JP5549981B2 - Mg-based alloy cold worked parts - Google Patents
Mg-based alloy cold worked parts Download PDFInfo
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- JP5549981B2 JP5549981B2 JP2010521754A JP2010521754A JP5549981B2 JP 5549981 B2 JP5549981 B2 JP 5549981B2 JP 2010521754 A JP2010521754 A JP 2010521754A JP 2010521754 A JP2010521754 A JP 2010521754A JP 5549981 B2 JP5549981 B2 JP 5549981B2
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- 229910045601 alloy Inorganic materials 0.000 title claims description 79
- 239000000956 alloy Substances 0.000 title claims description 79
- 239000013078 crystal Substances 0.000 claims description 32
- 238000005482 strain hardening Methods 0.000 claims description 29
- 239000006185 dispersion Substances 0.000 claims description 6
- 239000000470 constituent Substances 0.000 claims description 5
- 239000000203 mixture Substances 0.000 claims description 5
- 239000012535 impurity Substances 0.000 claims description 4
- 239000000126 substance Substances 0.000 claims description 4
- 229910052769 Ytterbium Inorganic materials 0.000 claims description 3
- 239000000463 material Substances 0.000 description 61
- 229910052727 yttrium Inorganic materials 0.000 description 29
- 239000011777 magnesium Substances 0.000 description 27
- 238000001125 extrusion Methods 0.000 description 26
- VWQVUPCCIRVNHF-UHFFFAOYSA-N yttrium atom Chemical compound [Y] VWQVUPCCIRVNHF-UHFFFAOYSA-N 0.000 description 20
- 238000000034 method Methods 0.000 description 19
- 230000000052 comparative effect Effects 0.000 description 14
- 230000006835 compression Effects 0.000 description 12
- 238000007906 compression Methods 0.000 description 12
- 238000010586 diagram Methods 0.000 description 12
- 238000012669 compression test Methods 0.000 description 11
- 238000003754 machining Methods 0.000 description 11
- 238000012360 testing method Methods 0.000 description 11
- FYYHWMGAXLPEAU-UHFFFAOYSA-N Magnesium Chemical compound [Mg] FYYHWMGAXLPEAU-UHFFFAOYSA-N 0.000 description 8
- 229910052749 magnesium Inorganic materials 0.000 description 8
- 238000000465 moulding Methods 0.000 description 8
- 230000008859 change Effects 0.000 description 7
- 238000001816 cooling Methods 0.000 description 7
- 238000009826 distribution Methods 0.000 description 7
- 229910000946 Y alloy Inorganic materials 0.000 description 6
- 238000012545 processing Methods 0.000 description 6
- 239000000523 sample Substances 0.000 description 6
- 238000007373 indentation Methods 0.000 description 5
- 238000000024 high-resolution transmission electron micrograph Methods 0.000 description 4
- 230000008569 process Effects 0.000 description 4
- 229910052761 rare earth metal Inorganic materials 0.000 description 4
- 229910000765 intermetallic Inorganic materials 0.000 description 3
- 229910052747 lanthanoid Inorganic materials 0.000 description 3
- 150000002602 lanthanoids Chemical class 0.000 description 3
- 239000011159 matrix material Substances 0.000 description 3
- XEEYBQQBJWHFJM-UHFFFAOYSA-N Iron Chemical compound [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 2
- 239000012298 atmosphere Substances 0.000 description 2
- 230000015572 biosynthetic process Effects 0.000 description 2
- 238000011156 evaluation Methods 0.000 description 2
- 238000004519 manufacturing process Methods 0.000 description 2
- 238000003825 pressing Methods 0.000 description 2
- 239000002994 raw material Substances 0.000 description 2
- 238000005728 strengthening Methods 0.000 description 2
- XUIMIQQOPSSXEZ-UHFFFAOYSA-N Silicon Chemical compound [Si] XUIMIQQOPSSXEZ-UHFFFAOYSA-N 0.000 description 1
- 229910001315 Tool steel Inorganic materials 0.000 description 1
- 239000012300 argon atmosphere Substances 0.000 description 1
- 230000008901 benefit Effects 0.000 description 1
- 238000005266 casting Methods 0.000 description 1
- 238000006073 displacement reaction Methods 0.000 description 1
- 238000000635 electron micrograph Methods 0.000 description 1
- 239000004519 grease Substances 0.000 description 1
- 229910052742 iron Inorganic materials 0.000 description 1
- 230000001788 irregular Effects 0.000 description 1
- 239000003562 lightweight material Substances 0.000 description 1
- 239000000314 lubricant Substances 0.000 description 1
- 238000005259 measurement Methods 0.000 description 1
- 238000002360 preparation method Methods 0.000 description 1
- 230000009467 reduction Effects 0.000 description 1
- 238000005070 sampling Methods 0.000 description 1
- 238000000926 separation method Methods 0.000 description 1
- 229910052710 silicon Inorganic materials 0.000 description 1
- 239000010703 silicon Substances 0.000 description 1
- 230000009466 transformation Effects 0.000 description 1
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 1
Classifications
-
- 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
- C22F1/06—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of magnesium or alloys based thereon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C23/00—Alloys based on magnesium
- C22C23/06—Alloys based on magnesium with a rare earth metal as the next major constituent
-
- 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
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- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Forging (AREA)
- Extrusion Of Metal (AREA)
Description
本発明は、イットリウム等のランタノイド系希土類元素が添加されたＭｇ基合金に関し、塑性加工が容易なＭｇ基合金に関する。 The present invention relates to an Mg-based alloy to which a lanthanoid rare earth element such as yttrium is added, and to an Mg-based alloy that can be easily plastically processed.
従来この種のＭｇ基合金では、冷間（室温程度の温度）温度域での塑性加工性が困難であったために構造用の軽量材料等への利用が望まれながら実現することが困難であった。 Conventionally, with this type of Mg-based alloy, plastic workability in the cold (temperature around room temperature) temperature range has been difficult, so it has been difficult to achieve it while being desired to be used for structural lightweight materials. It was.
本発明は、冷間塑性加工に必要な負荷加重を、著しく引き下げることができるＭｇ基合金冷間加工部材を提供し、その実用化を可能にすることを目的とする。 本発明は、Ｍｇ基合金を冷間加工にて所定の形状に成型したＭｇ基合金冷間加工部材であって、構成するＭｇ基合金は、Ｙを０．１ａｔ％以上３．０ａｔ％以下、および残部をＭｇと不可避的不純物から成る化学組成を有すると共に、前記Ｍｇ基合金は、温間で相当塑性ひずみ１以上が加えられ、その後３００〜５５０℃の範囲で等温保持が行われたＭｇ基合金に対して、さらに、冷間加工を公称ひずみにて０．１５以上を付与され、当該Ｍｇ基合金の組織中に当該冷間加工により分割され微細化された結晶粒が含まれて、その結晶粒径の平均値が３０μｍ以下であると共に、平均直径サイズとして２〜５０ｎｍ、分散間隔が２〜５０ｎｍのＹの高濃度領域を有することを特徴とする。
本発明のＭｇ基合金加工部材において、構成するＭｇ基合金にはランタノイド系希土類元素として、Ｙを０．１ａｔ％以上３．０ａｔ％以下とするのに代えて、Ｙｂを０．３ａｔ％としてもよい。
また、本発明のＭｇ基合金加工部材において、構成するＭｇ基合金にはランタノイド系希土類元素として、Ｙを０．１ａｔ％以上３．０ａｔ％以下とするのに代えて、Ｇｄを０．３ａｔ％としてもよい。
以上のような特徴のある内部構造により、ＡＺ３１合金に代表される従来展伸合金で通常認められる変形の異方性が解消され、例えば、引張荷重が作用する場合での降伏応力、すなわち塑性変形開始応力が、圧縮荷重が作用する場合での塑性変形開始応力の１．２〜１．４倍必要となるデメリットが解消される。 本合金は、変形の等方性を有しており、一定の荷重に対して、あらゆる方向へ同等の変形量を示す。同時に、変形加工に必要とされる荷重も負荷方向に依存せず、同等である。
An object of the present invention is to provide an Mg-based alloy cold-worked member capable of significantly reducing the load applied to cold plastic working, and to enable its practical use. The present invention is an Mg-based alloy cold-worked member formed by cold-working an Mg-based alloy into a predetermined shape, and the Mg-based alloy to be configured has a Y of 0.1 at% or more and 3.0 at% or less, And the remaining Mg has a chemical composition composed of Mg and inevitable impurities , and the Mg-based alloy is warmly applied with an equivalent plastic strain of 1 or more, and then kept isothermally in the range of 300 to 550 ° C. the alloy further granted 0.15 or more at strain nominal cold working, contains divided miniaturized crystal grain by the cold working to tissues of the Mg based alloy, the The crystal grain size has an average value of 30 μm or less, a high concentration region of Y having an average diameter size of 2 to 50 nm and a dispersion interval of 2 to 50 nm .
In the Mg-based alloy processed member of the present invention, the constituent Mg-based alloy may be a lanthanoid rare earth element, and Y may be 0.3 at% instead of 0.1 at% or more and 3.0 at% or less. Good .
Further, in the Mg-based alloy processed member of the present invention, the Mg-based alloy constituting the lanthanoid-based rare earth element has Gd of 0.3 at% instead of Y being 0.1 at% or more and 3.0 at% or less. It is good .
Due to the internal structure as described above, the anisotropy of deformation normally observed in a conventional wrought alloy represented by the AZ31 alloy is eliminated. For example, the yield stress when a tensile load is applied, that is, plastic deformation The disadvantage that the starting stress is required to be 1.2 to 1.4 times the plastic deformation starting stress when the compressive load is applied is eliminated. This alloy has an isotropic property of deformation and exhibits the same amount of deformation in all directions for a given load. At the same time, the load required for deformation processing does not depend on the load direction and is equivalent.
図１は、本合金（実施例４）を構成する内部組織の高分解能透過型電子顕微鏡写真である。
図２は、本合金（実施例４）を構成する内部組織のＺコントラスト法による高分解能透過型電子顕微鏡写真である。
図３は、上図は、本合金（実施例４）について、三次元アトムプローブにより観察したイットリウム原子の存在箇所をドットで表示した写真である。下図は、上図存在分布を元に、イットリウム原子が高濃度に偏在している領域をグレー色のコンター図でしめしている模式図である。
図４は、本合金の圧縮公称応力−公称ひずみ関係を、実施例４のＭｇ−０．６原子％Ｙ合金について示したグラフである。
図５は、図４に示した圧縮試験のうち、押出しと平行方向に、公称ひずみ０．４、すなわち、初期高さの６０％まで圧縮変形を加えたものから試験片を採取し、図４の場合と同様に静的に圧縮試験を行ったときのグラフである。
図６は、冷間加工性を評価するための金型セットの縦断正面図である。
図７は、図６に示す治具を用いて、冷間加工性を評価したグラフ。実施例１，実施例２，実施例４および比較例１に示す材料の結果を示す。
図８は、成形加工後の試料断面写真である。ここでは、押し込み速度０．０３ｍｍ／秒とし、４．５ｔｏｎを負荷した結果について示す。
図９は、成形加工後の試料断面写真を示す。ここでは、押し込み速度０．０３ｍｍ／秒とし、４．５ｔｏｎを負荷した結果について示す。上図はＡＺ３１合金、下図は本合金のＭｇ−０．６原子％Ｙ合金の例である。
図１０は、Ｍｇ−０．６原子％Ｙを４２５℃にて押出し、４００℃にて２４時間保持した材料の圧縮変形前後の結晶粒方位分布変化と平均結晶粒径（ｄ）を示す。ここでは、成形前、および、４％（公称ひずみ０．０４）、１５％（公称ひずみ０．１５）、２５％（公称ひずみ０．２５）変形の後に形成される内部組織を示している。
図１１は、図１０と同様材を４５％（公称ひずみ０．４５））変形の後に形成される内部組織および平均結晶粒径（ｄ）を示している。
図１２は、図１０と同様材を１５％（公称ひずみ０．１５））変形の後に形成される内部組織の拡大図を示している。
図１３は、従来材であるＡＺ３１合金を２５０℃で押出した後に４００℃にて２４時間保持した比較材を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、材料内部に形成される組織の変化を示している。
図１４は、Ｍｇ−０．６原子％Ｙを３２０℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、材料内部に形成される組織の変化を示している。
図１５は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、材料内部に形成される組織の変化を示している。
図１６は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料、ならびに、Ｍｇ−０．３原子％Ｙを３００℃にて押出し、４００℃にて２４時間保持した材料、を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒、ならびに３．０ｍｍ／秒にて、冷間加工した後に、形成されるボス状突起部の内部組織を冷間加工例として示している。
図１７は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料、ならびに、Ｍｇ−０．３原子％Ｙを３００℃にて押出し、４００℃にて２４時間保持した材料、を図６に示す方法で、上金型を速度３．０ｍｍ／秒にて、冷間加工した後に、形成されるボス状突起部の内部組織を冷間加工例として示している。
図１８は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、形成されるボス状突起部の内部組織を冷間加工例として示している。
図１９は、図６に記載の冷間加工方法で成形した後に形成される突起部の硬さを変形量の少ない部位と比較して示している。
図２０は、比較例として、純マグネシウムを３２８℃にて押出し、４００℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．１４で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。
図２１は、実施例として、Ｍｇ−０．３原子％Ｙを３００℃にて押出し、４００℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．４０で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。
図２２は、実施例として、Ｍｇ−１．０原子％Ｙを４２５℃にて押出し、４００℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．４０で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。
図２３は、実施例として、Ｍｇ−０．３原子％Ｙｂを３００℃にて押出し、４５０℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．４０で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。
図２４は、実施例として、Ｍｇ−０．３原子％Ｇｄを３００℃にて押出し、４５０℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．３５で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。
図２５は、比較例として、Ｍｇ−０．６原子％Ｙを押出比２５：１、温度３２０℃にて押出した材料の結晶粒組織を示している。図中の黒線は、結晶方位差が５°以上の界面を結晶粒界として示したものである。
図２６は、図２５で比較例として示したＭｇ−０．６原子％Ｙを押出比２５：１、温度３２０℃にて押出した材料から、押出平行および垂直方向に試験片を採取し、室温で圧縮試験した結果を示す。FIG. 1 is a high-resolution transmission electron micrograph of the internal structure constituting this alloy (Example 4).
FIG. 2 is a high-resolution transmission electron micrograph of the internal structure constituting this alloy (Example 4) by the Z contrast method.
FIG. 3 is a photograph in which the upper part of the present alloy (Example 4) is displayed with dots of yttrium atoms observed with a three-dimensional atom probe. The lower figure is a schematic diagram showing a region where yttrium atoms are unevenly distributed in a high concentration on the basis of the upper figure distribution by a gray contour diagram.
FIG. 4 is a graph showing the compression nominal stress-nominal strain relationship of this alloy for the Mg-0.6 atomic% Y alloy of Example 4.
FIG. 5 shows a test piece taken from the compression test shown in FIG. 4 in which a compressive deformation was applied to a nominal strain of 0.4, that is, 60% of the initial height in the direction parallel to the extrusion. It is a graph when performing a compression test statically similarly to the case of.
FIG. 6 is a longitudinal front view of a mold set for evaluating cold workability.
FIG. 7 is a graph showing an evaluation of cold workability using the jig shown in FIG. The result of the material shown in Example 1, Example 2, Example 4, and Comparative Example 1 is shown.
FIG. 8 is a cross-sectional photograph of the sample after the molding process. Here, the indentation speed is set to 0.03 mm / second, and 4.5 ton is shown.
FIG. 9 shows a cross-sectional photograph of the sample after the molding process. Here, the indentation speed is set to 0.03 mm / second, and 4.5 ton is shown. The upper figure is an example of an AZ31 alloy, and the lower figure is an example of an Mg-0.6 atomic% Y alloy of this alloy.
FIG. 10 shows the change in grain orientation distribution and the average crystal grain size (d) before and after compression deformation of the material extruded at 425 ° C. and held at 400 ° C. for 24 hours. Here, the internal structure formed before forming and after 4% (nominal strain 0.04), 15% (nominal strain 0.15), 25% (nominal strain 0.25) deformation is shown.
FIG. 11 shows the internal structure and average crystal grain size (d) formed after deformation of the material by 45% (nominal strain 0.45) as in FIG.
FIG. 12 shows an enlarged view of the internal structure formed after a 15% (nominal strain 0.15) deformation of the material as in FIG.
FIG. 13 shows a comparative material which was extruded at 250 ° C. and was held at 400 ° C. for 24 hours after extruding a conventional AZ31 alloy at a temperature of 0.0003 mm / sec. It shows the changes in the structure formed inside the material after processing.
FIG. 14 shows a material obtained by extruding Mg-0.6 atomic% Y at 320 ° C. and holding it at 400 ° C. for 24 hours by the method shown in FIG. 6 and cooling the upper mold at a speed of 0.0003 mm / sec. It shows the change in the structure formed inside the material after the interworking.
FIG. 15 shows a material obtained by extruding Mg-0.1 atomic% Y at 290 ° C. and holding it at 400 ° C. for 24 hours by the method shown in FIG. 6 and cooling the upper mold at a speed of 0.0003 mm / sec. It shows the change in the structure formed inside the material after the interworking.
FIG. 16 shows materials extruding Mg-0.1 atomic% Y at 290 ° C. and kept at 400 ° C. for 24 hours, and extruding Mg-0.3 atomic% Y at 300 ° C. at 400 ° C. The internal structure of the boss-shaped protrusion formed after the material held for 24 hours is cold worked at a speed of 0.0003 mm / second and 3.0 mm / second by the method shown in FIG. Is shown as an example of cold working.
FIG. 17 shows materials extruding Mg-0.1 atomic% Y at 290 ° C. and holding at 400 ° C. for 24 hours, and extruding Mg-0.3 atomic% Y at 300 ° C. at 400 ° C. 6 shows the internal structure of the boss-shaped protrusion formed as a cold working example after cold working the upper mold at a speed of 3.0 mm / second by the method shown in FIG. ing.
18 shows a material obtained by extruding Mg-0.1 atomic% Y at 290 ° C. and holding it at 400 ° C. for 24 hours by the method shown in FIG. 6 and cooling the upper mold at a speed of 0.0003 mm / sec. The internal structure of the boss-like projection formed after the hot working is shown as an example of cold working.
FIG. 19 shows the hardness of the protrusion formed after being formed by the cold working method shown in FIG. 6 in comparison with a portion having a small amount of deformation.
FIG. 20 shows, as a comparative example, a nominal stress-nominal strain curve (upper diagram) obtained when pure magnesium was extruded at 328 ° C. and compression-deformed at 400 ° C. for 24 hours, and the nominal strain. FIG. 6 shows a nominal stress-nominal strain curve obtained when the compression test piece taken by machining is compressed and deformed in a direction parallel to and perpendicular to extrusion after the deformation is stopped at 0.14.
FIG. 21 shows, as an example, a nominal stress-nominal strain curve obtained by compressing and deforming a material that was extruded at 300 ° C. and held at 400 ° C. for 24 hours. ), And after stopping the deformation at a nominal strain of 0.40, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in parallel and perpendicular to the extrusion. ing.
FIG. 22 shows, as an example, a nominal stress-nominal strain curve obtained by compressing and deforming a material which was extruded at 425 ° C. and held at 400 ° C. for 24 hours as an example (upper figure). ), And after stopping the deformation at a nominal strain of 0.40, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in parallel and perpendicular to the extrusion. ing.
FIG. 23 shows, as an example, a nominal stress-nominal strain curve (upper diagram) obtained when Mg-0.3 atomic% Yb was extruded at 300 ° C. and compression-deformed on a material held at 450 ° C. for 24 hours. ), And after stopping the deformation at a nominal strain of 0.40, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in parallel and perpendicular to the extrusion. ing.
FIG. 24 shows, as an example, a nominal stress-nominal strain curve obtained by compressing and deforming a material which was extruded at 300 ° C. and held at 450 ° C. for 24 hours as an example (upper diagram). ), And after stopping the deformation at a nominal strain of 0.35, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in the direction parallel and perpendicular to the extrusion. ing.
FIG. 25 shows a crystal grain structure of a material obtained by extruding Mg-0.6 atomic% Y at an extrusion ratio of 25: 1 and a temperature of 320 ° C. as a comparative example. The black line in the figure indicates an interface having a crystal orientation difference of 5 ° or more as a crystal grain boundary.
FIG. 26 shows specimens taken in parallel and perpendicular to the extrusion from a material obtained by extruding Mg-0.6 atomic% Y shown as a comparative example in FIG. 25 at an extrusion ratio of 25: 1 and a temperature of 320 ° C. The result of compression test is shown.
本発明の望ましい形態においては、Ｍｇ基合金は、その合金組織が、１μｍ３単位では、全体として均質ではあるが、１μｍ３内では、平均直径２〜５０ｎｍのＹ高濃度部が不規則に分散している。
本発明の更に望ましい形態においては、Ｍｇ基合金は、Ｙ高濃度部が１μｍ３単位でのＹ濃度の１．５倍以上の高濃度である。
本発明の材料の内部構造の特徴は、イットリウム原子が、素材中平均濃度の５割以上、すなわち、１．５倍以上の濃度で、高濃度に存在する領域が、平均直径２ｎｍ〜５０ｎｍの大きさを形成し、さらに、これら高濃度領域が、２ｎｍ〜５０ｎｍの間隔で素材結晶粒中に分散していることにある。
また、高濃度に分布するイットリウム原子は、母相であるマグネシウム原子と金属間化合物、すなわち、規則的な構造を形成せず、高濃度でありながらもランダムな分布を形成している。
本発明の材料は、冷間加工を公称ひずみにて０．１５以上（相当ひずみの絶対値として０．１７以上）加えることにより、内部結晶組織が分割され、微細化されることで、平均値が３０μｍ以下の結晶粒径を有することを特徴とする。
本発明のＭｇ基合金は、任意の長尺棒材、板材、ブロック材が製造可能である。従来、困難であると考えられていたマグネシウムの冷間加工性を確保することが可能となり、軽量構造材料として、あらゆる用途に貢献することが期待される。In a desirable form of the present invention, the Mg-based alloy is homogeneous as a whole when the alloy structure is 1 μm 3 units, but within 1 μm 3 the Y high concentration part having an average diameter of 2 to 50 nm is irregularly dispersed. doing.
In a more desirable embodiment of the present invention, the Mg-based alloy has a high concentration of 1.5 times or more the Y concentration in the unit of 1 μm 3 in the Y high concentration portion.
A feature of the internal structure of the material of the present invention is that a region where yttrium atoms are present at a high concentration at 50% or more of the average concentration in the material, that is, at a concentration of 1.5 times or more, is an average diameter of 2 nm to 50 nm. Further, these high concentration regions are dispersed in the material crystal grains at intervals of 2 nm to 50 nm.
Further, yttrium atoms distributed at a high concentration do not form an intermetallic compound with a magnesium atom as a parent phase, that is, form a regular distribution without forming a regular structure.
The material of the present invention is obtained by applying cold working at a nominal strain of 0.15 or more (the absolute value of the equivalent strain is 0.17 or more), so that the internal crystal structure is divided and refined to obtain an average value. Has a crystal grain size of 30 μm or less.
The Mg-based alloy of the present invention can produce any long bar material, plate material, and block material. It is possible to ensure the cold workability of magnesium, which has been considered difficult in the past, and is expected to contribute to all uses as a lightweight structural material.
＜合金の作製＞
イットリウム（Ｙ）と純マグネシウム（Ｍｇ）（純度９９．９５％）をアルゴン雰囲気にて完全に溶解し、鉄製鋳型に鋳込み、Ｙ含有量が０．１ａｔ％、０．３ａｔ％、０．６ａｔ％、１．０ａｔ％、１．２ａｔ％、１．５ａｔ％、２．０ａｔ％、２．２ａｔ％、３．０ａｔ％である９種類のＭｇ−Ｙ合金を作製した。表１に、実施例１〜１８、比較例１として各々示した。
得られた鋳造合金を、温度５００℃にて２４時間炉中保持（大気雰囲気）後に水冷することにより、溶体化処理を施した。
その後、機械加工により、直径４０ｍｍ、長さ７０ｍｍの円柱材とした。
この円柱材を表１に示す各押出し温度に保持したコンテナ内（大気中）で３０分間保持した後、押出し比２５：１にて押出しによる強ひずみ熱間加工を行なった。断面減少率から求めた平均相当塑性ひずみは３．７となる。
この押出し材を温度３００〜５５０℃の炉中に２４時間等温保持した後、炉外で空冷した。押出温度は表１に示す各温度を用いた。平均再結晶粒径（μｍ）、引張降伏応力（Ａ）、圧縮降伏応力（Ｂ）、降伏応力比（Ｂ／Ａ）、圧縮破断ひずみを測定した。結果をまとめて表１に示す。
しかしながら、部分的に配列が乱れている箇所が点在している。これは、原子半径が比較的大きいイットリウム原子が、マグネシウム母相原子中に分散して、配列単位である格子をゆがめているためである。
さらに、複数のイットリウム原子が不規則に集まることにより、格子のゆがみは著しくなる。その結果、図中に白色波線円、もしくは、白色波線楕円で示すような、著しく格子が歪められた領域が形成される。
従って、本合金ではイットリウム原子がマグネシウム母相原子と規則的な構造、いわゆる金属間化合物を形成しないで、イットリウムの高濃度領域を形成することが特徴である。
この格子がゆがんだ領域のサイズは、この図例で示したような電子顕微鏡写真などから測定することが可能である。測定の結果をもとに、格子が歪んだ領域の平均直径サイズ：２〜５０ｎｍ、分散間隔：２〜５０ｎｍであることが確認された。ただし、一部の領域では不可避の金属間化合物の形成が認められるため、半分以上のイットリウム原子がランダムな高濃度領域を形成することを本合金の特徴とする。
なお、イットリウムの濃度は、０．１原子％以上、３．０原子％以下の範囲とすることが出来る。
素材の製造方法は、所定のイットリウム濃度を含む合金を鋳造等により製造し、押出しなどにより、温間で相当塑性ひずみ１以上を素材に加え、その後、３００〜５５０℃の範囲で等温保持を行うことである。
図２は、本合金（実施例４）を構成する内部組織の高分解能透過型電子顕微鏡写真である。この観察写真は、Ｚコントラスト法という手法を用いて、マグネシウム母相原子より、重い原子であるイットリウム原子がコントラストの異なるドットとして表示されている。例えば、図中に白色波線円もしくは白色波線楕円で示した領域は、イットリウム原子が数多く集まっている領域を示している。
これら領域のサイズと分散間隔は、図１で示した、格子の歪んだ領域のものと一致するため、本合金の特徴ある構造は、イットリウム原子が高濃度に、かつ、ランダムに集まることにより形成されることは明白である。
図３の上図は、本合金（実施例４）について、三次元アトムプローブにより観察したイットリウム原子の存在箇所をドットで表示した写真である。下図は、上図存在分布を元に、イットリウム原子が高濃度に偏在している領域をグレー色のコンター図でしめしている模式図である。
ここでは、合金の例として、０．６原子％のイットリウムを含む合金に関する結果を示している。
例示合金におけるイットリウムの平均濃度は０．６原子％であり、ここでは１．０原子％以上の濃度領域、すなわち、平均濃度の１．６７倍以上の濃度となっている領域の大きさと分布間隔が示されている。
高濃度領域の大きさは、５〜１５ｎｍであり、領域間も５〜１５ｎｍとなっており、図１のひずみ領域や図２の高濃度領域と同様の結果となっている。
図４は、本合金の圧縮公称応力−公称ひずみ関係を、実施例４のＭｇ−０．６原子％Ｙ合金について示したグラフである。圧縮試験の方向は、押出材の押出方向に対して平行なもの（１８０°）、垂直なもの（９０°）、両者の中間であるもの（４５°）の３通りを選択している。
降伏応力、すなわち塑性変形開始応力は６０ＭＰａ程度で、その後の加工硬化、さらにはひずみ０．１２程度で加工硬化の勾配が緩やかになり、公称ひずみ０．４３〜０．５程度まで破断しない状態を維持する。３通りの方向に変形させた結果を比較すると明らかなように、変形の方向依存性は、ほとんど無いことがわかる。
以上に示すような、方向に依存しない圧縮変形挙動は、従来展伸材であるＡＺ３１合金などでは、発現しない特性である。
図５は、図４に示した圧縮試験のうち、押出しと平行方向に、公称ひずみ０．４、すなわち、初期高さの６０％まで圧縮変形を加えたものから試験片を採取し、図４の場合と同様に静的に圧縮試験を行ったときのグラフである。
その結果、初期ひずみとして公称圧縮ひずみ０．４を加えた材料は、変形開始応力が２００ＭＰａ程度まで増加することがわかった。また、試験結果から明らかなように、変形開始応力の大きさと、その後の加工硬化の割合は、圧縮試験方向によらず、同様であることがわかった。
また、図４の試験方向と同じ方向に変形させた材料では、公称破断ひずみが０．３７と大きな値が得られることがわかった。
以上の図４・図５に示す知見より、本合金は室温での塑性加工性、すなわち冷間加工性を有し、さらには、塑性加工後の機械的性質も良好であることが示唆された。
本合金の冷間加工性を明確にするために、図６に示すような金型セットを用いて、塑性加工性を評価した。
変形前の試験片は、直径８ｍｍ、高さ６ｍｍの円柱形状とし、同じ直径の円柱断面穴を有する工具鋼の下金型にセットした。その後、中心軸に直径３ｍｍの円柱断面穴と肩に半径１．５ｍｍのＲ部を有する上金型を試験片上面に接触させ、図の上から下に上金型を移動させることにより、下金型に拘束された試験片を上金型に設けた中心穴に沿って、塑性流動させることで、成形性を確認した。なお、試験片と金型の表面には、潤滑剤としてシリコングリスを塗布した。
成形試験の過程で、成形に必要な負荷荷重と上金型押し込み量を測定し、成形性の指標として用いた。
素材の成形性が十分であれば、図６に示すような上金型と同じ直径のボスが張り出す形に成形されるため、加工後の断面観察から、成形性と加工に伴う割れの形成の有無を直接、確認することが可能である。
図７は、図６に示す治具を用いて、冷間加工性を評価したグラフである。
試験片として、本合金の例として、実施例１のＭｇ−０．１原子％Ｙ、実施例２のＭｇ−０．３原子％Ｙ、実施例４のＭｇ−０．６原子％Ｙ、ならびに比較例１として、同一の条件でＡＺ３１合金を用いた結果を示している。
ここでは、上金型の押し込み速度として、０．０００３、０．０３ｍｍ／秒を選択した。
なお、本試験では、付与する最大荷重を４．５ｔｏｎ（４５ｋＮ）とした。
荷重と押し込み変位量の関係から明らかなように、同じ成形高さを得るための本合金の成形に必要な荷重は、従来材ＡＺ３１の場合と比較して、２割〜４割程度低いことがわかる。
図８は、成形加工後の試料断面写真を示す。ここでは、押し込み速度０．０００３ｍｍ／秒とし、４．５ｔｏｎを負荷した結果について示す。上図は、比較例１に示すＡＺ３１合金の場合で、Ｒ部を含む成形高さは、１．８ｍｍであった。
下図は、本合金（実施例４）のＭｇ−０．６原子％Ｙ合金の例であり、成形高さは３．７ｍｍであり、ＡＺ３１合金の２倍以上の成形高さが得られ、本合金の成形性が確認された。
図９は、成形加工後の試料断面写真を示す。ここでは、押し込み速度０．０３ｍｍ／秒とし、４．５ｔｏｎを負荷した結果について示す。上図は、比較例１に示すＡＺ３１合金の場合で、Ｒ部を含む成形高さは、１．４ｍｍであった。下図は、本合金（実施例４）のＭｇ−０．６原子％Ｙ合金の例であり、成形高さは２．９ｍｍであり、ＡＺ３１合金の２倍以上の成形高さが得られ、本合金の成形性が確認された。
以下に、前記結晶の変形を示す更なるデータにより、冷間加工の実態と、それによって強度が増した理由とを明示する。実験例により、当該現象が生じる数値的な限界を明示する。また、Ｙ以外の希土類元素についても、同様の現象が起こることを示す。
図１０は、Ｍｇ−０．６原子％Ｙを４２５℃にて押出し、４００℃にて２４時間保持した材料の圧縮変形前後の結晶粒方位分布変化と平均結晶粒径を示す。ここでは、成形前、および、４％（公称ひずみ０．０４）、１５％（公称ひずみ０．１５）、２５％（公称ひずみ０．２５）変形の後に形成される内部組織を示している。公称ひずみ０．１５以上の変形を加えた材料では、変形前の材料と比較して、平均結晶粒径が３０μｍ以下であり、微細化されている。
図１１は、図１０と同様材を４５％（公称ひずみ０．４５））変形の後に形成される内部組織を示している。図中の黒線で示した境界線は、結晶方位差が５度以上のものを結晶粒界として示している。図１０に示す変形前組織と比較して、結晶粒径が１／５に微細化されていることがわかる。
図１２は、図１０と同様材を１５％（公称ひずみ０．１５））変形の後に形成される内部組織の拡大図を示している。図中のＬおよびＴで示すライン上の結晶方位変化を右図の実線で示している。濃淡が変化している箇所、例えば図中の白抜き矢印で示した部位は、右下グラフから方位角が５度増加していることが明示されている。すなわち、本発明の冷間加工部材は、冷間加工を加えることによって、ある結晶粒内部の方位が変化してゆき、付与するひずみの増加とともに結晶方位差が大きくなり、やがては、結晶粒界が形成されることにより、結晶粒が分断され、材料内部の平均結晶粒径が微細化されていくことを裏付けるものである。
図１３は、従来材であるＡＺ３１合金を２５０℃で押出した後に４００℃にて２４時間保持した比較材を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、材料内部に形成される組織の変化を示している。加工材の中心部Ｄでは、結晶粒の分断が見られず、変形双晶がバンド状組織として、斜め方向に形成されている。
図１４は、Ｍｇ−０．６原子％Ｙを３２０℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、材料内部に形成される組織の変化を示している。加工材の中心部Ｄでは、変形双晶のような特定の方向へバンド状組織が形成されることなく、ランダムな方向へ新しい粒界が形成されていることがわかる。
図１５は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、材料内部に形成される組織の変化を示している。加工材の中心部Ｄでは、変形双晶のような特定の方向へバンド状組織が形成されることなく、ランダムな方向へ新しい粒界が形成されていることがわかる。
図１６は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料、ならびに、Ｍｇ−０．３原子％Ｙを３００℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒、ならびに３．０ｍｍ／秒にて、冷間加工した後に、形成されるボス状突起部の内部組織を冷間加工例として示している。
図１７は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料、ならびに、Ｍｇ−０．３原子％Ｙを３００℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度３．０ｍｍ／秒にて、冷間加工した後に、形成されるボス状突起部の内部組織を冷間加工例として示している。図１３の変形模式図中Ｄで示す部位の組織を示している。速度３．０ｍｍ／秒の加工速度で、同様の結晶粒微細化が起こることがわかる。
図１８は、Ｍｇ−０．１原子％Ｙを２９０℃にて押出し、４００℃にて２４時間保持した材料を図６に示す方法で、上金型を速度０．０００３ｍｍ／秒にて、冷間加工した後に、形成されるボス状突起部の内部組織を冷間加工例として示している。変形後組織の方位分布図より、結晶粒組織が微細化されていることがわかる。
図１９は、図６に記載の冷間加工方法で成形した後に形成される突起部の硬さが変形量の少ない部位と比較して、増加していることがわかる。冷間加工による結晶粒の微細化が強度増加をもたらしていることがわかる。
図２０は、比較例として、純マグネシウムを３２８℃にて押出し、４００℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．１４で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。降伏強度が大きく異なり、変形異方性が確認できる。
図２１は、実施例として、Ｍｇ−０．３原子％Ｙを３００℃にて押出し、４００℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．４０で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。降伏強度の異方性が低減されていることが確認できる。
図２２は、実施例として、Ｍｇ−１．０原子％Ｙを４２５℃にて押出し、４００℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．４０で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。降伏強度の異方性が低減されていることが確認できる。
図２３は、実施例として、Ｍｇ−０．３原子％Ｙｂを３００℃にて押出し、４５０℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．４０で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。降伏強度ならびに加工硬化率の異方性が低減されていることが確認できる。
図２４は、実施例として、Ｍｇ−０．３原子％Ｇｄを３００℃にて押出し、４５０℃にて２４時間保持した材料を圧縮変形させた際に得られる公称応力−公称ひずみ曲線（上図）、ならびに、公称ひずみ０．３５で変形を停止後、再度、機械加工により採取した圧縮試験片を、押出と平行および垂直方向へ圧縮変形させた際に得られる公称応力−公称ひずみ曲線を示している。降伏強度ならびに加工硬化率の異方性が低減されていることが確認できる。
図２５は、比較例として、Ｍｇ−０．６原子％Ｙを押出比２５：１、温度３２０℃にて押出した材料の結晶粒組織を示している。図中の黒線は、結晶方位差が５°以上の界面を結晶粒界として示したものである。図１１に示す実施例と比較して、当該比較例の冷間加工後組織では、図中の左部ならびに中央部に分割が不十分で粗大な結晶粒組織が残存していることがわかる。
図２６は、図２５で比較例として示したＭｇ−０．６原子％Ｙを押出比２５：１、温度３２０℃にて押出した材料から、押出平行および垂直方向に試験片を採取し、室温で圧縮試験した結果を示す。破断時の公称ひずみが０．１３以下であり、図４ならびに図５に示す本実施例合金と同様の組成を有しながらも、冷間加工性が低いことがわかる。また、降伏後の加工硬化率が採取方向により、大きく異なっており、かつ、押出平行方向に圧縮試験した際の破断直前の公称応力が、押出垂直方向のものと比較して２倍近い値を示していることから、変形異方性が強いことがわかる。<Preparation of alloy>
Yttrium (Y) and pure magnesium (Mg) (purity 99.95%) are completely dissolved in an argon atmosphere and cast into an iron mold, and the Y content is 0.1 at%, 0.3 at%, 0.6 at% , 1.0 at%, 1.2 at%, 1.5 at%, 2.0 at%, 2.2 at%, and 3.0 at%, nine kinds of Mg—Y alloys were produced. Table 1 shows Examples 1 to 18 and Comparative Example 1.
The obtained cast alloy was subjected to a solution treatment by holding it in a furnace at a temperature of 500 ° C. for 24 hours (air atmosphere) and then water cooling.
Thereafter, a cylindrical material having a diameter of 40 mm and a length of 70 mm was obtained by machining.
This columnar material was held in a container (in the atmosphere) held at each extrusion temperature shown in Table 1 for 30 minutes, and then subjected to high strain hot working by extrusion at an extrusion ratio of 25: 1. The average equivalent plastic strain obtained from the cross-sectional reduction rate is 3.7.
This extruded material was kept isothermal in a furnace at a temperature of 300 to 550 ° C. for 24 hours, and then air-cooled outside the furnace. Each temperature shown in Table 1 was used as the extrusion temperature. Average recrystallized grain size (μm), tensile yield stress (A), compressive yield stress (B), yield stress ratio (B / A), and compressive breaking strain were measured. The results are summarized in Table 1.
However, portions where the arrangement is partially disturbed are scattered. This is because yttrium atoms having a relatively large atomic radius are dispersed in the magnesium matrix phase atoms, and the lattice which is the arrangement unit is distorted.
Furthermore, the lattice distortion becomes significant due to the irregular collection of a plurality of yttrium atoms. As a result, a region in which the lattice is remarkably distorted as shown by a white wavy circle or a white wavy ellipse in the drawing is formed.
Therefore, the present alloy is characterized in that the yttrium atom forms a high-concentration region of yttrium without forming a regular structure with the magnesium matrix phase atom, that is, a so-called intermetallic compound.
The size of the region where the lattice is distorted can be measured from an electron micrograph as shown in this example. Based on the measurement results, it was confirmed that the average diameter size of the region where the lattice was distorted: 2 to 50 nm and the dispersion interval: 2 to 50 nm. However, since formation of inevitable intermetallic compounds is observed in some regions, it is a feature of this alloy that more than half of yttrium atoms form random high concentration regions.
Note that the concentration of yttrium can be in the range of 0.1 atomic% to 3.0 atomic%.
The raw material manufacturing method is to manufacture an alloy containing a predetermined yttrium concentration by casting or the like, add an equivalent plastic strain of 1 or more to the raw material by extrusion or the like, and then hold isothermally in the range of 300 to 550 ° C That is.
FIG. 2 is a high-resolution transmission electron micrograph of the internal structure constituting this alloy (Example 4). In this observed photograph, yttrium atoms that are heavier than magnesium matrix atoms are displayed as dots having different contrasts using a technique called Z contrast method. For example, a region indicated by a white wavy circle or a white wavy ellipse in the figure indicates a region where many yttrium atoms are gathered.
Since the size and dispersion interval of these regions coincide with those of the strained region of the lattice shown in FIG. 1, the characteristic structure of this alloy is formed by gathering yttrium atoms at a high concentration and randomly. It is clear that it will be done.
The upper diagram of FIG. 3 is a photograph showing the presence location of yttrium atoms observed with a three-dimensional atom probe for this alloy (Example 4). The lower figure is a schematic diagram showing a region where yttrium atoms are unevenly distributed in a high concentration on the basis of the upper figure distribution by a gray contour diagram.
Here, the result regarding the alloy containing 0.6 atomic% yttrium is shown as an example of the alloy.
The average concentration of yttrium in the example alloy is 0.6 atomic%. Here, the size and distribution interval of a concentration region of 1.0 atomic percent or more, that is, a region having a concentration of 1.67 times or more of the average concentration. It is shown.
The size of the high concentration region is 5 to 15 nm, and the region is 5 to 15 nm, which is the same result as the strain region of FIG. 1 and the high concentration region of FIG.
FIG. 4 is a graph showing the compression nominal stress-nominal strain relationship of this alloy for the Mg-0.6 atomic% Y alloy of Example 4. The direction of the compression test is selected from three types: one parallel to the extrusion direction of the extruded material (180 °), one perpendicular (90 °), and one intermediate between them (45 °).
The yield stress, that is, the plastic deformation starting stress is about 60 MPa, the subsequent work hardening, and further, the strain of work hardening becomes gentle when the strain is about 0.12, and the fracture does not break until the nominal strain is about 0.43 to 0.5. maintain. As is apparent from comparison of the results of deformation in three directions, it can be seen that there is almost no direction dependency of deformation.
The direction-independent compressive deformation behavior as described above is a characteristic that does not appear in AZ31 alloy, which is a conventional wrought material.
FIG. 5 shows a test piece taken from the compression test shown in FIG. 4 in which a compressive deformation was applied to a nominal strain of 0.4, that is, 60% of the initial height in the direction parallel to the extrusion. It is a graph when performing a compression test statically similarly to the case of.
As a result, it was found that the material to which the nominal compressive strain 0.4 was added as the initial strain increased the deformation start stress to about 200 MPa. Further, as apparent from the test results, it was found that the magnitude of the deformation starting stress and the subsequent work hardening ratio were the same regardless of the compression test direction.
In addition, it was found that a material deformed in the same direction as the test direction of FIG. 4 can obtain a large value at a nominal breaking strain of 0.37.
The above findings shown in FIGS. 4 and 5 suggest that this alloy has plastic workability at room temperature, that is, cold workability, and also has good mechanical properties after plastic working. .
In order to clarify the cold workability of this alloy, the plastic workability was evaluated using a mold set as shown in FIG.
The test piece before deformation had a cylindrical shape with a diameter of 8 mm and a height of 6 mm, and was set in a lower mold of tool steel having a cylindrical cross-sectional hole with the same diameter. Thereafter, an upper mold having a cylindrical cross-sectional hole with a diameter of 3 mm on the central axis and an R portion with a radius of 1.5 mm on the shoulder is brought into contact with the upper surface of the test piece, and the upper mold is moved downward from the top of the figure, The moldability was confirmed by plastically flowing the test piece restrained by the mold along the center hole provided in the upper mold. Silicon grease was applied as a lubricant to the surface of the test piece and the mold.
In the course of the molding test, the load required for molding and the amount of pressing of the upper mold were measured and used as an index of moldability.
If the formability of the material is sufficient, the boss with the same diameter as the upper mold as shown in Fig. 6 is formed into a shape that projects, so from the cross-sectional observation after processing, formability and formation of cracks accompanying processing It is possible to directly check for the presence or absence.
FIG. 7 is a graph showing an evaluation of cold workability using the jig shown in FIG.
As test pieces, examples of this alloy include Mg-0.1 atomic% Y of Example 1, Mg-0.3 atomic% Y of Example 2, Mg-0.6 atomic% Y of Example 4, and As Comparative Example 1, a result using an AZ31 alloy under the same conditions is shown.
Here, 0.0003 and 0.03 mm / second were selected as the pressing speed of the upper mold.
In this test, the maximum load to be applied was 4.5 ton (45 kN).
As is apparent from the relationship between the load and the amount of indentation displacement, the load required for forming the present alloy to obtain the same forming height is about 20% to 40% lower than that of the conventional material AZ31. Recognize.
FIG. 8 shows a cross-sectional photograph of the sample after the molding process. Here, it shows about the result of having applied indentation speed 0.0003mm / sec and loading 4.5ton. The upper figure shows the case of the AZ31 alloy shown in Comparative Example 1, and the molding height including the R portion was 1.8 mm.
The following figure is an example of the Mg-0.6 atomic% Y alloy of this alloy (Example 4), the forming height is 3.7 mm, and a forming height more than twice that of the AZ31 alloy is obtained. The formability of the alloy was confirmed.
FIG. 9 shows a cross-sectional photograph of the sample after the molding process. Here, the indentation speed is set to 0.03 mm / second, and 4.5 ton is shown. The upper figure shows the case of the AZ31 alloy shown in Comparative Example 1, and the molding height including the R portion was 1.4 mm. The following figure is an example of the Mg-0.6 atomic% Y alloy of this alloy (Example 4), the forming height is 2.9 mm, and a forming height more than twice that of the AZ31 alloy is obtained. The formability of the alloy was confirmed.
In the following, further data showing the deformation of the crystal will clarify the actual state of cold working and the reason why the strength is increased thereby. An experimental example clearly shows the numerical limits at which this phenomenon occurs. Moreover, it shows that the same phenomenon occurs also about rare earth elements other than Y.
FIG. 10 shows the grain orientation distribution change and the average crystal grain size before and after compression deformation of the material extruded with Mg-0.6 atomic% Y at 425 ° C. and held at 400 ° C. for 24 hours. Here, the internal structure formed before forming and after 4% (nominal strain 0.04), 15% (nominal strain 0.15), 25% (nominal strain 0.25) deformation is shown. In a material to which a deformation having a nominal strain of 0.15 or more is applied, the average crystal grain size is 30 μm or less as compared with the material before the deformation, and the material is refined.
FIG. 11 shows the internal structure formed after 45% (nominal strain 0.45) deformation of the material as in FIG. The boundary line indicated by the black line in the figure indicates a crystal grain boundary having a crystal orientation difference of 5 degrees or more. Compared to the structure before deformation shown in FIG. 10, it can be seen that the crystal grain size is refined to 1/5.
FIG. 12 shows an enlarged view of the internal structure formed after a 15% (nominal strain 0.15) deformation of the material as in FIG. The change in crystal orientation on the lines indicated by L and T in the figure is indicated by the solid line in the right figure. It is clearly shown in the lower right graph that the azimuth angle is increased by 5 degrees at the portion where the shading changes, for example, the portion indicated by the white arrow in the figure. That is, the cold-worked member of the present invention changes the orientation inside a certain crystal grain by applying cold work, and the crystal orientation difference increases with increasing strain to be applied. This confirms that the crystal grains are divided and the average crystal grain size inside the material is refined.
FIG. 13 shows a comparative material which was extruded at 250 ° C. and was held at 400 ° C. for 24 hours after extruding a conventional AZ31 alloy at a temperature of 0.0003 mm / sec. It shows the changes in the structure formed inside the material after processing. In the center portion D of the processed material, no crystal grain separation is observed, and deformation twins are formed in a diagonal direction as a band-like structure.
FIG. 14 shows a material obtained by extruding Mg-0.6 atomic% Y at 320 ° C. and holding it at 400 ° C. for 24 hours by the method shown in FIG. 6 and cooling the upper mold at a speed of 0.0003 mm / sec. It shows the change in the structure formed inside the material after the interworking. It can be seen that a new grain boundary is formed in a random direction at the center D of the workpiece without forming a band-like structure in a specific direction such as a deformation twin.
FIG. 15 shows a material obtained by extruding Mg-0.1 atomic% Y at 290 ° C. and holding it at 400 ° C. for 24 hours by the method shown in FIG. 6 and cooling the upper mold at a speed of 0.0003 mm / sec. It shows the change in the structure formed inside the material after the interworking. It can be seen that a new grain boundary is formed in a random direction at the center D of the workpiece without forming a band-like structure in a specific direction such as a deformation twin.
FIG. 16 shows materials extruding Mg-0.1 atomic% Y at 290 ° C. and kept at 400 ° C. for 24 hours, and extruding Mg-0.3 atomic% Y at 300 ° C. at 400 ° C. The material held for 24 hours was cold worked at a speed of 0.0003 mm / second and 3.0 mm / second by the method shown in FIG. It is shown as an example of cold working.
FIG. 17 shows materials extruding Mg-0.1 atomic% Y at 290 ° C. and holding at 400 ° C. for 24 hours, and extruding Mg-0.3 atomic% Y at 300 ° C. at 400 ° C. 6 shows the internal structure of the boss-like protrusion formed as a cold working example after cold working the upper mold at a speed of 3.0 mm / sec with the method shown in FIG. Yes. The structure | tissue of the site | part shown by D in the deformation | transformation schematic diagram of FIG. 13 is shown. It can be seen that similar grain refinement occurs at a processing speed of 3.0 mm / sec.
18 shows a material obtained by extruding Mg-0.1 atomic% Y at 290 ° C. and holding it at 400 ° C. for 24 hours by the method shown in FIG. 6 and cooling the upper mold at a speed of 0.0003 mm / sec. The internal structure of the boss-like projection formed after the hot working is shown as an example of cold working. From the orientation distribution map of the post-deformation structure, it can be seen that the crystal grain structure is refined.
FIG. 19 shows that the hardness of the protrusions formed after being formed by the cold working method shown in FIG. 6 is increased as compared with the part having a small amount of deformation. It can be seen that the refinement of crystal grains by cold working has led to an increase in strength.
FIG. 20 shows, as a comparative example, a nominal stress-nominal strain curve (upper diagram) obtained when pure magnesium was extruded at 328 ° C. and compression-deformed at 400 ° C. for 24 hours, and the nominal strain. FIG. 6 shows a nominal stress-nominal strain curve obtained when the compression test piece taken by machining is compressed and deformed in a direction parallel to and perpendicular to extrusion after the deformation is stopped at 0.14. The yield strength differs greatly, and deformation anisotropy can be confirmed.
FIG. 21 shows, as an example, a nominal stress-nominal strain curve obtained by compressing and deforming a material that was extruded at 300 ° C. and held at 400 ° C. for 24 hours. ), And after stopping the deformation at a nominal strain of 0.40, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in parallel and perpendicular to the extrusion. ing. It can be confirmed that the anisotropy of the yield strength is reduced.
FIG. 22 shows, as an example, a nominal stress-nominal strain curve obtained by compressing and deforming a material which was extruded at 425 ° C. and held at 400 ° C. for 24 hours as an example (upper figure). ), And after stopping the deformation at a nominal strain of 0.40, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in parallel and perpendicular to the extrusion. ing. It can be confirmed that the anisotropy of the yield strength is reduced.
FIG. 23 shows, as an example, a nominal stress-nominal strain curve (upper diagram) obtained when Mg-0.3 atomic% Yb was extruded at 300 ° C. and compression-deformed on a material held at 450 ° C. for 24 hours. ), And after stopping the deformation at a nominal strain of 0.40, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in parallel and perpendicular to the extrusion. ing. It can be confirmed that the anisotropy of yield strength and work hardening rate is reduced.
FIG. 24 shows, as an example, a nominal stress-nominal strain curve obtained by compressing and deforming a material which was extruded at 300 ° C. and held at 450 ° C. for 24 hours as an example (upper diagram). ), And after stopping the deformation at a nominal strain of 0.35, again shows a nominal stress-nominal strain curve obtained when the compression specimen taken by machining is compressed and deformed in the direction parallel and perpendicular to the extrusion. ing. It can be confirmed that the anisotropy of yield strength and work hardening rate is reduced.
FIG. 25 shows a crystal grain structure of a material obtained by extruding Mg-0.6 atomic% Y at an extrusion ratio of 25: 1 and a temperature of 320 ° C. as a comparative example. The black line in the figure indicates an interface having a crystal orientation difference of 5 ° or more as a crystal grain boundary. Compared with the example shown in FIG. 11, it can be seen that in the structure after cold working of the comparative example, a coarse crystal grain structure remains insufficiently divided in the left part and the central part in the figure.
FIG. 26 shows specimens taken in parallel and perpendicular to the extrusion from a material obtained by extruding Mg-0.6 atomic% Y shown as a comparative example in FIG. 25 at an extrusion ratio of 25: 1 and a temperature of 320 ° C. The result of compression test is shown. It can be seen that the nominal strain at break is 0.13 or less, and the cold workability is low while having the same composition as the alloy of this example shown in FIGS. In addition, the work hardening rate after yield varies greatly depending on the sampling direction, and the nominal stress immediately before fracture when subjected to a compression test in the parallel direction of extrusion is nearly twice as large as that in the vertical direction of extrusion. As can be seen, the deformation anisotropy is strong.
本発明によれば、冷間塑性加工に必要な負荷加重を、著しく引き下げることができるＭｇ基合金冷間加工部材が提供され、その実用化が可能になる。 ADVANTAGE OF THE INVENTION According to this invention, the Mg base alloy cold work member which can reduce remarkably the load load required for cold plastic working is provided, and the practical use is attained.
Claims (3)
構成するＭｇ基合金は、Ｙを０．１ａｔ％以上３．０ａｔ％以下、および残部をＭｇと不可避的不純物から成る化学組成を有すると共に、
前記Ｍｇ基合金は、温間で相当塑性ひずみ１以上が加えられ、その後３００〜５５０℃の範囲で等温保持が行われたＭｇ基合金に対して、さらに、冷間加工を公称ひずみにて０．１５以上を付与され、
当該Ｍｇ基合金の組織中に当該冷間加工により分割され微細化された結晶粒が含まれて、その結晶粒径の平均値が３０μｍ以下であると共に、平均直径サイズとして２〜５０ｎｍ、分散間隔が２〜５０ｎｍのＹの高濃度領域を有することを特徴とするＭｇ基合金冷間加工部材。 A Mg-based alloy cold-worked member formed by cold-working a Mg-based alloy into a predetermined shape,
The constituent Mg-based alloy has a chemical composition of Y being 0.1 at% or more and 3.0 at% or less and the balance being Mg and inevitable impurities ,
For the Mg-based alloy, the equivalent plastic strain of 1 or more is applied warm, and then the Mg-based alloy is kept isothermally in the range of 300 to 550 ° C. .15 or more,
Includes the Mg based alloy in the tissue divided by the cold working miniaturized crystal grains, along with their mean value of the grain diameter is 30μm or less, 2 to 50 nm as an average diameter size, dispersion spacing Has a high concentration region of Y of 2 to 50 nm .
構成するＭｇ基合金は、Ｙｂを０．３ａｔ％、残部をＭｇと不可避的不純物から成る化学組成を有し、
前記Ｍｇ基合金は、温間で相当塑性ひずみ１以上が加えられ、その後３００〜５５０℃の範囲で等温保持が行われたＭｇ基合金に対して、さらに、冷間加工を公称ひずみにて０．１５以上を付与され、
当該Ｍｇ基合金の組織中に当該冷間加工により分割され微細化された結晶粒が含まれて、その結晶粒径の平均値が３０μｍ以下であると共に、平均直径サイズとして２〜５０ｎｍ、分散間隔が２〜５０ｎｍのＹｂの高濃度領域を有することを特徴とするＭｇ基合金冷間加工部材。 A Mg-based alloy cold-worked member formed by cold-working a Mg-based alloy into a predetermined shape,
The constituent Mg-based alloy has a chemical composition consisting of 0.3 at% Yb, the balance Mg and inevitable impurities,
For the Mg-based alloy, the equivalent plastic strain of 1 or more is applied warm, and then the Mg-based alloy is kept isothermally in the range of 300 to 550 ° C. .15 or more,
Includes the Mg based alloy in the tissue divided by the cold working miniaturized crystal grains, along with their mean value of the grain diameter is 30μm or less, 2 to 50 nm as an average diameter size, dispersion spacing Has a high concentration region of Yb with a thickness of 2 to 50 nm .
構成するＭｇ基合金は、Ｇｄを０．３ａｔ％、残部をＭｇと不可避的不純物から成る化学組成を有し、
前記Ｍｇ基合金は、温間で相当塑性ひずみ１以上が加えられ、その後３００〜５５０℃の範囲で等温保持が行われたＭｇ基合金に対して、さらに、冷間加工を公称ひずみにて０．１５以上を付与され、
当該Ｍｇ基合金の組織中に当該冷間加工により分割され微細化された結晶粒が含まれて、その結晶粒径の平均値が３０μｍ以下であると共に、平均直径サイズとして２〜５０ｎｍ、分散間隔が２〜５０ｎｍのＧｄの高濃度領域を有することを特徴とするＭｇ基合金冷間加工部材。
A Mg-based alloy cold-worked member formed by cold-working a Mg-based alloy into a predetermined shape,
The constituent Mg-based alloy has a chemical composition consisting of 0.3 at% Gd, the remainder Mg and inevitable impurities,
For the Mg-based alloy, the equivalent plastic strain of 1 or more is applied warm, and then the Mg-based alloy is kept isothermally in the range of 300 to 550 ° C. .15 or more,
Includes the Mg based alloy in the tissue divided by the cold working miniaturized crystal grains, along with their mean value of the grain diameter is 30μm or less, 2 to 50 nm as an average diameter size, dispersion spacing Has a high concentration region of Gd of 2 to 50 nm, a Mg-based alloy cold-worked member.
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| JP2010521754A JP5549981B2 (en) | 2008-07-22 | 2009-07-22 | Mg-based alloy cold worked parts |
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| JP2008188397 | 2008-07-22 | ||
| PCT/JP2009/063452 WO2010010965A1 (en) | 2008-07-22 | 2009-07-22 | COLD-WORKED Mg-BASE ALLOY PRODUCT |
| JP2010521754A JP5549981B2 (en) | 2008-07-22 | 2009-07-22 | Mg-based alloy cold worked parts |
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| EP (1) | EP2319949B1 (en) |
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| US9980470B2 (en) | 2013-03-14 | 2018-05-29 | Erasmus University Medical Center | Antibody production |
Citations (7)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPH0941066A (en) * | 1995-08-01 | 1997-02-10 | Mitsui Mining & Smelting Co Ltd | Magnesium alloy capable of cold press working |
| JP2001294966A (en) * | 2000-04-14 | 2001-10-26 | Matsushita Electric Ind Co Ltd | Magnesium alloy sheet, manufacturing method thereof and molding using the same |
| JP2003328065A (en) * | 2002-05-10 | 2003-11-19 | Toyo Kohan Co Ltd | Wrought magnesium thin-sheet superior in formability, and manufacturing method therefor |
| WO2006003899A1 (en) * | 2004-06-30 | 2006-01-12 | Sumitomo Electric Industries, Ltd. | Method for producing magnesium alloy product |
| JP2006016658A (en) * | 2004-06-30 | 2006-01-19 | National Institute For Materials Science | High strength and high ductility magnesium alloy and its production method |
| WO2007058276A1 (en) * | 2005-11-16 | 2007-05-24 | National Institute For Materials Science | Magnesium-based biodegradable metal material |
| JP2008214668A (en) * | 2007-02-28 | 2008-09-18 | National Institute Of Advanced Industrial & Technology | Magnesium alloy press-formed body and its manufacturing method |
-
2009
- 2009-07-22 CN CN2009801285689A patent/CN102105613A/en active Pending
- 2009-07-22 US US13/055,214 patent/US20110135532A1/en not_active Abandoned
- 2009-07-22 EP EP09800476.5A patent/EP2319949B1/en not_active Not-in-force
- 2009-07-22 JP JP2010521754A patent/JP5549981B2/en not_active Expired - Fee Related
- 2009-07-22 WO PCT/JP2009/063452 patent/WO2010010965A1/en active Application Filing
Patent Citations (7)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPH0941066A (en) * | 1995-08-01 | 1997-02-10 | Mitsui Mining & Smelting Co Ltd | Magnesium alloy capable of cold press working |
| JP2001294966A (en) * | 2000-04-14 | 2001-10-26 | Matsushita Electric Ind Co Ltd | Magnesium alloy sheet, manufacturing method thereof and molding using the same |
| JP2003328065A (en) * | 2002-05-10 | 2003-11-19 | Toyo Kohan Co Ltd | Wrought magnesium thin-sheet superior in formability, and manufacturing method therefor |
| WO2006003899A1 (en) * | 2004-06-30 | 2006-01-12 | Sumitomo Electric Industries, Ltd. | Method for producing magnesium alloy product |
| JP2006016658A (en) * | 2004-06-30 | 2006-01-19 | National Institute For Materials Science | High strength and high ductility magnesium alloy and its production method |
| WO2007058276A1 (en) * | 2005-11-16 | 2007-05-24 | National Institute For Materials Science | Magnesium-based biodegradable metal material |
| JP2008214668A (en) * | 2007-02-28 | 2008-09-18 | National Institute Of Advanced Industrial & Technology | Magnesium alloy press-formed body and its manufacturing method |
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| Publication number | Publication date |
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| JPWO2010010965A1 (en) | 2012-01-05 |
| EP2319949A1 (en) | 2011-05-11 |
| EP2319949A4 (en) | 2012-12-12 |
| EP2319949B1 (en) | 2015-02-18 |
| US20110135532A1 (en) | 2011-06-09 |
| WO2010010965A1 (en) | 2010-01-28 |
| CN102105613A (en) | 2011-06-22 |
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