JP2012162773A - Method for manufacturing grain-oriented electrical steel sheet - Google Patents
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Abstract
Description
本発明は、結晶粒の方位が、ミラー指数で、板面に｛１１０｝、圧延方向に＜００１＞に高度に集積した、いわゆる方向性電磁鋼板の製造方法に関するものである。 The present invention relates to a method for producing a so-called grain-oriented electrical steel sheet in which the orientation of crystal grains is highly integrated with a Miller index, {110} on the plate surface, and <001> in the rolling direction.
方向性電磁鋼板は、軟磁性材料であり、主に変圧器等の電気機器の鉄芯として用いられている。この方向性電磁鋼板は、二次再結晶焼鈍を施して、結晶粒を｛１１０｝＜００１＞方位（以降、「Ｇｏｓｓ方位」と称す。）に高度に集積させたものであり、優れた磁気特性を示すことが知られている（例えば、特許文献１参照。）。なお、電磁鋼板の磁気特性を評価する指標としては、磁場の強さ８００Ａ／ｍにおける磁束密度Ｂ８と、励磁周波数５０Ｈｚの交流磁場で１．７Ｔまで磁化させたときの鋼板１ｋｇあたりの鉄損Ｗ１７／５０が多く用いられている。 The grain-oriented electrical steel sheet is a soft magnetic material and is mainly used as an iron core of electrical equipment such as a transformer. This grain-oriented electrical steel sheet is obtained by performing secondary recrystallization annealing and highly accumulating crystal grains in {110} <001> orientation (hereinafter referred to as “Goss orientation”). It is known to show characteristics (for example, see Patent Document 1). In addition, as an index for evaluating the magnetic properties of the electromagnetic steel sheet, the iron loss per kg of the steel sheet when magnetized to 1.7 T with a magnetic flux density B 8 at a magnetic field strength of 800 A / m and an alternating magnetic field with an excitation frequency of 50 Hz. W 17/50 is often used.
ところで、方向性電磁鋼板の磁気特性は、二次再結晶焼鈍させる前、すなわち、一次再結晶焼鈍後の鋼板の集合組織を制御することで改善されることが知られている。例えば、特許文献２には、一次再結晶焼鈍後の鋼板の表層近傍の集合組織が、Ｂｕｎｇｅのオイラー角表示で、φ１＝０°、Φ＝１５°、φ２＝０°の方位から１０°以内、または、φ１＝５°、Φ＝２０°、φ２＝７０°の方位から１０°以内に極大方位を有し、かつ、鋼板の中心層の集合組織が、同じくＢｕｎｇｅのオイラー角表示で、φ１＝９０°、Φ＝６０°、φ２＝４５°の方位から５°以内に極大方位を有する場合に、二次再結晶焼鈍後に優れた磁気特性を有する方向性電磁鋼板が得られることが記載されている。 By the way, it is known that the magnetic properties of grain-oriented electrical steel sheets are improved by controlling the texture of the steel sheets before secondary recrystallization annealing, that is, after primary recrystallization annealing. For example, in Patent Document 2, the texture in the vicinity of the surface layer of the steel sheet after the primary recrystallization annealing is 10 from the direction of φ 1 = 0 °, φ = 15 °, φ 2 = 0 ° in Bunge's Euler angle display. Or a maximum orientation within 10 ° from the orientation of φ 1 = 5 °, Φ = 20 °, φ 2 = 70 °, and the texture of the central layer of the steel plate is also a Bunge Euler angle. A grain-oriented electrical steel sheet having excellent magnetic properties after secondary recrystallization annealing when it has a maximum orientation within 5 ° from the orientation of φ 1 = 90 °, φ = 60 °, φ 2 = 45 °. It is described that it is obtained.
一次再結晶後の鋼板（以降、「一次再結晶板」ともいう。）の集合組織を改善する方法については、これまでにも多くの提案がなされており、例えば、特許文献３には、従来の一般的な方向性電磁鋼板の製造方法では、最終冷間圧延の圧下率を７０〜９１％の範囲とすることで、安定して優れた磁気特性が得られることが開示されている。そして、この技術は、二次再結晶におけるＧｏｓｓ方位粒の粒成長性に関して、粒界エネルギーの観点から検討を加え、一次再結晶組織におけるＧｏｓｓ方位粒に対する高エネルギー粒界（Ｇｏｓｓ方位粒に対する方位差角が２０〜４５°である粒界）の易動度が重要な働きをすることを明らかにしている。 Many proposals have been made to improve the texture of a steel sheet after primary recrystallization (hereinafter, also referred to as “primary recrystallization board”). For example, Patent Document 3 discloses a conventional technique. In the general method for producing a grain-oriented electrical steel sheet, it is disclosed that excellent magnetic properties can be stably obtained by setting the rolling reduction ratio of the final cold rolling in the range of 70 to 91%. And this technique adds the examination from the viewpoint of grain boundary energy about the grain growth property of the Goss orientation grain in the secondary recrystallization, and the high energy grain boundary (the orientation difference with respect to the Goss orientation grain) in the primary recrystallization structure. It has been clarified that the mobility of the grain boundaries having an angle of 20 to 45 ° plays an important role.
すなわち、二次再結晶は、拡散律速であるインヒビターの析出物の成長によって発現することが知られている。高エネルギー粒界は、粒界内の自由空間が大きく、乱雑な構造をしているため、拡散速度が速いのが特徴である。そのため、高エネルギー粒界上の析出物は、仕上焼鈍中に優先的に粗大化が進行するため、ピン止めが外れて粒界が移動を開始し、Ｇｏｓｓ方位粒が成長する。そして、最終冷間圧延での圧下率を７０〜９１％の範囲に制御した場合には、Ｇｏｓｓ方位粒に対する高エネルギー粒界の頻度が増加し、二次再結晶においては、Ｇｏｓｓ方位粒が安定して成長するとしている。 That is, it is known that secondary recrystallization is manifested by growth of inhibitor precipitates that are diffusion-limited. High energy grain boundaries are characterized by a high diffusion rate due to the large free space in the grain boundaries and a messy structure. For this reason, the precipitate on the high energy grain boundary is preferentially coarsened during finish annealing, so that the pinning is released and the grain boundary starts to move, and Goss orientation grains grow. When the reduction ratio in the final cold rolling is controlled within the range of 70 to 91%, the frequency of high energy grain boundaries with respect to the Goss orientation grains increases, and the Goss orientation grains are stable in the secondary recrystallization. And will grow.
最終冷間圧延の圧下率以外に、一次再結晶板の集合組織に影響を及ぼす因子としては、冷間圧延前のＣの存在形態が挙げられる。例えば、特許文献４には、１回の冷間圧延で方向性電磁鋼板を製造する方法において、冷間圧延前の熱延板焼鈍後の冷却に際して７７０〜４００℃間の滞留時間を６０秒未満、４００〜３００℃間の滞留時間を６０秒未満、３００〜２００℃間の滞留時間を３０秒以上とする制御冷却を施すことで炭化物の析出形態を制御し、磁気特性を改善する技術が開示されている。しかし、この技術は、冷延１回法による製造方法であるため、一次再結晶板の集合組織の改善効果が十分に得られていない。そのため、磁気特性のさらなる向上のためには、中間焼鈍を挟む２回以上の冷間圧延を施す製造方法が望ましいと考えられる。 In addition to the rolling reduction of the final cold rolling, as a factor affecting the texture of the primary recrystallized sheet, the existence form of C before cold rolling can be mentioned. For example, in Patent Document 4, in a method of producing a grain-oriented electrical steel sheet by one cold rolling, the residence time between 770-400 ° C. is less than 60 seconds during cooling after hot-rolled sheet annealing before cold rolling. Discloses a technique for improving the magnetic properties by controlling the precipitation form of carbides by controlling cooling so that the residence time between 400 and 300 ° C. is less than 60 seconds and the residence time between 300 and 200 ° C. is 30 seconds or more. Has been. However, since this technique is a manufacturing method based on a single cold rolling method, the effect of improving the texture of the primary recrystallized plate is not sufficiently obtained. Therefore, in order to further improve the magnetic properties, it is considered that a manufacturing method in which cold rolling is performed twice or more with intermediate annealing interposed therebetween is desirable.
中間焼鈍を挟む複数回の冷間圧延を行う２回冷延法で、最終冷間圧延前のＣの存在形態を制御する他の技術としては、例えば、特許文献５には、最終冷間圧延前の中間焼鈍における冷却を、鋼板表面温度が９００〜２００℃の間を冷却速度２５℃／ｓ以上で急冷して炭化物の析出を防止することで、一次再結晶集合組織を改善し、優れた磁気特性を安定して得る技術が開示されている。また、特許文献６には、最終冷間圧前の中間焼鈍の冷却速度を７００〜１５０℃間において１０℃／ｓ以上として急冷することで、Ｃの析出を抑制して固溶Ｃ量を増加させ、一次再結晶板集合組織を改善する技術が開示されている。しかし、これらの技術は、熱延板焼鈍での冷却過程における規定が十分ではないため、一次再結晶板の集合組織の改善効果が小さい。 As another technique for controlling the existence form of C before the final cold rolling by a two-time cold rolling method in which cold rolling is performed a plurality of times with intermediate annealing, for example, Patent Document 5 discloses final cold rolling. The cooling in the previous intermediate annealing is rapidly cooled at a cooling rate of 25 ° C./s or higher between the steel plate surface temperature of 900 to 200 ° C. to prevent the precipitation of carbides, thereby improving the primary recrystallization texture. A technique for stably obtaining magnetic characteristics is disclosed. Patent Document 6 discloses that the cooling rate of the intermediate annealing before the final cold pressure is quenched at 700 ° C./s to 10 ° C./s or more, thereby suppressing the precipitation of C and increasing the amount of solid solution C. And a technique for improving the primary recrystallized plate texture is disclosed. However, these techniques are not sufficiently defined in the cooling process in the hot-rolled sheet annealing, so the effect of improving the texture of the primary recrystallized sheet is small.
また、上述した従来技術は、いずれも、最終冷間圧延前の中間焼鈍における冷却を制御することで、一次再結晶板の集合組織の改善を図ろうとする技術である。しかし、一次再結晶板の集合組織は、前工程の中間焼鈍後の鋼板（以降、「中間焼鈍板」ともいう。）の集合組織の影響を大きく受け、さらに、中間焼鈍板の集合組織は、その前の熱延板焼鈍の影響を大きく受けるにも拘らず、従来技術に於いては、熱延板焼鈍の影響については十分な検討がなされていない。特に、熱延板焼鈍における冷却過程は、冷間圧延前のＣの存在形態を制御する上で重要な工程であり、一次再結晶板の集合組織にも大きな影響を及ぼすため、厳密な制御が行われてしかるべきである。 In addition, any of the conventional techniques described above is a technique for improving the texture of the primary recrystallized plate by controlling the cooling in the intermediate annealing before the final cold rolling. However, the texture of the primary recrystallized plate is greatly influenced by the texture of the steel plate after the intermediate annealing in the previous process (hereinafter also referred to as “intermediate annealed plate”), and the texture of the intermediate annealed plate is In spite of being greatly affected by the previous hot-rolled sheet annealing, the prior art has not sufficiently studied the effects of hot-rolled sheet annealing. In particular, the cooling process in hot-rolled sheet annealing is an important process for controlling the existence form of C before cold rolling, and has a great influence on the texture of the primary recrystallized sheet. Should be done.
なお、熱延板焼鈍における冷却を制御する技術としては、幾つかの提案があり、例えば、特許文献６には、熱延板焼鈍後の冷却速度を７００〜１５０℃間において１０℃／ｓ以上とすることが、また特許文献７には、熱延板焼鈍後、室温までの冷却速度を１０℃／ｓ以上で急冷することが開示されている。 In addition, as a technique for controlling cooling in hot-rolled sheet annealing, there are several proposals. For example, in Patent Document 6, the cooling rate after hot-rolled sheet annealing is 10 ° C./s or more between 700 to 150 ° C. In addition, Patent Document 7 discloses that after hot-rolled sheet annealing, rapid cooling is performed at a cooling rate to room temperature of 10 ° C./s or more.
しかしながら、特許文献６や７の従来技術では、熱延板焼鈍での冷却速度を１０℃／ｓ以上とすることだけしか規定されておらず、斯かる規定だけでは、Ｃの存在形態を制御する、具体的には、冷間圧延前の炭化物（カーバイド）を均一微細に析出させるには十分ではなく、その結果、一次再結晶後の鋼板の集合組織を改善する効果が小さいのが実情である。 However, in the prior arts of Patent Documents 6 and 7, only the cooling rate in hot-rolled sheet annealing is set to 10 ° C./s or more, and the existence form of C is controlled only by such specifications. Specifically, it is not sufficient to precipitate carbide (carbide) before cold rolling uniformly and finely, and as a result, the effect of improving the texture of the steel sheet after primary recrystallization is small. .
そこで、本発明の目的は、冷間圧延前のＣの存在形態を最適化することで、その後の中間焼鈍板の集合組織および一次再結晶板の集合組織を改善し、ひいては二次再結晶焼鈍後に優れた磁気特性を安定して得ることができる方向性電磁鋼板の有利な製造方法を提案することにある。 Therefore, an object of the present invention is to improve the texture of the intermediate annealed sheet and the texture of the primary recrystallized sheet by optimizing the existence form of C before cold rolling, and thus secondary recrystallized annealing. It is to propose an advantageous manufacturing method of grain-oriented electrical steel sheet that can stably obtain excellent magnetic properties later.
発明者らは、上述した課題を解決するべく、熱延圧延後から冷間圧延前までの工程条件に着目して鋭意検討を重ねた。その結果、熱間圧延した鋼板表面に生成するスケールと熱延板焼鈍での雰囲気ガスとを反応させて鋼板表層のＣ濃度を低下させ、さらに、熱延板焼鈍後の冷却過程を厳密に制御し、熱延板焼鈍後の鋼板中心部に炭化物（カーバイド）を均一微細に析出させることで、鋼板表層部と中心部それぞれの中間焼鈍板の集合組織とその後の一次再結晶板の集合組織を最適化することができ、ひいては、二次再結晶焼鈍後に優れた磁気特性を有する方向性電磁鋼板を得ることができること、さらに、熱延板焼鈍後の鋼板表面にショットブラスト加工を施して鋼板表面近傍に歪を付与した場合には、より上記効果が高まることを見出し、本発明を完成するに至った。 In order to solve the above-described problems, the inventors made extensive studies by paying attention to process conditions from after hot rolling to before cold rolling. As a result, the scale generated on the surface of the hot-rolled steel sheet reacts with the atmospheric gas in the hot-rolled sheet annealing to lower the C concentration of the steel sheet surface layer, and the cooling process after hot-rolled sheet annealing is strictly controlled. Then, by uniformly and finely depositing carbide (carbide) in the central part of the steel sheet after the hot-rolled sheet annealing, the texture of the intermediate annealed sheet in the steel sheet surface layer part and the central part and the texture of the primary recrystallized sheet thereafter It is possible to optimize and, in turn, to obtain a grain-oriented electrical steel sheet having excellent magnetic properties after secondary recrystallization annealing, and further, subjecting the steel sheet surface after hot-rolled sheet annealing to shot blast processing The inventors have found that the above effect is further enhanced when strain is applied in the vicinity, and have completed the present invention.
すなわち、本発明は、Ｃ：０．００５〜０．１５ｍａｓｓ％、Ｓｉ：２．５〜７．０ｍａｓｓ％、Ｍｎ：０．００５〜０．３ｍａｓｓ％、ｓｏｌ．Ａｌ：０．０１〜０．０５ｍａｓｓ％、Ｎ：０．００２〜０．０１２ｍａｓｓ％、ＳおよびＳｅのうちの１種または２種を合計で０．００５〜０．０５ｍａｓｓ％を含有し、残部がＦｅおよび不可避的不純物からなる成分組成を有する鋼素材を再加熱した後、熱間圧延し、熱延板焼鈍し、中間焼鈍を挟む２回以上の冷間圧延して最終板厚の冷延板とし、その後、一次再結晶焼鈍し、二次再結晶焼鈍する一連の工程からなる方向性電磁鋼板の製造方法において、熱間圧延後の巻取温度を７６０〜４６０℃の範囲として鋼板表面にＦｅ３Ｏ４を主体とする酸化物を生成させた後、酸化性雰囲気中で、８００℃以上に加熱後、８００℃から３５０〜２００℃間の冷却停止温度までを平均冷却速度１０〜１００℃／ｓで冷却し、その後、冷却停止温度から冷却速度５℃／ｓ以下で４０〜２００ｓ間徐冷する熱延板焼鈍を施すことを特徴とする方向性電磁鋼板の製造方法である。 That is, the present invention relates to C: 0.005-0.15 mass%, Si: 2.5-7.0 mass%, Mn: 0.005-0.3 mass%, sol. Al: 0.01-0.05 mass%, N: 0.002-0.012 mass%, one or two of S and Se in total contain 0.005-0.05 mass%, with the balance being After reheating a steel material having a component composition composed of Fe and inevitable impurities, hot rolling, hot rolled sheet annealing, cold rolling at least twice with intermediate annealing in between, and then cold rolling a final sheet thickness Then, in the method for producing a grain-oriented electrical steel sheet comprising a series of steps of primary recrystallization annealing and secondary recrystallization annealing, the coiling temperature after hot rolling is set to a range of 760 to 460 ° C., and Fe on the steel sheet surface. After generating an oxide mainly composed of 3 O 4 , after heating to 800 ° C. or higher in an oxidizing atmosphere, an average cooling rate of 10 to 100 ° C./800° C. to a cooling stop temperature of 350 to 200 ° C. s, then stop cooling temperature It is a manufacturing method of a grain-oriented electrical steel sheet characterized by subjecting a hot-rolled sheet annealing to slow cold 40~200s below the cooling rate of 5 ° C. / s from.
本発明の方向性電磁鋼板の製造方法は、上記熱延板焼鈍した後、鋼板表面にショットブラスト加工を施すことを特徴とする。 The method for producing a grain-oriented electrical steel sheet according to the present invention is characterized by performing shot blasting on the steel sheet surface after the hot-rolled sheet annealing.
また、本発明の方向性電磁鋼板の製造方法における上記鋼素材は、上記成分組成に加えてさらに、Ｎｉ：０．００５〜１．５ｍａｓｓ％、Ｓｎ：０．００５〜０．５０ｍａｓｓ％、Ｓｂ：０．００５〜０．５０ｍａｓｓ％、Ｍｏ：０．００５〜０．１ｍａｓｓ％、Ｃｕ：０．００５〜１．５ｍａｓｓ％、Ｃｒ：０．０３〜１．５０ｍａｓｓ％およびＰ：０．００５〜０．５０ｍａｓｓ％の内から選ばれる１種または２種以上を含有することを特徴とする。 Moreover, in addition to the said component composition, the said steel raw material in the manufacturing method of the grain-oriented electrical steel sheet of this invention is further Ni: 0.005-1.5mass%, Sn: 0.005-0.50mass%, Sb: 0.005-0.50 mass%, Mo: 0.005-0.1 mass%, Cu: 0.005-1.5 mass%, Cr: 0.03-1.50 mass%, and P: 0.005-0. 1 type or 2 types or more chosen from 50 mass% are contained.
本発明によれば、熱延板焼鈍条件を適正化することで、一次再結晶板の集合組織を改善し、二次再結晶焼鈍後に従来にも増して優れた磁気特性を有する方向性電磁鋼板を、安定して製造することができる。その結果、良好な磁気特性を得ることが難しい板厚０．２３ｍｍの方向性電磁鋼板でも、磁束密度Ｂ８が１．９３Ｔ以上で、かつ鉄損Ｗ１７／５０が０．９４Ｗ／ｋｇ以下の磁気特性が安定して得られ、さらには、鉄損Ｗ１７／５０が０．９０Ｗ／ｋｇ以下の磁気特性を達成することが可能となる。 According to the present invention, by optimizing the hot-rolled sheet annealing conditions, the texture of the primary recrystallized sheet is improved, and the grain-oriented electrical steel sheet having superior magnetic properties after the secondary recrystallized annealing as compared with the prior art. Can be manufactured stably. As a result, even with a grain-oriented electrical steel sheet having a thickness of 0.23 mm, it is difficult to obtain good magnetic properties, the magnetic flux density B 8 is 1.93 T or more and the iron loss W 17/50 is 0.94 W / kg or less. Magnetic characteristics can be obtained stably, and furthermore, magnetic characteristics with an iron loss W 17/50 of 0.90 W / kg or less can be achieved.
まず、本発明の方向性電磁鋼板の製造に用いる鋼素材は、その成分組成がＣ：０．００５〜０．１５ｍａｓｓ％、Ｓｉ：２．５〜７．０ｍａｓｓ％、Ｍｎ：０．００５〜０．３ｍａｓｓ％、ｓｏｌ．Ａｌ：０．０１〜０．０５ｍａｓｓ％、Ｎ：０．００２〜０．０１２ｍａｓｓ％、ＳおよびＳｅのうちの１種または２種を合計で０．００５〜０．０５ｍａｓｓ％を含有するものであることが必要である。以下、その限定理由について説明する。 First, as for the steel raw material used for manufacture of the grain-oriented electrical steel sheet of this invention, the component composition is C: 0.005-0.15mass%, Si: 2.5-7.0mass%, Mn: 0.005-0 .3 mass%, sol. Al: 0.01 to 0.05 mass%, N: 0.002 to 0.012 mass%, one or two of S and Se in total contain 0.005 to 0.05 mass% It is necessary. Hereinafter, the reason for limitation will be described.
Ｃ：０．００５〜０．１５ｍａｓｓ％
Ｃは、熱延および熱延板焼鈍の均熱時におけるγ−α変態を利用して、熱延板組織の改善を図るのに必要な元素である。しかし、０．００５ｍａｓｓ％未満では、熱延板組織の改善効果が小さく、｛１１１｝＜１１２＞と｛１１０｝＜００１＞のバランスがとれた一次再結晶集合組織を得ることが難しくなる。一方、０．１５ｍａｓｓ％を超えると、脱炭焼鈍での負荷が増大して脱炭が不完全となり、製品板において磁気時効を起こす原因となる。よって、Ｃは０．００５〜０．１５ｍａｓｓ％の範囲とする。好ましくは、０．０２〜０．１０ｍａｓｓ％の範囲である。
C: 0.005-0.15 mass%
C is an element necessary for improving the hot-rolled sheet structure by utilizing the γ-α transformation during soaking of hot-rolled and hot-rolled sheet annealing. However, if it is less than 0.005 mass%, the effect of improving the hot-rolled sheet structure is small, and it becomes difficult to obtain a primary recrystallized texture that balances {111} <112> and {110} <001>. On the other hand, when it exceeds 0.15 mass%, the load in decarburization annealing increases, decarburization becomes incomplete, and causes a magnetic aging in a product plate. Therefore, C is in the range of 0.005 to 0.15 mass%. Preferably, it is the range of 0.02-0.10 mass%.
Ｓｉ：２．５〜７．０ｍａｓｓ％
Ｓｉは、鋼の電気抵抗を増大させ、鉄損の一部を構成する渦電流損を低減するのに有効な元素であり、上記効果を得るため、本発明では、２．５ｍａｓｓ％以上添加する。また、Ｓｉが２．５ｍａｓｓ％未満では、α−γ変態の存在によって、最終仕上焼鈍における二次再結晶が阻害されて、磁気特性が低下するという問題もある。一方、Ｓｉ添加による上記鉄損低減効果は、１１ｍａｓｓ％まで得られるが、７．０ｍａｓｓ％を超えて添加すると、加工性が著しく低下し、製造することが難しくなる。よって、Ｓｉは２．５〜７．０ｍａｓｓ％の範囲とする。好ましくは、３．０〜６．５ｍａｓｓ％の範囲である。
Si: 2.5-7.0 mass%
Si is an element effective for increasing the electrical resistance of steel and reducing eddy current loss constituting a part of iron loss. In order to obtain the above effect, Si is added in an amount of 2.5 mass% or more. . Further, if Si is less than 2.5 mass%, the presence of the α-γ transformation hinders secondary recrystallization in the final finish annealing, resulting in a problem that the magnetic properties are deteriorated. On the other hand, the iron loss reduction effect due to the addition of Si can be obtained up to 11 mass%, but if added over 7.0 mass%, the workability is remarkably deteriorated and it is difficult to manufacture. Therefore, Si is set to a range of 2.5 to 7.0 mass%. Preferably, it is in the range of 3.0 to 6.5 mass%.
Ｍｎ：０．００５〜０．３ｍａｓｓ％
Ｍｎは、二次再結晶焼鈍での昇温過程において、正常粒成長を抑制するインヒビターの働きをするＭｎＳおよびＭｎＳｅを形成する、本発明においては重要な元素である。しかし、Ｍｎ含有量が０．００５ｍａｓｓ％未満では、必要なインヒビターの絶対量が不足するため、十分な抑制力が得られない。一方、０．３ｍａｓｓ％を超える添加は、インヒビターを完全固溶させるための熱延前のスラブ加熱温度を高温にする必要があったり、インヒビターが粗大析出して抑制力が不十分となったりする。よって、Ｍｎは０．００５〜０．３ｍａｓｓ％の範囲とする。好ましくは、０．０２〜０．１ｍａｓｓ％の範囲である。
Mn: 0.005 to 0.3 mass%
Mn is an important element in the present invention that forms MnS and MnSe that act as an inhibitor that suppresses the growth of normal grains in the temperature rising process during secondary recrystallization annealing. However, if the Mn content is less than 0.005 mass%, a sufficient inhibitory force cannot be obtained because the absolute amount of the necessary inhibitor is insufficient. On the other hand, when the addition exceeds 0.3 mass%, it is necessary to increase the slab heating temperature before hot rolling to completely dissolve the inhibitor, or the inhibitor is coarsely precipitated and the inhibitory power becomes insufficient. . Therefore, Mn is in the range of 0.005 to 0.3 mass%. Preferably, it is the range of 0.02-0.1 mass%.
ｓｏｌ．Ａｌ：０．０１〜０．０５ｍａｓｓ％
Ａｌは、二次再結晶焼鈍での昇温過程において、正常粒成長を抑制するインヒビターの働きをするＡｌＮを構成する、本発明においては重要な元素である。しかし、Ａｌの含有量がｓｏｌ．Ａｌ（酸可溶性Ａｌ）で０．０１ｍａｓｓ％未満では、インヒビターの絶対量が不足し、抑制力が不十分となる。一方、０．０５ｍａｓｓ％を超えると、ＡｌＮが粗大析出し、やはり抑制力が不十分となる。よって、ｓｏｌ．Ａｌは０．０１〜０．０５ｍａｓｓ％の範囲とする。好ましくは、０．０１５〜０．０４ｍａｓｓ％の範囲である。
sol. Al: 0.01-0.05 mass%
Al is an important element in the present invention that constitutes AlN that acts as an inhibitor that suppresses the growth of normal grains in the temperature rising process during secondary recrystallization annealing. However, the content of Al is sol. If it is less than 0.01 mass% with Al (acid-soluble Al), the absolute amount of the inhibitor is insufficient and the inhibitory power becomes insufficient. On the other hand, when it exceeds 0.05 mass%, AlN coarsely precipitates, and the suppression force is still insufficient. Therefore, sol. Al is in the range of 0.01 to 0.05 mass%. Preferably, it is the range of 0.015-0.04 mass%.
Ｎ：０．００２〜０．０１２ｍａｓｓ％
Ｎは、Ａｌと結合してインヒビターを形成する元素である。しかし、含有量が０．００２ｍａｓｓ％未満では、インヒビターの絶対量が不足するため、抑制力が不十分となる。一方、０．０１２ｍａｓｓ％を超えると、冷間圧延時にブリスターと呼ばれる空孔欠陥を生じるようになる。よって、Ｎは０．００２〜０．０１２ｍａｓｓ％の範囲とする。好ましくは、０．００６〜０．０１０ｍａｓｓ％の範囲である。
N: 0.002-0.012 mass%
N is an element that combines with Al to form an inhibitor. However, if the content is less than 0.002 mass%, the inhibitory power is insufficient because the absolute amount of the inhibitor is insufficient. On the other hand, if it exceeds 0.012 mass%, void defects called blisters are produced during cold rolling. Therefore, N is set to a range of 0.002 to 0.012 mass%. Preferably, it is the range of 0.006-0.010 mass%.
Ｓ，Ｓｅ：１種または２種を合計で０．００５〜０．０５ｍａｓｓ％
ＳおよびＳｅは、Ｍｎと結合してインヒビターを形成するのに必要な元素であり、十分なインヒビター量を確保するため、合計で０．００５ｍａｓｓ％以上添加する。しかし、合計含有量が０．０５ｍａｓｓ％を超えると、仕上焼鈍（純化焼鈍）における脱Ｓ、脱Ｓｅが不完全となり、鉄損特性の低下を引き起こす。よって、本発明では、ＳおよびＳｅは、合計で０．０５ｍａｓｓ％以下添加する。好ましくは、０．０１０〜０．０３５ｍａｓｓ％の範囲である。
S, Se: 1 type or 2 types in total 0.005-0.05 mass%
S and Se are elements necessary for binding to Mn to form an inhibitor. In order to secure a sufficient amount of inhibitor, a total of 0.005 mass% or more is added. However, if the total content exceeds 0.05 mass%, the removal of S and the removal of Se in finish annealing (purification annealing) become incomplete, causing a reduction in iron loss characteristics. Therefore, in the present invention, S and Se are added in a total amount of 0.05 mass% or less. Preferably, it is the range of 0.010-0.035 mass%.
本発明の方向性電磁鋼板は、上記の必須成分に加えてさらに、Ｎｉ，Ｓｎ，Ｓｂ，Ｍｏ，Ｃｕ，ＣｒおよびＰから選ばれる１種または２種以上を下記の範囲で添加することができる。
Ｎｉ：０．００５〜１．５ｍａｓｓ％
Ｎｉは、オーステナイト生成元素であるため、γ−α変態を利用して熱延板組織を改善し、磁気特性を向上するのに有効な元素である。しかし、含有量が０．００５ｍａｓｓ％未満では、上記磁気特性の改善効果が小さく、一方、１．５ｍａｓｓ％を超えると、加工性が低下して製造性が悪化したり、二次再結晶が不安定となって磁気特性が低下したりする。よって、Ｎｉを添加する場合は、０．００５〜１．５ｍａｓｓ％の範囲とするのが好ましい。
In addition to the above essential components, the grain-oriented electrical steel sheet of the present invention can further contain one or more selected from Ni, Sn, Sb, Mo, Cu, Cr and P in the following range. .
Ni: 0.005-1.5 mass%
Since Ni is an austenite-generating element, it is an effective element for improving the hot rolled sheet structure by utilizing the γ-α transformation and improving the magnetic properties. However, if the content is less than 0.005 mass%, the effect of improving the magnetic properties is small. On the other hand, if the content exceeds 1.5 mass%, the workability is deteriorated and the productivity is deteriorated, or secondary recrystallization is not achieved. It becomes stable and the magnetic properties are lowered. Therefore, when adding Ni, it is preferable to set it as the range of 0.005-1.5 mass%.
Ｓｎ：０．００５〜０．５０ｍａｓｓ％、Ｓｂ：０．００５〜０．５０ｍａｓｓ％、Ｍｏ：０．００５〜０．１ｍａｓｓ％、Ｃｕ：０．００５〜１．５ｍａｓｓ％、Ｃｒ：０．０３〜１．５０ｍａｓｓ％およびＰ：０．００５〜０．５０ｍａｓｓ％の内から選ばれる１種または２種以上
Ｓｎ，Ｓｂ，Ｍｏ，Ｃｕ，ＣｒおよびＰは、磁気特性の向上に有用な元素である。しかし、いずれの元素も、含有量が上記下限値未満であると、磁気特性改善効果が小さく、一方、含有量が上記上限値を超えると、二次再結晶が不安定になって磁気特性が低下するようになる。よって、Ｓｎ，Ｓｂ，Ｍｏ，Ｃｕ，ＣｒおよびＰは、Ｓｎ：０．００５〜０．５０ｍａｓｓ％、Ｓｂ：０．００５〜０．５０ｍａｓｓ％、Ｍｏ：０．００５〜０．１ｍａｓｓ％、Ｃｕ：０．００５〜１．５ｍａｓｓ％、Ｃｒ：０．０３〜１．５０ｍａｓｓ％およびＰ：０．００５〜０．５０ｍａｓｓ％の範囲で添加することができる。
Sn: 0.005-0.50 mass%, Sb: 0.005-0.50 mass%, Mo: 0.005-0.1 mass%, Cu: 0.005-1.5 mass%, Cr: 0.03- One or more selected from 1.50 mass% and P: 0.005 to 0.50 mass% Sn, Sb, Mo, Cu, Cr, and P are elements useful for improving magnetic properties. However, if the content of any element is less than the above lower limit value, the effect of improving the magnetic properties is small. On the other hand, if the content exceeds the upper limit value, secondary recrystallization becomes unstable and the magnetic properties are reduced. It begins to decline. Therefore, Sn, Sb, Mo, Cu, Cr and P are Sn: 0.005-0.50 mass%, Sb: 0.005-0.50 mass%, Mo: 0.005-0.1 mass%, Cu: It can be added in the range of 0.005 to 1.5 mass%, Cr: 0.03 to 1.50 mass%, and P: 0.005 to 0.50 mass%.
本発明の方向性電磁鋼板は、上記の必須成分および任意添加成分以外の残部は、Ｆｅおよび不可避的不純物である。ただし、本発明の効果を害しない範囲であれば、その他の元素の添加を拒むものではない。 In the grain-oriented electrical steel sheet of the present invention, the balance other than the above essential components and optional additive components is Fe and inevitable impurities. However, addition of other elements is not rejected as long as the effects of the present invention are not impaired.
次に、本発明の方向性電磁鋼板の製造方法について説明する。
本発明の方向性電磁鋼板の製造方法は、まず、上記成分組成を有する鋼素材（スラブ）を製造し、その後、その鋼素材を、再加熱し、熱間圧延して熱延板とする。
上記スラブ加熱および熱間圧延は、常法で行えばよいが、熱間圧延後のコイルの巻取温度は７６０〜４５０℃の温度範囲とし、鋼板表面にＦｅ３Ｏ４を主体とする酸化皮膜を生成させる必要がある。そのためには、上記温度範囲でコイルに巻取り後、酸素濃度１０ｖｏｌ％以上の雰囲気で少なくとも５ｈｒはコイル状態のまま保持することが好ましい。
Next, the manufacturing method of the grain-oriented electrical steel sheet of this invention is demonstrated.
In the method for producing a grain-oriented electrical steel sheet according to the present invention, first, a steel material (slab) having the above component composition is produced, and then the steel material is reheated and hot-rolled to obtain a hot-rolled sheet.
The slab heating and hot rolling may be performed by a conventional method, but the coil winding temperature after hot rolling is set to a temperature range of 760 to 450 ° C., and the oxide film mainly composed of Fe 3 O 4 on the steel plate surface. Must be generated. For this purpose, it is preferable to keep the coil in the coil state in an atmosphere having an oxygen concentration of 10 vol% or more after winding the coil in the above temperature range.
その後、上記熱延板の表層にＦｅ３Ｏ４を生成させた状態で、酸化性ガス雰囲気中にて熱延板焼鈍を施して、鋼板表層のＣ量を若干低下させることで、熱延板の鋼板組織の改善を行う必要がある。すなわち、熱延板の鋼板表面に生成したスケールと熱延板焼鈍での雰囲気ガスとを反応させて鋼板表層のＣ濃度を低下させることで鋼板表層部に脱炭層を形成し、さらに、その後に冷却を制御することで、熱延板焼鈍後の鋼板中心部に炭化物（カーバイド）を均一微細に析出させて、鋼板表層部と中心部それぞれの中間焼鈍板の集合組織とその後の一次再結晶板の集合組織を最適化することができ、ひいては、二次再結晶焼鈍後に優れた磁気特性を有する方向性電磁鋼板を得ることができる。 Then, in a state where Fe 3 O 4 is generated on the surface layer of the hot-rolled sheet, hot-rolled sheet annealing is performed in an oxidizing gas atmosphere to slightly reduce the amount of C in the surface layer of the steel sheet. It is necessary to improve the steel sheet structure. That is, a decarburized layer is formed on the surface of the steel sheet by reducing the C concentration of the steel sheet surface layer by reacting the scale gas generated on the surface of the steel sheet of the hot rolled sheet with the atmospheric gas in the hot-rolled sheet annealing. By controlling the cooling, carbide (carbide) is uniformly and finely precipitated at the center of the steel sheet after hot-rolled sheet annealing, and the texture of the intermediate annealed sheet at each of the steel sheet surface layer and the center and the subsequent primary recrystallized sheet Thus, a grain-oriented electrical steel sheet having excellent magnetic properties after secondary recrystallization annealing can be obtained.
そのためには、上記熱延板焼鈍は、８００℃以上に加熱・均熱する必要があり、具体的には、均熱温度が８００〜１２００℃で、均熱時間が２〜３００ｓの範囲として行うのが好ましい。均熱温度が８００℃未満あるいは均熱時間が２ｓ未満では、未再結晶部が残存するため、熱延板組織の改善が完全ではない。一方、均熱温度が１２００℃より高温あるいは均熱時間が３００ｓを超えると、ＡｌＮ、ＭｎＳｅおよびＭｎＳの溶解が進行して、二次再結晶過程でインヒビターの抑制力が不足し、二次再結晶し難くなるため、磁気特性が劣化するようになるからである。また、熱延板焼鈍は、表層脱炭を行うため、酸化性ガスの雰囲気下で行うことが必要であり、例えば、燃焼ガス雰囲気下で行うのが好ましい。 For that purpose, the hot-rolled sheet annealing needs to be heated and soaked to 800 ° C. or higher. Specifically, the soaking temperature is 800 to 1200 ° C. and the soaking time is 2 to 300 s. Is preferred. If the soaking temperature is less than 800 ° C. or the soaking time is less than 2 s, an unrecrystallized portion remains, so that the improvement of the hot rolled sheet structure is not complete. On the other hand, if the soaking temperature is higher than 1200 ° C. or the soaking time exceeds 300 s, the dissolution of AlN, MnSe and MnS proceeds, and the inhibitor repressing power is insufficient in the secondary recrystallization process, and the secondary recrystallization. This is because the magnetic characteristics are deteriorated. Moreover, in order to perform surface layer decarburization, hot-rolled sheet annealing needs to be performed in the atmosphere of oxidizing gas, for example, it is preferable to perform in a combustion gas atmosphere.
上記条件で熱延板焼鈍した鋼板は、その後、図１に示す条件で冷却する、すなわち、焼鈍温度から３５０〜２００℃の間の冷却停止温度Ｘまでを平均冷却速度１０〜１００℃／ｓで急速冷却し、その後、冷却停止温度から冷却速度５℃／ｓ以下で４０〜２００ｓ間徐冷することが必要である。
熱延板焼鈍温度から２００〜３５０℃の間の冷却停止温度Ｘまでの冷却速度が１０℃／ｓ未満では、カーバイドがオストワルド成長して粗大化するため、中間焼鈍板やその後の一次再結晶板の集合組織を改善する効果が弱まる結果、二次再結晶後の磁気特性が低下する。一方、上記範囲の平均冷却速度が１００℃／ｓを超えると、固溶Ｃ量が増加して、極微細なカーバイドの均一分散析出が不可能となったり、硬質のマルテンサイト相が生成したりするため、中間焼鈍板やその後の一次再結晶板の集合組織の改善効果が弱まって、やはり、磁気特性の低下を引き起こすからである。
The steel sheet hot-rolled sheet annealed under the above conditions is then cooled under the conditions shown in FIG. 1, that is, from the annealing temperature to the cooling stop temperature X between 350 to 200 ° C. at an average cooling rate of 10 to 100 ° C./s. It is necessary to rapidly cool and then gradually cool from the cooling stop temperature at a cooling rate of 5 ° C./s or less for 40 to 200 s.
When the cooling rate from the hot-rolled sheet annealing temperature to the cooling stop temperature X between 200-350 ° C. is less than 10 ° C./s, the carbide grows and grows Ostwald, so the intermediate annealed plate or the subsequent primary recrystallized plate As a result of the weakening of the effect of improving the texture, the magnetic properties after secondary recrystallization deteriorate. On the other hand, when the average cooling rate in the above range exceeds 100 ° C./s, the amount of dissolved C increases, and uniform dispersion precipitation of ultrafine carbide becomes impossible, or a hard martensite phase is generated. Therefore, the effect of improving the texture of the intermediate annealed plate and the subsequent primary recrystallized plate is weakened, which also causes a decrease in magnetic properties.
冷却停止温度Ｘまで冷却した後の冷却条件について、以下の実験を基に説明する。
Ｃ：０．０５ｍａｓｓ％、Ｓｉ：３．２ｍａｓｓ％、Ｍｎ：０．１ｍａｓｓ％、ｓｏｌ．Ａｌ：０．０２ｍａｓｓ％、Ｎ：０．００７ｍａｓｓ％、Ｓ：０．００３０ｍａｓｓ％およびＳｅ：０．０３ｍａｓｓ％を含有するスラブを、１４００℃に加熱した後、熱間圧延して板厚２．２ｍｍの熱延板とし、６５０℃でコイルに巻き取った後、酸素濃度２０ｖｏｌ％の雰囲気中で２０ｈｒ放置し、その後、１０２５℃×４０ｓの焼鈍後、８００℃から平均冷却速度５０℃／ｓで室温〜５００℃の間の種々の冷却停止温度まで急冷し、その急冷停止温度から３℃／ｓで２０〜４００ｓの間の種々の時間徐冷する熱延板焼鈍を施した。次いで、一次冷間圧延して中間板厚１．５ｍｍの冷延板とし、１０００℃×８０ｓの中間焼鈍を施し、さらに、二次冷間圧延して板厚０．２３ｍｍの冷延板とし、その後、８００℃×１２０ｓの脱炭を兼ねた一次再結晶焼鈍を施した後、鋼板表面にＭｇＯを主成分とする焼鈍分離剤を塗布してから、１１５０℃×５０ｈｒの純化を兼ねた二次再結晶焼鈍を施して方向性電磁鋼板とした。
The cooling conditions after cooling to the cooling stop temperature X will be described based on the following experiment.
C: 0.05 mass%, Si: 3.2 mass%, Mn: 0.1 mass%, sol. A slab containing Al: 0.02 mass%, N: 0.007 mass%, S: 0.0030 mass%, and Se: 0.03 mass% was heated to 1400 ° C. and then hot-rolled to a thickness of 2.2 mm. After being wound on a coil at 650 ° C. and left in an atmosphere with an oxygen concentration of 20 vol% for 20 hours, after annealing at 1025 ° C. × 40 s, the room temperature is increased from 800 ° C. to an average cooling rate of 50 ° C./s. The steel sheet was rapidly cooled to various cooling stop temperatures between ˜500 ° C., and subjected to hot rolled sheet annealing that was gradually cooled from the rapid cooling stop temperature at 3 ° C./s for 20 to 400 s for various times. Next, primary cold-rolled to obtain a cold-rolled sheet having an intermediate thickness of 1.5 mm, subjected to intermediate annealing at 1000 ° C. × 80 s, and further subjected to secondary cold-rolling to obtain a cold-rolled sheet having a thickness of 0.23 mm, Then, after performing primary recrystallization annealing also serving as decarburization at 800 ° C. × 120 s, and then applying an annealing separator mainly composed of MgO to the steel sheet surface, the secondary serving as purification at 1150 ° C. × 50 hr. Reoriented annealing was performed to obtain a grain-oriented electrical steel sheet.
この方向性電磁鋼板から試験片を採取し、鉄損Ｗ１７／５０を測定し、その結果を、熱延板焼鈍の冷却過程における急冷停止温度Ｘとその後の徐冷時間との関係として図２に示した。この図から、熱延板焼鈍後の冷却過程において、急冷停止温度Ｘを２００〜３５０℃の範囲とし、かつ、その後の徐冷時間を４０〜２００ｓの範囲とすることによって、良好な鉄損特性が得られることがわかる。 A test piece was taken from this grain- oriented electrical steel sheet, and the iron loss W 17/50 was measured. The result is shown in FIG. 2 as the relationship between the quenching stop temperature X in the cooling process of hot-rolled sheet annealing and the subsequent slow cooling time. It was shown to. From this figure, in the cooling process after hot-rolled sheet annealing, by setting the rapid cooling stop temperature X in the range of 200 to 350 ° C. and the subsequent slow cooling time in the range of 40 to 200 s, good iron loss characteristics are obtained. It can be seen that
上記条件で鉄損が改善される理由は、以下のように考えている。
一般に、熱延焼鈍板は、中間焼鈍板と比較して、転位密度が高いため、カーバイドの核生成も転位上で頻繁に起こる。したがって、熱延板焼鈍での冷却過程において、上記のように比較的低温で徐冷を行うことで、微細カーバイドの均一析出が促進され、その結果、その後の中間焼鈍板の集合組織、さらには一次再結晶板の集合組織が改善されて、二次再結晶後の磁気特性が改善させるものと考えられる。なお、徐冷時間が４０ｓ未満では、均一微細なカーバイドの析出が十分ではなく、一方、２００ｓを超えるとカーバイトが粗大化し、却って鉄損特性が低下するようになる。
The reason why the iron loss is improved under the above conditions is considered as follows.
In general, since a hot-rolled annealed plate has a higher dislocation density than an intermediate annealed plate, carbide nucleation frequently occurs on the dislocation. Therefore, in the cooling process in hot-rolled sheet annealing, by performing slow cooling at a relatively low temperature as described above, uniform precipitation of fine carbide is promoted, and as a result, the texture of the subsequent intermediate annealing plate, It is considered that the texture of the primary recrystallization plate is improved and the magnetic properties after the secondary recrystallization are improved. When the slow cooling time is less than 40 s, uniform fine carbide is not sufficiently precipitated. On the other hand, when it exceeds 200 s, the carbide is coarsened and the iron loss characteristics are lowered.
さらに、本発明では、上記熱延板焼鈍を施した後の鋼板表面に、ショットブラスト加工を施すことによって、磁気特性の改善を図ることが好ましい。というのは、圧延安定方位である斜めキューブ方位（｛１００｝＜１１０＞方位）は、通常の圧延では方位回転が起こり難く、その後の中間焼鈍等でも斜めキューブ方位として残存し、二次再結晶を阻害する。しかし、ショットブラスト加工で表層の斜めキューブ方位に歪を加えることで、中間焼鈍時の再結晶により、異なる方位とすることができ、斜めキューブ方位の残存が防止できるからである。 Furthermore, in this invention, it is preferable to aim at the improvement of a magnetic characteristic by giving shot blasting to the steel plate surface after giving the said hot-rolled sheet annealing. This is because the oblique cube orientation ({100} <110> orientation), which is a stable rolling orientation, hardly undergoes orientation rotation in normal rolling, and remains as an oblique cube orientation even in subsequent intermediate annealing, and secondary recrystallization. Inhibits. However, by applying strain to the oblique cube orientation of the surface layer by shot blasting, different orientations can be achieved by recrystallization during intermediate annealing, and the residual oblique cube orientation can be prevented.
上記の条件で熱延板焼鈍した鋼板、あるいは、さらにショットブラスト加工した鋼板は、その後、中間焼鈍を挟む２回以上の冷間圧延によって最終板厚の冷延板とする。ここで、最終冷間圧延前の中間焼鈍は、均熱温度８００〜１２００℃、均熱時間２〜３００ｓの範囲で行うのが好ましく、また、上記中間焼鈍における冷却は、８００〜４００℃の区間を冷却速度１０〜２００℃／ｓで急冷することが好ましい。均熱温度が８００℃未満あるいは均熱時間が２ｓ未満では、未再結晶組織が残存するため、一次再結晶板の組織が整粒組織とならず、二次再結晶粒が成長できない結果、磁気特性が低下する。一方、均熱温度が１２００℃より高温あるいは均熱時間が３００ｓ超えでは、ＡｌＮ、ＭｎＳｅおよびＭｎＳの溶解が進行し、二次再結晶過程でのインヒビターの抑制力が不足し、二次再結晶しなくなる結果、やはり磁気特性の低下を引き起こすからである。
また、中間焼鈍における冷却において、８００〜４００℃の区間での冷却速度を１０℃／ｓ未満とすると、カーバイドの粗大化が進行し、一次再結晶板の集合組織を改善する効果が弱まるため、磁気特性が低下する。一方、上記区間での冷却速度を２００℃／ｓより大きくすると、硬質のマルテンサイト相率が増加し、一次再結晶焼鈍後の鋼板を、一次再結晶粒が整粒で、集合組織的に優れた組織とすることができず、やはり、磁気特性の低下を引き起こす。
A steel sheet that has been hot-rolled sheet annealed under the above conditions, or a steel sheet that has been further shot blasted, is then made into a cold-rolled sheet having a final thickness by cold rolling two or more times with intermediate annealing. Here, the intermediate annealing before the final cold rolling is preferably performed in the range of a soaking temperature of 800 to 1200 ° C. and a soaking time of 2 to 300 s, and the cooling in the intermediate annealing is performed in a section of 800 to 400 ° C. Is preferably quenched at a cooling rate of 10 to 200 ° C./s. When the soaking temperature is less than 800 ° C. or the soaking time is less than 2 s, an unrecrystallized structure remains, so that the structure of the primary recrystallized plate does not become a sized structure and secondary recrystallized grains cannot grow. Characteristics are degraded. On the other hand, when the soaking temperature is higher than 1200 ° C. or the soaking time exceeds 300 s, the dissolution of AlN, MnSe and MnS proceeds, the inhibitor repressing power in the secondary recrystallization process is insufficient, and the secondary recrystallization occurs. As a result, the magnetic characteristics are also deteriorated.
Moreover, in the cooling in the intermediate annealing, when the cooling rate in the section of 800 to 400 ° C. is less than 10 ° C./s, the coarsening of the carbide proceeds and the effect of improving the texture of the primary recrystallized plate is weakened. Magnetic properties are degraded. On the other hand, when the cooling rate in the above section is larger than 200 ° C./s, the hard martensite phase ratio is increased, and the primary recrystallized grains are sized and the texture is excellent in the steel sheet after the primary recrystallization annealing. It is not possible to obtain a tissue, which again causes a decrease in magnetic properties.
中間焼鈍後、最終冷間圧延して最終板厚とした冷延板は、その後、一次再結晶焼鈍するが、この焼鈍での均熱温度は７００〜１０００℃とするのが好ましい。均熱温度が７００℃未満では、未再結晶組織が残存し、一次再結晶粒が整粒で、集合組織的に優れた組織を得ることができない。一方、均熱温度が１０００℃を超えると、二次再結晶を起こしてＧｏｓｓ方位粒が発生するおそれがあるからである。なお、この一次再結晶焼鈍は、焼鈍後のＣ含有量が０．００４ｍａｓｓ％以下となるよう、湿水素雰囲気中で脱炭を兼ねて行うことが好ましい。 After the intermediate annealing, the cold-rolled sheet having the final sheet thickness by final cold rolling is subsequently subjected to primary recrystallization annealing, and the soaking temperature in this annealing is preferably 700 to 1000 ° C. When the soaking temperature is less than 700 ° C., an unrecrystallized structure remains, primary recrystallized grains are sized, and a texture excellent in texture cannot be obtained. On the other hand, if the soaking temperature exceeds 1000 ° C., secondary recrystallization may occur and Goss orientation grains may be generated. In addition, it is preferable to perform this primary recrystallization annealing to serve as decarburization in a wet hydrogen atmosphere so that C content after annealing may be 0.004 mass% or less.
上記、一次再結晶焼鈍後の鋼板は、その後、必要に応じて、鋼板表面にＭｇＯを主成分とする焼鈍分離剤を塗布した後、二次再結晶焼鈍（仕上焼鈍）を行う。この二次再結晶焼鈍は、常法に準じて行えばよく、特に制限はない。また、この二次再結晶焼鈍は、水素雰囲気中で、純化焼鈍も兼ねて行ってよい。 The steel sheet after the primary recrystallization annealing is then subjected to secondary recrystallization annealing (finish annealing) after applying an annealing separator mainly composed of MgO to the steel sheet surface as necessary. This secondary recrystallization annealing may be performed according to a conventional method and is not particularly limited. Further, this secondary recrystallization annealing may be performed in a hydrogen atmosphere also as purification annealing.
仕上焼鈍で二次再結晶させた鋼板は、その後、絶縁コーティングを塗布・焼付ける絶縁被膜塗布工程および平坦化焼鈍工程を経て、方向性電磁鋼板（製品板）とするのが好ましい。 The steel sheet that has been secondarily recrystallized by finish annealing is preferably a grain-oriented electrical steel sheet (product board) through an insulating coating application process for applying and baking an insulating coating and a flattening annealing process.
なお、本発明の方向性電磁鋼板の製造方法においては、一次再結晶焼鈍後から仕上焼鈍で二次再結晶が開始までの間に、鋼中にＮを含有させる窒化処理を施すことも可能である。その方法としては、一次再結晶焼鈍後、ＮＨ３雰囲気中で熱処理を施したり、窒化物を焼鈍分離剤中に含有させたり、仕上焼鈍前段の雰囲気を窒化ガスとしたりする公知の技術が適用できる。 In the method for producing a grain-oriented electrical steel sheet of the present invention, it is possible to perform a nitriding treatment in which N is contained in the steel during the period from the first recrystallization annealing to the finish annealing and the second recrystallization. is there. As the method, a known technique in which heat treatment is performed in an NH 3 atmosphere after primary recrystallization annealing, nitride is contained in an annealing separator, or the atmosphere before the final annealing is used as a nitriding gas can be applied. .
さらに、最終冷間圧延後から一次再結晶焼鈍の間に鋼板表面に複数の人工溝を形成したり、平坦化焼鈍後の鋼板表面にプラズマジェットやレーザー照射、電子ビーム照射を線状に施したり、突起ロールによる線状の凹みを付与したりする磁区細分化処理を施して、鉄損の低減を図ってもよい。 Furthermore, multiple artificial grooves are formed on the steel sheet surface during the primary recrystallization annealing after the final cold rolling, or plasma jet, laser irradiation and electron beam irradiation are linearly applied to the steel sheet surface after flattening annealing. Further, it is possible to reduce the iron loss by applying a magnetic domain subdivision process such as providing a linear dent by a protruding roll.
表１に示したＡ〜Ｃの成分組成を有する鋼スラブを１４００℃に加熱した後、熱間圧延して板厚２．２ｍｍの熱延板とし、６００℃でコイルに巻取った後、２１ｖｏｌ％の酸素濃度雰囲気中に１０ｈｒ保持した。ただし、一部コイルは、コイルに巻取り、２１ｖｏｌ％の酸素濃度雰囲気中に２ｈｒ保持後、巻き解いた。その後、上記熱延板を、酸化性雰囲気下、１０２５℃×４０ｓで焼鈍した後、８００℃〜冷却停止温度Ｘ（３５０℃）までを冷却速度３０℃／ｓで急冷し、冷却停止温度Ｘから１℃／ｓで１００ｓ間徐冷した後、室温まで急冷する冷却パターンをベースとし、成分組成がＡの鋼スラブは、８００℃〜冷却停止温度Ｘ（３５０℃）までを冷却速度を５〜２００℃／ｓの範囲で変化させ、成分組成がＢの鋼スラブは、冷却停止温度Ｘを室温から５００℃の範囲で変化させ、成分組成がＣの鋼スラブは、冷却停止温度からの徐冷時間を１０〜４００ｓの範囲で変化させて、熱延板焼鈍を施した。なお、一部の熱延板には、熱延板焼鈍後、ショットブラスト加工を施し、鋼板表層に歪を付与した。
次いで、上記熱延焼鈍板を一次冷間圧延して中間板厚１．５ｍｍの冷延板とし、１０００℃×８０ｓの中間焼鈍を施した後、二次冷間圧延して最終板厚が０．２３ｍｍの冷延板とした後、８００℃×１２０ｓの脱炭を兼ねた一次再結晶焼鈍を施し、その後、鋼板表面にＭｇＯを主成分とする焼鈍分離剤を塗布してから、１１５０℃×５０ｈｒの純化を兼ねた二次再結晶焼鈍（仕上焼鈍）を施して方向性電磁鋼板とした。
上記のようにして得た各種鋼板から、試験片を採取し、磁気特性（磁束密度Ｂ８、鉄損Ｗ１７／５０）を測定し、その結果を表１に示した。
After heating a steel slab having the composition of A to C shown in Table 1 to 1400 ° C., it was hot-rolled to form a hot-rolled sheet having a thickness of 2.2 mm, wound on a coil at 600 ° C., and then 21 vol. It was kept for 10 hours in an atmosphere of oxygen concentration of%. However, some coils were wound around the coil and unwound after being kept in a 21 vol% oxygen concentration atmosphere for 2 hr. Thereafter, the hot-rolled sheet is annealed at 1025 ° C. × 40 s in an oxidizing atmosphere, and then rapidly cooled from 800 ° C. to a cooling stop temperature X (350 ° C.) at a cooling rate of 30 ° C./s. A steel slab having a component composition A based on a cooling pattern that is gradually cooled to room temperature after being gradually cooled at 1 ° C./s for 100 s, has a cooling rate of 5 to 200 from 800 ° C. to a cooling stop temperature X (350 ° C.). For steel slabs with component composition B, the cooling stop temperature X is changed from room temperature to 500 ° C., and for steel slabs with component composition C, the slow cooling time from the cooling stop temperature is used. Was changed in the range of 10 to 400 s and subjected to hot-rolled sheet annealing. Some of the hot-rolled sheets were subjected to shot blasting after the hot-rolled sheet annealing to impart strain to the steel sheet surface layer.
Next, the hot-rolled annealed sheet is subjected to primary cold rolling to obtain a cold-rolled sheet having an intermediate thickness of 1.5 mm, subjected to intermediate annealing at 1000 ° C. × 80 s, and then subjected to secondary cold rolling to obtain a final thickness of 0. After forming a cold rolled sheet of .23 mm, first recrystallization annealing was performed which also served as decarburization at 800 ° C. × 120 s, and then an annealing separator mainly composed of MgO was applied to the steel plate surface, and then 1150 ° C. × A grain-oriented electrical steel sheet was obtained by performing secondary recrystallization annealing (finish annealing) that also serves as a purification for 50 hours.
Test pieces were collected from the various steel plates obtained as described above, and magnetic properties (magnetic flux density B 8 , iron loss W 17/50 ) were measured. The results are shown in Table 1.
表１のＮｏ．１〜１１は、８００℃〜冷却停止温度Ｘ（３５０℃）までの冷却速度を変化させた例であり、上記冷却速度を１０〜１００℃／ｓの範囲としたＮｏ．２〜１０（ただし、Ｎｏ．５を除く）では、磁束密度、鉄損とも優れた特性が得られている。特に、ショットブラスト加工で熱延板表層に歪を付与したＮｏ．４では、他のものと比較して鉄損特性が優れている。
これに対して、冷却速度を５℃／ｓとしたＮｏ．１の例では、磁気特性の劣化が認められる。これは、冷却速度が遅いため、カーバイドのオストワルド成長が進行して、中間焼鈍板や一次再結晶板の集合組織の改善効果が弱まった結果であると考えられる。また、冷却速度を２００℃／ｓとしたＮｏ．１１の例でも、磁気特性の劣化が認められる。これは、カーバイド析出温度域を急冷することで、極微細カーバイドの均一分散析出が不可能となって固溶Ｃ量が増加し、さらに、硬質のマルテンサイト相が増加した結果、中間焼鈍板やその後の一次再結晶板の集合組織の改善効果が弱まった結果であると考えられる。
No. in Table 1 Nos. 1 to 11 are examples in which the cooling rate from 800 ° C. to the cooling stop temperature X (350 ° C.) was changed. In 2 to 10 (excluding No. 5), excellent characteristics of both magnetic flux density and iron loss are obtained. In particular, No. 1 was applied to the surface layer of the hot-rolled sheet by shot blasting. In No. 4, the iron loss characteristics are superior to those of the others.
On the other hand, No. 1 with a cooling rate of 5 ° C./s. In the example of 1, the deterioration of the magnetic characteristics is observed. This is considered to be a result of the fact that the Ostwald growth of carbide progresses due to the slow cooling rate, and the effect of improving the texture of the intermediate annealed plate or the primary recrystallized plate is weakened. In addition, in the case of No. In the example 11 as well, deterioration of the magnetic properties is observed. This is because by rapidly cooling the carbide precipitation temperature region, uniform dispersion precipitation of ultrafine carbide becomes impossible and the amount of solid solution C increases, and further, the hard martensite phase increases. It is thought that this is a result of the subsequent improvement effect of the texture of the primary recrystallized plate weakening.
また、表１のＮｏ．１２〜１７は、冷却停止温度を変化させた例であり、上記冷却停止温度を３５０〜２００℃の範囲としたＮｏ．１３〜１６（ただし、Ｎｏ．１５を除く）の例では良好な磁気特性が得られている。特に、ショットブラスト加工で熱延板表層に歪を付与したＮｏ．１４では、他のものと比較して鉄損特性が優れている。
これに対して、Ｎｏ．１２のように、冷却停止温度を室温とした例では、磁気特性の劣化が認められる。これは、極微細カーバイドが析出する時効時間が確保されなかったため、中間焼鈍板やその後の一次再結晶板の集合組織の改善効果が弱まったためと考えられる。また、Ｎｏ．１７のように、冷却停止温度Ｘを５００℃とし、その後、徐冷した例でも、磁気特性の劣化が認められる。これは、徐冷過程でカーバイドのオストワルド成長が進行し、中間焼鈍板やその後の一次再結晶板の集合組織の改善効果が弱まったためと考えられる。
In Table 1, No. Nos. 12 to 17 are examples in which the cooling stop temperature was changed. In the examples of 13 to 16 (excluding No. 15), good magnetic properties are obtained. In particular, No. 1 was applied to the surface layer of the hot-rolled sheet by shot blasting. In No. 14, the iron loss characteristics are superior to those of others.
In contrast, no. As shown in FIG. 12, in the example in which the cooling stop temperature is set to room temperature, deterioration of the magnetic characteristics is observed. This is presumably because the effect of improving the texture of the intermediate annealed plate and the subsequent primary recrystallized plate was weakened because the aging time for precipitation of ultrafine carbide was not secured. No. As shown in FIG. 17, even when the cooling stop temperature X is set to 500 ° C. and then gradually cooled, deterioration of the magnetic properties is observed. This is thought to be because the Ostwald growth of carbide progressed during the slow cooling process, and the effect of improving the texture of the intermediate annealed plate and the subsequent primary recrystallized plate was weakened.
また、表１のＮｏ．１８〜２４は、冷却停止温度からの徐冷時間を変化させた例であり、上記徐冷時間を４０〜２００ｓの範囲としたＮｏ．１９〜２３（ただし、Ｎｏ．２２を除く）の例では良好な磁気特性が得られている。特に、ショットブラスト加工で熱延板表層に歪を付与したＮｏ．２１では、他のものと比較して鉄損特性が優れている。
これに対して、徐冷時間が１０ｓしかないＮｏ．１８の例では、磁気特性の劣化が認められる。これは、固溶Ｃの拡散時間が十分確保されなかったことで、極微細カーバイドの析出量が不足し、中間焼鈍板の集合組織や一次再結晶板の集合組織の改善効果が弱まったためと考えられる。また、Ｎｏ．２４のように３００ｓ間徐冷した例でも、磁気特性の劣化が認められる。これは、カーバイドのオストワルド成長が進行し、中間焼鈍板や一次再結晶板の集合組織の改善効果が弱まったためと考えられる。
In Table 1, No. Nos. 18 to 24 are examples in which the slow cooling time from the cooling stop temperature was changed, and No. 18 in which the slow cooling time was in the range of 40 to 200 s. In the examples of 19 to 23 (excluding No. 22), good magnetic properties are obtained. In particular, No. 1 was applied to the surface layer of the hot-rolled sheet by shot blasting. In No. 21, the iron loss characteristics are superior to those of others.
On the other hand, no. In the example of 18, the deterioration of the magnetic characteristics is recognized. This is thought to be due to the fact that the diffusion time of the solid solution C was not sufficiently secured, the precipitation amount of ultrafine carbide was insufficient, and the effect of improving the texture of the intermediate annealing plate and the texture of the primary recrystallized plate was weakened. It is done. No. Even in the case of slow cooling for 300 s, such as 24, deterioration of the magnetic properties is observed. This is thought to be due to the fact that the Ostwald growth of carbide progressed and the effect of improving the texture of the intermediate annealed plate and the primary recrystallized plate was weakened.
なお、Ｎｏ．５，１５および２２は、熱間圧延したコイルの保持条件が不適切で、鋼板表面にＦｅ３Ｏ４を主体とする酸化物を十分に生成させることができなかった例であり、いずれも磁気特性に劣るものしか得られていない。 In addition, No. Nos. 5, 15 and 22 are examples in which the holding conditions of the hot-rolled coil were inappropriate, and oxides mainly composed of Fe 3 O 4 could not be sufficiently formed on the steel sheet surface, both of which were magnetic Only inferior properties are obtained.
以上の結果から、熱延鋼板表面にＦｅ３Ｏ４を主体とする酸化物を生成させた上で、熱延板焼鈍の冷却過程における８００℃〜冷却停止温度Ｘの間の冷却速度を１０〜１００℃／ｓの範囲し、冷却停止温度Ｘを３５０〜２００℃の範囲とし、かつ、その後の徐冷時間を４０〜２００ｓの範囲とすることで、固溶Ｃの拡散時間が十分に確保されて極微細カーバイドの均一分散析出が促進され、かつカーバイドのオストワルド成長も抑制さるので、中間焼鈍板の集合組織やその後の一次再結晶板の集合組織が改善されて、二次再結晶後の磁気特性も改善されるものと考えられる。 From the above results, after generating an oxide mainly composed of Fe 3 O 4 on the surface of the hot-rolled steel sheet, the cooling rate between 800 ° C. and the cooling stop temperature X in the cooling process of hot-rolled sheet annealing is 10 to 10. By setting the cooling stop temperature X in the range of 100 ° C./s, the cooling stop temperature X in the range of 350 to 200 ° C., and the subsequent slow cooling time in the range of 40 to 200 s, the diffusion time of the solid solution C is sufficiently secured. Therefore, the uniform dispersion precipitation of ultrafine carbide is promoted and the Ostwald growth of carbide is suppressed, so the texture of the intermediate annealed plate and the texture of the subsequent primary recrystallized plate are improved, and the magnetism after secondary recrystallization is improved. The characteristics are also expected to be improved.
Ｓｉ：３．２ｍａｓｓ％、Ｍｎ：０．０１ｍａｓｓ％、ｓｏｌ．Ａｌ：０．０２ｍａｓｓ％、Ｎ：０．０１ｍａｓｓ％、Ｓ：０．００２０ｍａｓｓ％およびＳｅ：０．０３ｍａｓｓ％を含有し、Ｃ，Ｎｉ，Ｓｎ，Ｓｂ，Ｍｏ，Ｃｕ，ＣｒおよびＰを表４に記載した範囲で含有するスラブを、１４００℃に加熱した後、熱間圧延し、板厚２．２ｍｍの熱延板とし、７２０℃でコイルに巻取り、酸素濃度１８ｖｏｌ％の雰囲気中に２４ｈｒ放置して、鋼板表層にＦｅ３Ｏ４を主体とする酸化物を生成させた。次いで、上記熱延板に、１０００℃×４０ｓで焼鈍後、８００〜３５０℃間を３０℃／ｓで冷却し、３５０℃に到達後、１℃／ｓで１００ｓ間徐冷した後、室温まで急冷する熱延板焼鈍を施した。次いで、一部のものにはショットブラスト加工を施し、他のものはそのままで酸洗した後、一次冷間圧延して板厚１．５ｍｍの冷延板とし、１０００℃×８０ｓの中間焼鈍を施した後、二次冷間圧延して板厚０．２３ｍｍの冷延板とした。その後、この冷延板に、８００℃×１２０ｓの脱炭を兼ねた一次再結晶焼鈍し、鋼板表面にＭｇＯを主成分とする焼鈍分離剤を塗布してから、１１５０℃×５０ｈｒの純化を兼ねた二次再結晶焼鈍（仕上焼鈍）を施して、方向性電磁鋼板とした。上記のようにして得た鋼板から試験片を採取し、磁気特性を測定し結果を表２に示した。 Si: 3.2 mass%, Mn: 0.01 mass%, sol. Al: 0.02 mass%, N: 0.01 mass%, S: 0.0020 mass%, and Se: 0.03 mass%, C, Ni, Sn, Sb, Mo, Cu, Cr, and P in Table 4 The slab contained within the stated range is heated to 1400 ° C. and then hot rolled to form a hot-rolled sheet having a thickness of 2.2 mm, wound on a coil at 720 ° C., and left in an atmosphere with an oxygen concentration of 18 vol% for 24 hours. Then, an oxide mainly composed of Fe 3 O 4 was formed on the steel sheet surface layer. Next, after annealing at 1000 ° C. × 40 s to the above hot-rolled sheet, the temperature between 800 to 350 ° C. is cooled at 30 ° C./s, and after reaching 350 ° C., gradually cooled at 1 ° C./s for 100 s, and then to room temperature Hot-rolled sheet annealing for rapid cooling was performed. Next, some parts are shot blasted, others are pickled as they are, and then cold-rolled to a cold rolled sheet with a thickness of 1.5 mm, and subjected to intermediate annealing at 1000 ° C. × 80 s. After the application, secondary cold rolling was performed to obtain a cold-rolled sheet having a thickness of 0.23 mm. Thereafter, this cold-rolled sheet is subjected to primary recrystallization annealing that also serves as decarburization at 800 ° C. for 120 s, and after applying an annealing separator mainly composed of MgO to the steel sheet surface, it also serves as purification at 1150 ° C. for 50 hours. Secondary recrystallization annealing (finish annealing) was performed to obtain a grain-oriented electrical steel sheet. A test piece was collected from the steel plate obtained as described above, the magnetic properties were measured, and the results are shown in Table 2.
表２のＮｏ．１〜５に示すように、Ｃ含有量を０．００３〜０．２０ｍａｓｓ％の範囲で変化させた場合には、Ｎｏ．２〜４の範囲、つまり、Ｃ：０．００５〜０．１５ｍａｓｓ％の範囲で良好な磁気特性が得られることがわかる。Ｎｏ．１の磁気特性が劣化した原因は、Ｃ含有量が少ないため、一次再結晶板の集合組織改善効果が弱かったためと考えられる。また、Ｎｏ．５の磁気特性が低下した原因は、Ｃ含有量が多いため、一次再結晶焼鈍での脱炭が不完全となって一次再結晶板の集合組織が劣化したためと考えられる。
また、Ｎｏ．６〜２６は、Ｃ含有量を０．０５ｍａｓｓ％一定とし、Ｎｉ，Ｓｎ，Ｓｂ，Ｍｏ，Ｃｕ，ＣｒおよびＰの含有量を変化させた例であるが、表２に示した本発明に適合する添加量の範囲では、いずれも優れた磁気特性が得られていることがわかる。特に、ショットブラスト加工で鋼板表層に歪を付与した例では、鉄損Ｗ１７／５０が０．９０Ｗ／ｋｇ以下を安定して達成できている。
No. in Table 2 As shown to 1-5, when changing C content in the range of 0.003-0.20 mass%, it is No. It can be seen that good magnetic properties can be obtained in the range of 2 to 4, that is, in the range of C: 0.005 to 0.15 mass%. No. The reason for the deterioration of the magnetic characteristics of No. 1 is considered to be that the texture improvement effect of the primary recrystallized plate was weak because the C content was small. No. The reason for the decrease in magnetic properties of No. 5 is considered to be that the decarbonization in the primary recrystallization annealing was incomplete and the texture of the primary recrystallized plate was deteriorated because of the high C content.
No. Examples 6 to 26 are examples in which the C content is kept constant at 0.05 mass% and the contents of Ni, Sn, Sb, Mo, Cu, Cr, and P are changed, but are suitable for the present invention shown in Table 2. It can be seen that excellent magnetic properties are obtained in the range of the amount to be added. In particular, in the example in which strain is applied to the steel sheet surface layer by shot blasting, the iron loss W 17/50 can stably achieve 0.90 W / kg or less.
Claims (3)
熱間圧延後の巻取温度を７６０〜４６０℃の範囲として鋼板表面にＦｅ３Ｏ４を主体とする酸化物を生成させた後、
酸化性雰囲気中で、８００℃以上に加熱後、８００℃から３５０〜２００℃間の冷却停止温度までを平均冷却速度１０〜１００℃／ｓで冷却し、その後、冷却停止温度から冷却速度５℃／ｓ以下で４０〜２００ｓ間徐冷する熱延板焼鈍を施すことを特徴とする方向性電磁鋼板の製造方法。 C: 0.005-0.15 mass%, Si: 2.5-7.0 mass%, Mn: 0.005-0.3 mass%, sol. Al: 0.01-0.05 mass%, N: 0.002-0.012 mass%, one or two of S and Se in total contain 0.005-0.05 mass%, with the balance being After reheating a steel material having a component composition composed of Fe and inevitable impurities, hot rolling, hot rolled sheet annealing, cold rolling at least twice with intermediate annealing in between, and then cold rolling a final sheet thickness And then, in the method for producing a grain-oriented electrical steel sheet comprising a series of steps of primary recrystallization annealing and secondary recrystallization annealing,
After generating the oxide mainly composed of Fe 3 O 4 on the steel sheet surface with the coiling temperature after hot rolling in the range of 760 to 460 ° C.,
After heating to 800 ° C. or higher in an oxidizing atmosphere, cooling is performed at an average cooling rate of 10 to 100 ° C./s from 800 ° C. to a cooling stop temperature of 350 to 200 ° C., and then the cooling stop temperature is set to 5 ° C. The manufacturing method of the grain-oriented electrical steel sheet characterized by performing the hot-rolled sheet annealing which anneals for 40-200 s below / s.
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| JP2017095771A (en) * | 2015-11-25 | 2017-06-01 | 新日鐵住金株式会社 | Oriented electromagnetic steel sheet and manufacturing method therefor, hot rolled sheet for oriented electromagnetic steel sheet and manufacturing method therefor |
| JP2019183271A (en) * | 2018-03-30 | 2019-10-24 | Ｊｆｅスチール株式会社 | Production method of grain-oriented electromagnetic steel sheet |
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