase was created and used as a target for analysis. The α crystal grains were determined with the angle difference (misorientation angle) of the orientations (c-axis directions) of adjacent EBSD measurement points being 5 ° or less.
次に、ＰａｒｔａｔｉｏｎのＧｒａｉｎ ＰｒｏｐｅｒｔｉｅｓでＧｒａｉｎ Ｓｉｚｅを２０μｍ超とした後、Ｃｒｙｓｔａｌ Ｄｉｒｅｃｔｉｏｎ Ｍａｐを作成し、α結晶粒の（０００１）面の法線方向と、チタン合金棒材の径方向および周方向とのなす角度θ１が０°以上２５°以下の範囲にあるα結晶粒の面積率（Ｔｏｔａｌ Ｆｒａｃｔｉｏｎ）を求めた。 Next, after setting the Grain Size to more than 20 μm in the Grain Properties of the Partation, a Crystal Direction Map was created, and the normal direction of the (0001) plane of the α crystal grain and the radial direction and the circumferential direction of the titanium alloy rod were set. The area ratio (Total Fraction) of α crystal grains in which the angle θ 1 formed was in the range of 0 ° or more and 25 ° or less was determined.
また、ＰａｒｔａｔｉｏｎのＧｒａｉｎ ＰｒｏｐｅｒｔｉｅｓでＧｒａｉｎ Ｓｉｚｅを２０μｍ超とした後、Ｃｒｙｓｔａｌ Ｄｉｒｅｃｔｉｏｎ Ｍａｐを作成し、α結晶粒の（０００１）面の法線方向と、チタン合金棒材の径方向および周方向とのなす角度θ１が２５°以上５５°以下の範囲にあるα結晶粒の面積率（Ｔｏｔａｌ Ｆｒａｃｔｉｏｎ）を求めた。 Further, after setting the Grain Size to more than 20 μm in the Grain Properties of the Partation, a Crystal Direction Map is created, and the normal direction of the (0001) plane of the α crystal grain is formed with the radial direction and the circumferential direction of the titanium alloy rod. The area ratio (Total Fraction) of α crystal grains in the range where the angle θ 1 is 25 ° or more and 55 ° or less was determined.
また、ＰａｒｔａｔｉｏｎのＧｒａｉｎ ＰｒｏｐｅｒｔｉｅｓでＧｒａｉｎ Ｓｉｚｅを２０μｍ超とした後、Ｃｒｙｓｔａｌ Ｄｉｒｅｃｔｉｏｎ Ｍａｐを作成し、α結晶粒の（１０−１０）面の法線方向と、チタン合金棒材の径方向および周方向とのなす角θ２度が０°以上３０°以下の範囲にあるα結晶粒の面積率（Ｔｏｔａｌ Ｆｒａｃｔｉｏｎ）を求めた。 In addition, after setting the Grain Size to more than 20 μm in the Grain Properties of the Partation, a Crystal Direction Map was created to determine the normal direction of the (10-10) plane of the α crystal grain and the radial and circumferential directions of the titanium alloy rod. The area ratio (Total Fraction) of α crystal grains in the range of 0 ° or more and 30 ° or less of the angle θ 2 degrees between them was determined.
また、得られたチタン合金棒材のDwell疲労特性を測定した。
試験片として、チタン合金棒材の長軸方向が長手方向となるように引張試験片と疲労試験片を採取した。
In addition, the Dwell fatigue characteristics of the obtained titanium alloy rod were measured.
As test pieces, tensile test pieces and fatigue test pieces were taken so that the long axis direction of the titanium alloy rod was the longitudinal direction.
引張試験の測定条件は以下の通りとした。
試験片形状：平行部φ５×３０ｍｍ、ゲージ長さ２５ｍｍ、ひずみ速度：８．３×１０−５ｓ−１。
The measurement conditions for the tensile test were as follows.
Specimen shape: Parallel part φ5 × 30 mm, gauge length 25 mm, strain rate: 8.3 × 10-5 s -1 .
疲労試験の測定条件は以下の通りとした。
疲労試験片形状：平行部φ５．０８ｍｍ×１５．２４ｍｍ、ゲージ長さ１２ｍｍ。
疲労試験方法：軸力、片振り、応力比０．０５。最大応力＝同材料（同方向）の０．２％耐力の９５％。
通常疲労：三角波、負荷１ｓ、除荷１ｓ
Ｄｗｅｌｌ疲労：台形波、負荷１ｓ、保持１２０ｓ、除荷１ｓ
The measurement conditions for the fatigue test were as follows.
Fatigue test piece shape: Parallel part φ5.08 mm × 15.24 mm, gauge length 12 mm.
Fatigue test method: axial force, swing, stress ratio 0.05. Maximum stress = 0.2% proof stress of the same material (in the same direction) 95%.
Normal fatigue: triangular wave, load 1s, unloading 1s
Dwell fatigue: trapezoidal wave, load 1s, holding 120s, unloading 1s
表３に、α結晶粒の（０００１）面の法線方向と棒材の長軸方向とのなす角度θ１が０°以上２５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率、α結晶粒の（０００１）面の法線方向と棒材の長軸方向とのなす角度θ１が２５°以上５５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率、α結晶粒の（１０−１０）面の法線方向と棒材の長軸方向とのなす角度θ２が０°以上３０°以下の範囲にあるα結晶粒の面積率、チタン合金棒材のＤｗｅｌｌ疲労寿命比＝（通常疲労の破断寿命）／（Ｄｗｅｌｌ疲労の破断寿命）を示す。本発明の範囲にある実施例では、通常疲労の破断寿命は１６０００回以上であり、Ｄｗｅｌｌ疲労の破断寿命は８０００回以上であった。 Table 3 shows α crystal grains having a circle-equivalent diameter of more than 20 μm in which the angle θ 1 between the normal direction of the (0001) plane of the α crystal grain and the major axis direction of the bar is in the range of 0 ° or more and 25 ° or less. The angle θ 1 between the normal direction of the (0001) plane of the α crystal grain and the long axis direction of the bar is in the range of 25 ° or more and 55 ° or less, and the equivalent circle diameter of the α crystal grain is more than 20 μm. Area ratio of α crystal grains, the area ratio of α crystal grains in the range where the angle θ 2 between the normal direction of the (10-10) plane of the α crystal grains and the long axis direction of the bar is 0 ° or more and 30 ° or less, titanium The Dwell fatigue life ratio of the alloy rod = (normal fatigue fracture life) / (Dwell fatigue fracture life) is shown. In the examples within the scope of the present invention, the rupture life of normal fatigue was 16000 times or more, and the rupture life of Dwell fatigue was 8000 times or more.
表３に示すように、本発明の範囲にある実施例は、（通常疲労の破断寿命）／（Ｄｗｅｌｌ疲労の破断寿命）の値が２以下と小さく、通常の疲労特性に対するＤｗｅｌｌ疲労特性の低下代が小さくなっていることが分かる。一方、本発明の範囲外である比較例では、通常の疲労特性に対するＤｗｅｌｌ疲労特性の低下代が大きくなっていることが分かる。 As shown in Table 3, in the examples within the scope of the present invention, the value of (breaking life of normal fatigue) / (breaking life of Dwell fatigue) is as small as 2 or less, and the Dwell fatigue characteristic is lowered with respect to the normal fatigue characteristic. You can see that the cost is getting smaller. On the other hand, in the comparative example outside the scope of the present invention, it can be seen that the reduction allowance of the Dwell fatigue characteristic is larger than that of the normal fatigue characteristic.
１…金敷、２…ビレット。 1 ... Kinjiki, 2 ... Billet.
Claims (5)
α結晶粒を構成する稠密六方結晶の（０００１）面の法線方向と、前記チタン合金棒材の長軸方向とのなす角度θ１が０°以上２５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が５．０％以下であるとともに、
前記（０００１）面の法線方向と、前記長軸方向とのなす角度θ１が２５°以上５５°以下の範囲にある円相当直径が２０μｍ超のα結晶粒の面積率が２．０％以下であり、
かつ、α結晶粒を構成する稠密六方結晶の（１０−１０）面の法線方向のうちのひとつの方向と、前記長軸方向とのなす角度θ２が０°以上３０°以下の範囲にあるα結晶粒の面積率が４０％以上であることを特徴とする、チタン合金棒材。 A titanium alloy bar having a metal structure having an α phase as a main phase and a β phase as a second phase at 25 ° C.
The circle-equivalent diameter in which the angle θ 1 between the normal direction of the (0001) plane of the dense hexagonal crystal constituting the α crystal grain and the long axis direction of the titanium alloy bar is in the range of 0 ° or more and 25 ° or less The area ratio of α crystal grains over 20 μm is 5.0% or less, and
The area ratio of α crystal grains having a circle-equivalent diameter of more than 20 μm in the range where the angle θ 1 formed by the normal direction of the (0001) plane and the semimajor axis is 25 ° or more and 55 ° or less is 2.0%. Is below
In addition, the angle θ 2 formed by one of the normal directions of the (10-10) planes of the dense hexagonal crystals constituting the α crystal grain and the major axis direction is in the range of 0 ° or more and 30 ° or less. A titanium alloy rod having an area ratio of a certain α crystal grain of 40% or more.
前記チタン合金ビレットをα＋β二相域の温度に加熱し、前記チタン合金ビレットを鍛造した後に冷却する第２の工程と、
前記チタン合金ビレットを、α＋β二相域の温度であって前記第２の工程の加熱温度以下の温度に加熱し、前記チタン合金ビレットを鍛造する処理を１回以上行い、少なくとも最後に３００℃以下まで冷却する処理を行う第３の工程と、
をこの順で行う際に、
前記第２の工程における前記鍛造は、前記チタン合金ビレットを送り量Ｌiniで長軸方向に送りつつ金敷で圧下する加工であって、鍛造前の前記チタン合金ビレットの幅をＷiniとしたときにＬini／Ｗiniが０．８０以下を満たし、鍛造後の前記チタン合金ビレットの高さＨafterと幅Ｗafterとの比Ｈafter／Ｗafterが０．６７以上１．５以下となるように、かつ、前記Ｗiniと前記Ｗafterとの比ΔＷ（ΔＷ＝Ｗafter／Ｗini）が１．０５以上１．１５以下になるように圧下する鍛造であり、この鍛造を少なくとも２回以上行い、また、前記チタン合金ビレットを長軸周りに回転させて前記チタン合金ビレットに対する圧下方向を各回毎に変更させることとし、
前記第２の工程における鍛錬比を１．５以上とし、前記第３の工程の鍛錬比を３．０以上とする、
ことを特徴とする請求項１〜３のいずれか一項に記載のチタン合金棒材の製造方法。 A titanium alloy billet having a metal structure having an α phase as a main phase and a β phase as a second phase, obtained by hot working an ingot, is heated to a temperature in the β single phase region and then rapidly cooled. Step 1 and
A second step of heating the titanium alloy billet to a temperature in the α + β two-phase region, forging the titanium alloy billet, and then cooling the billet.
The titanium alloy billet is heated to a temperature in the α + β two-phase region and equal to or lower than the heating temperature of the second step, and the titanium alloy billet is forged at least once, and at least finally at 300 ° C. or lower. The third step of cooling to
When doing in this order
The forging in the second step is a process in which the titanium alloy billet is fed in the major axis direction with a feed amount of Lini and pressed down with a metal pad, and when the width of the titanium alloy billet before forging is Wini, the Lini. / Wini satisfies 0.80 or less, the ratio Hafter / Wafter of the height Hafter and the width Wafter of the forged titanium alloy billet is 0.67 or more and 1.5 or less, and the Wini and the said Forging is forging so that the ratio ΔW (ΔW = Wafter / Wini) with Wafter is 1.05 or more and 1.15 or less. This forging is performed at least twice, and the titanium alloy billet is placed around the long axis. It was decided to change the rolling direction with respect to the titanium alloy billet each time.
The forging ratio in the second step is 1.5 or more, and the forging ratio in the third step is 3.0 or more.
The method for producing a titanium alloy bar according to any one of claims 1 to 3.
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| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| CN114273581A (en) * | 2021-12-26 | 2022-04-05 | 贵州安大航空锻造有限责任公司 | Multidirectional forging forming method for titanium alloy complex die forging |
| CN114850367A (en) * | 2022-05-28 | 2022-08-05 | 西部超导材料科技股份有限公司 | Preparation method of Ti6246 titanium alloy large-size bar |
| CN117102273A (en) * | 2023-10-24 | 2023-11-24 | 成都先进金属材料产业技术研究院股份有限公司 | Titanium alloy seamless pipe and method for improving rotation bending fatigue performance thereof |
Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS63206442A (en) * | 1987-02-23 | 1988-08-25 | Hitachi Ltd | Alpha+beta-type ti alloy excellent in fatigue strength |
| JPH01242743A (en) * | 1988-03-23 | 1989-09-27 | Nkk Corp | Heat-resistant titanium alloy |
| JP2012224935A (en) * | 2011-04-22 | 2012-11-15 | Kobe Steel Ltd | Titanium alloy billet and method for producing titanium alloy billet, and method for producing titanium alloy forged material |
| JP2016199796A (en) * | 2015-04-13 | 2016-12-01 | 新日鐵住金株式会社 | Titanium alloy bar and method for producing the same |
-
2019
- 2019-03-20 JP JP2019053488A patent/JP7307314B2/en active Active
Patent Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPS63206442A (en) * | 1987-02-23 | 1988-08-25 | Hitachi Ltd | Alpha+beta-type ti alloy excellent in fatigue strength |
| JPH01242743A (en) * | 1988-03-23 | 1989-09-27 | Nkk Corp | Heat-resistant titanium alloy |
| JP2012224935A (en) * | 2011-04-22 | 2012-11-15 | Kobe Steel Ltd | Titanium alloy billet and method for producing titanium alloy billet, and method for producing titanium alloy forged material |
| JP2016199796A (en) * | 2015-04-13 | 2016-12-01 | 新日鐵住金株式会社 | Titanium alloy bar and method for producing the same |
Cited By (6)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| CN114273581A (en) * | 2021-12-26 | 2022-04-05 | 贵州安大航空锻造有限责任公司 | Multidirectional forging forming method for titanium alloy complex die forging |
| CN114273581B (en) * | 2021-12-26 | 2023-06-23 | 贵州安大航空锻造有限责任公司 | Multidirectional forging forming method for titanium alloy complex die forging |
| CN114850367A (en) * | 2022-05-28 | 2022-08-05 | 西部超导材料科技股份有限公司 | Preparation method of Ti6246 titanium alloy large-size bar |
| CN114850367B (en) * | 2022-05-28 | 2023-11-03 | 西部超导材料科技股份有限公司 | Preparation method of Ti6246 titanium alloy large-size bar |
| CN117102273A (en) * | 2023-10-24 | 2023-11-24 | 成都先进金属材料产业技术研究院股份有限公司 | Titanium alloy seamless pipe and method for improving rotation bending fatigue performance thereof |
| CN117102273B (en) * | 2023-10-24 | 2024-02-02 | 成都先进金属材料产业技术研究院股份有限公司 | Titanium alloy seamless pipe and method for improving rotation bending fatigue performance thereof |
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