JP5487215B2 - Manufacturing method of high-strength, high-stretched steel sheet, hot-rolled steel sheet, cold-rolled steel sheet, galvanized steel sheet and galvanized alloyed steel sheet - Google Patents
Manufacturing method of high-strength, high-stretched steel sheet, hot-rolled steel sheet, cold-rolled steel sheet, galvanized steel sheet and galvanized alloyed steel sheet Download PDFInfo
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- JP5487215B2 JP5487215B2 JP2011543380A JP2011543380A JP5487215B2 JP 5487215 B2 JP5487215 B2 JP 5487215B2 JP 2011543380 A JP2011543380 A JP 2011543380A JP 2011543380 A JP2011543380 A JP 2011543380A JP 5487215 B2 JP5487215 B2 JP 5487215B2
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- 229910000831 Steel Inorganic materials 0.000 title claims description 114
- 239000010959 steel Substances 0.000 title claims description 114
- 239000010960 cold rolled steel Substances 0.000 title claims description 31
- 238000004519 manufacturing process Methods 0.000 title claims description 25
- 229910001335 Galvanized steel Inorganic materials 0.000 title claims description 20
- 239000008397 galvanized steel Substances 0.000 title claims description 20
- 230000009466 transformation Effects 0.000 claims description 39
- 229910001566 austenite Inorganic materials 0.000 claims description 38
- 229910001563 bainite Inorganic materials 0.000 claims description 31
- 238000005098 hot rolling Methods 0.000 claims description 31
- 238000004804 winding Methods 0.000 claims description 19
- 230000000717 retained effect Effects 0.000 claims description 18
- 238000005097 cold rolling Methods 0.000 claims description 15
- 239000012535 impurity Substances 0.000 claims description 13
- 229910052748 manganese Inorganic materials 0.000 claims description 12
- 229910000734 martensite Inorganic materials 0.000 claims description 12
- 229910052717 sulfur Inorganic materials 0.000 claims description 12
- 238000000137 annealing Methods 0.000 claims description 11
- 238000007747 plating Methods 0.000 claims description 10
- 230000032683 aging Effects 0.000 claims description 9
- 238000001816 cooling Methods 0.000 claims description 9
- 238000010438 heat treatment Methods 0.000 claims description 6
- 238000005275 alloying Methods 0.000 claims description 3
- 230000015572 biosynthetic process Effects 0.000 description 10
- 238000000034 method Methods 0.000 description 7
- 229910001562 pearlite Inorganic materials 0.000 description 5
- 229910000794 TRIP steel Inorganic materials 0.000 description 4
- 229910052799 carbon Inorganic materials 0.000 description 4
- 230000000052 comparative effect Effects 0.000 description 4
- 230000000694 effects Effects 0.000 description 4
- 238000005096 rolling process Methods 0.000 description 4
- 229910000859 α-Fe Inorganic materials 0.000 description 4
- 229910000797 Ultra-high-strength steel Inorganic materials 0.000 description 3
- 229910045601 alloy Inorganic materials 0.000 description 3
- 239000000956 alloy Substances 0.000 description 3
- 229910052804 chromium Inorganic materials 0.000 description 3
- 239000013585 weight reducing agent Substances 0.000 description 3
- CURLTUGMZLYLDI-UHFFFAOYSA-N Carbon dioxide Chemical compound O=C=O CURLTUGMZLYLDI-UHFFFAOYSA-N 0.000 description 2
- 239000000463 material Substances 0.000 description 2
- 229910052750 molybdenum Inorganic materials 0.000 description 2
- 229910001568 polygonal ferrite Inorganic materials 0.000 description 2
- 230000000087 stabilizing effect Effects 0.000 description 2
- 239000002436 steel type Substances 0.000 description 2
- QVGXLLKOCUKJST-UHFFFAOYSA-N atomic oxygen Chemical compound [O] QVGXLLKOCUKJST-UHFFFAOYSA-N 0.000 description 1
- 229910002092 carbon dioxide Inorganic materials 0.000 description 1
- 239000001569 carbon dioxide Substances 0.000 description 1
- 150000001875 compounds Chemical class 0.000 description 1
- 230000001276 controlling effect Effects 0.000 description 1
- 238000005261 decarburization Methods 0.000 description 1
- 238000009792 diffusion process Methods 0.000 description 1
- 238000005516 engineering process Methods 0.000 description 1
- 239000012467 final product Substances 0.000 description 1
- 239000000446 fuel Substances 0.000 description 1
- 229910052742 iron Inorganic materials 0.000 description 1
- 239000002184 metal Substances 0.000 description 1
- 229910052751 metal Inorganic materials 0.000 description 1
- 150000001247 metal acetylides Chemical class 0.000 description 1
- 239000000203 mixture Substances 0.000 description 1
- 229910052759 nickel Inorganic materials 0.000 description 1
- 229910052760 oxygen Inorganic materials 0.000 description 1
- 239000001301 oxygen Substances 0.000 description 1
- 230000000704 physical effect Effects 0.000 description 1
- 239000000047 product Substances 0.000 description 1
- 230000001105 regulatory effect Effects 0.000 description 1
- 238000007711 solidification Methods 0.000 description 1
- 230000008023 solidification Effects 0.000 description 1
- 239000000126 substance Substances 0.000 description 1
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/001—Austenite
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/002—Bainite
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- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Heat Treatment Of Sheet Steel (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
Description
本発明は、自動車の内板及び構造部材に使用することができる鋼板及びその製造方法に関し、より詳しくは、高強度、高延伸鋼板及び熱延鋼板、冷延鋼板、亜鉛メッキ鋼板及び亜鉛メッキ合金化鋼板の製造方法に関する。 The present invention relates to a steel plate that can be used for an inner plate and a structural member of an automobile, and a method for manufacturing the same, and more particularly, a high-strength, high-stretched steel plate and hot-rolled steel plate, cold-rolled steel plate, galvanized steel plate, and galvanized alloy. The present invention relates to a method for producing a chemical steel sheet.
最近、環境保存のために二酸化炭素の排出量を規制しており、自動車の燃費改善に対する要求が増加されている。また、衝突の際に乗客の安全を確保するために自動車その車体の衝突特性を中心とする安全性向上も要求されている。前記のような自動車その車体の軽量化と安定性をともに達成するための方案として高強度鋼板の加工性向上が要求されている。 Recently, carbon dioxide emissions have been regulated to preserve the environment, and demands for improving automobile fuel efficiency have increased. In addition, in order to ensure the safety of passengers in the event of a collision, it is also required to improve safety centering on the collision characteristics of the automobile and the vehicle body. Improvement of workability of high-strength steel sheets is required as a method for achieving both the weight reduction and stability of the automobile as described above.
使用の鋼板の降伏強度と引張強度が増加すればするほど車体の軽量化効果が増加されるので自動車業界では、高強度鋼板を採用しようとする工夫が続いている。近頃には、引張強度９８０ＭＰａ級以上のＡＨＳＳ（Ａｄｖａｎｃｅｄ Ｈｉｇｈ Ｓｔｒｅｎｇｔｈ Ｓｔｅｅｌ）鋼の適用が行われている。しかし、９８０ＭＰａ級以上の高い強度の場合には、相対的に延伸率が下がるという短所がある。 As the yield strength and tensile strength of the steel sheet used increase, the weight reduction effect of the vehicle body increases, and therefore the automobile industry continues to devise to adopt high-strength steel sheets. Recently, AHSS (Advanced High Strength Steel) steel having a tensile strength of 980 MPa or higher is being applied. However, in the case of a high strength of 980 MPa or higher, there is a disadvantage that the stretching ratio is relatively lowered.
これを補完するために残留オーステナイトを活用するＴＲＩＰ鋼の開発が行っているが、相変わらず引張強度と延伸率の積が２５、０００ＭＰａ％を超えることができず、強度と延伸率の均衡がよくないという短所がある。 In order to compensate for this, TRIP steel utilizing retained austenite has been developed, but the product of tensile strength and stretch ratio cannot exceed 25,000 MPa% as usual, and the balance between strength and stretch ratio is not good. There are disadvantages.
特許文献１を基に特許文献２と特許文献３などに高炭素高合金鋼を低温にて恒温変態させ微細なベイナイトを得て高軽度、高強度、高延性の鋼鉄を製造する方法が記述されている。
Based on
前記の発明には、一定性分の鋼を均質化処理した後空冷を行って、この鋼をさらに加熱した後恒温変態させて少なくとも５０％のベイナイト組織を有するか、または少なくとも６５％以上のベイナイト組織を有し、組織の残りは残留オーステナイト相に構成される鋼を製造する方法を提案している。これを介してＨｖ４０９以上の超高強度鋼を製造する方案を提示しているが、変態時間がかかりすぎて熱延や焼鈍工程に適用できないという短所がある。 In the above invention, the steel of a certain amount is homogenized and then air-cooled, and the steel is further heated and isothermally transformed to have at least 50% bainite structure, or at least 65% or more bainite. Proposes a method for producing a steel having a structure, the rest of the structure being composed of residual austenite phase. Through this, a method for producing ultra high strength steel of Hv409 or higher is presented, but it has a disadvantage that it takes too much transformation time and cannot be applied to hot rolling and annealing processes.
また、特許文献４では、一定成分からなる鋼材をオーステナイト化処理した後、Ｂｓ温度直下からＢｓ−１００℃間の温度にて恒温変態させることで１０００ＭＰａ以上の引張強度と１０％以上の延伸率を得る方法を提案しているが、Ｃｒ，Ｎｉ，Ｍｏなど高価の合金元素が添加されて製造単価が非常に高価であることと熱延鋼板にだけ制限されるという短所がある。 Moreover, in patent document 4, after carrying out the austenitizing process of the steel material which consists of a fixed component, the tensile strength of 1000 Mpa or more and the extending | stretching rate of 10% or more are carried out by carrying out the isothermal transformation at the temperature between Bs-100 degreeC right under Bs temperature. However, there are disadvantages in that expensive alloy elements such as Cr, Ni, and Mo are added and the manufacturing unit price is very expensive, and that only the hot-rolled steel sheet is limited.
Ｓａｎｄｖｉｋなどは、非特許文献１で、Ｃ：０．６５〜０．９９重量％、Ｓｉ：２〜２．７８重量％、Ｍｎ：０．５重量％、Ｃｒ：０．０２〜０．７９重量％、Ｔｉ：０〜０．０３重量％の添加鋼を活用してオーステナイト熱処理した後、２９０〜３８０℃の温度で熱処理した超高強度鋼を提案しているが、Ｓｉの含量が非常に高く、圧延性が非常に劣るため、熱間圧延及び冷間圧延の適用が困難で、熱延鋼板だけに制限されるという短所がある。
Sandvik et al.,
本発明は、前記従来の問題を解決するために、鋼の成分系及びベイナイトと残留オーステナイトのラス（ｌａｔｈ）幅を制御して高強度、高延伸の鋼板を提供して、また、本発明の他の目的は、高強度、高延伸の熱延鋼板、冷延鋼板、亜鉛メッキ鋼板及び亜鉛メッキ合金化鋼板を製造することができる方法を提供することにある。 In order to solve the above-mentioned conventional problems, the present invention provides a high-strength, high-stretched steel sheet by controlling the composition system of steel and the lath width of bainite and retained austenite. Another object is to provide a method capable of producing a high-strength, high-stretched hot-rolled steel sheet, cold-rolled steel sheet, galvanized steel sheet, and galvanized alloyed steel sheet.
本発明の一実施例は、重量％で、Ｃ：０．５〜１．０％、Ｍｎ：０．０１〜２．０％、Ｓｉ：１．０〜２．０％、Ａｌ：０．０１〜２．０％、Ｓ：０．０１２％以下、残部Ｆｅ、及びその他の不可避な不純物を含む鋼板であって、前記鋼板の内部組織は、ベイナイトと残留オーステナイトからなり、ベイナイトの平均レス（ｌａｔｈ）幅：０．３μｍ以下、残留オーステナイトの平均レス幅：０．３μｍ以下である鋼板に関する。前記鋼板は、降伏強度：７００ＭＰａ以上、引張強度：９８０ＭＰａ以上、及び延伸率：２０％以上であることができる。また、前記鋼板は重量％で、０．５％以下のＭｏ及び１．０％以下のＣｒのうち１種または２種をさらに含むことができる。また、前記鋼板は、重量％で、０．００５％以下のＢ及び０．０６％以下のＴｉをさらに含むことができる。 In one embodiment of the present invention, by weight, C: 0.5 to 1.0%, Mn: 0.01 to 2.0%, Si: 1.0 to 2.0%, Al: 0.01 -2.0%, S: 0.012% or less, the remainder Fe, and other inevitable impurities, the internal structure of the steel plate is composed of bainite and residual austenite, the average bainite less ) Width: 0.3 μm or less, the average loess width of retained austenite: relating to a steel sheet having a width of 0.3 μm or less. The steel sheet may have a yield strength of 700 MPa or more, a tensile strength of 980 MPa or more, and a draw ratio of 20% or more. The steel sheet may further include one or two of 0.5% or less of Mo and 1.0% or less of Cr. The steel sheet may further include 0.005% or less B and 0.06% or less Ti by weight%.
また、本発明の他の実施例は、重量％で、Ｃ：０．５〜１．０％、Ｍｎ：０．０１〜２．０％、Ｓｉ：１．０〜２．０％，Ａｌ：０．０１〜２．０％、Ｓ：０．０１２％以下、残部Ｆｅ、及びその他の不可避な不純物を含むスラブを熱間圧延してＡｒ３変態点以上で熱間圧延を仕上げして、Ｍｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で巻取りすることを含む、高強度高延伸熱延鋼板の製造方法に関する。 Moreover, the other Example of this invention is weight%, C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: A slab containing 0.01 to 2.0%, S: 0.012% or less, the balance Fe, and other inevitable impurities is hot-rolled and hot-rolled at the Ar3 transformation point or higher, and Ms ( The present invention relates to a method for producing a high-strength, high-stretched hot-rolled steel sheet including winding in a range of martensite transformation start temperature) to 450 ° C.
また、本発明の他の実施例は、重量％で、Ｃ：０．５〜１．０％、Ｍｎ：０．０１〜２．０％、Ｓｉ：１．０〜２．０％，Ａｌ：０．０１〜１．０％、Ｓ：０．０１２％以下、残部Ｆｅ、及びその他の不可避な不純物を含む鋼を熱間圧延して、Ａｒ３変態点以上で熱間圧延を仕上げる熱間圧延ステップ、前記のように熱間圧延された熱延鋼板を６５０℃以下で巻取りする巻取りステップ、前記熱延鋼板を冷間圧延した後、Ａｃ３温度以上で焼鈍した後冷却するステップ、及び前記冷延鋼板をＭｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で過時効処理する時効処理ステップを含む、高強度高延伸冷延鋼板の製造方法に関する。 Moreover, the other Example of this invention is weight%, C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: Hot rolling step of hot-rolling steel containing 0.01 to 1.0%, S: 0.012% or less, the balance Fe, and other inevitable impurities, and finishing hot rolling at an Ar3 transformation point or higher A winding step of winding the hot-rolled steel sheet hot-rolled as described above at 650 ° C. or lower, a step of cold-rolling the hot-rolled steel sheet, annealing it at an Ac3 temperature or higher, and cooling the cold-rolled steel sheet; The present invention relates to a method for producing a high-strength, high-stretched cold-rolled steel sheet, comprising an aging treatment step of over-aging the rolled steel sheet in a range of Ms (martensitic transformation start temperature) to 450 ° C.
また、本発明の他の実施例は、重量％で、Ｃ：０．５〜１．０％、Ｍｎ：０．０１〜２．０％、Ｓｉ：１．０〜２．０％，Ａｌ：０．０１〜１．０％、Ｓ：０．０１２％以下、残部Ｆｅ及びその他の不可避な不純物を含むスラブを熱間圧延して、Ａｒ３変態点以上で熱間圧延を仕上げる熱間圧延ステップ、前記のように熱間圧延された熱延鋼板を６５０℃以下で巻取りする巻取りステップ、前記熱延鋼板を冷間圧延した後、Ａｃ３温度以上にて焼鈍した後冷却するステップ、前記冷延鋼板をＭｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で過時効処理する時効処理ステップ、及び前記時効処理した冷延鋼板をメッキ槽を通過させて鋼板表面にメッキ層を形成させるステップを含む高強度高延伸亜鉛メッキ鋼板の製造方法に関する。 Moreover, the other Example of this invention is weight%, C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: Hot rolling step of hot rolling a slab containing 0.01 to 1.0%, S: 0.012% or less, the balance Fe and other inevitable impurities, and finishing the hot rolling at an Ar3 transformation point or higher, A winding step of winding the hot-rolled steel sheet hot-rolled as described above at 650 ° C. or lower, a step of cold-rolling the hot-rolled steel sheet, annealing at an Ac3 temperature or higher, and cooling the cold-rolled steel sheet, the cold-rolling An aging treatment step of over-aging the steel sheet in a range of Ms (martensitic transformation start temperature) to 450 ° C., and a step of passing the aging-treated cold-rolled steel sheet through a plating tank to form a plating layer on the steel sheet surface. The present invention relates to a method for producing a high-strength, highly-stretched galvanized steel sheet.
また、本発明の他の実施例は、重量％で、Ｃ：０．５〜１．０％、Ｍｎ：０．０１〜２．０％、Ｓｉ：１．０〜２．０％，Ａｌ：０．０１〜１．０％、Ｓ：０．０１２％以下、残部Ｆｅ及びその他の不可避な不純物を含むスラブを熱間圧延して、Ａｒ３変態点以上で熱間圧延を仕上げる熱間圧延ステップ、前記のように熱間圧延された熱延鋼板を６５０℃以下で巻取りする巻取りステップ、前記熱延鋼板を冷間圧延した後、Ａｃ３温度以上にて焼鈍した後冷却するステップ、前記冷延鋼板をＭｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で過時効処理する時効処理ステップ、前記時効処理した冷延鋼板をメッキ槽を通過させて鋼板表面にメッキ層を形成させるステップ及び前記メッキ層を形成した亜鉛メッキ鋼板を合金化熱処理して亜鉛メッキ合金化鋼板を製造するステップを含む、高強度高延伸亜鉛メッキ合金化鋼板の製造方法に関する。 Moreover, the other Example of this invention is weight%, C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: Hot rolling step of hot rolling a slab containing 0.01 to 1.0%, S: 0.012% or less, the balance Fe and other inevitable impurities, and finishing the hot rolling at an Ar3 transformation point or higher, A winding step of winding the hot-rolled steel sheet hot-rolled as described above at 650 ° C. or lower, a step of cold-rolling the hot-rolled steel sheet, annealing at an Ac3 temperature or higher, and cooling the cold-rolled steel sheet, the cold-rolling An aging treatment step of over-aging the steel sheet in a range of Ms (Martensite transformation start temperature) to 450 ° C., a step of passing the aging-treated cold-rolled steel sheet through a plating tank and forming a plating layer on the surface of the steel sheet, and the plating Alloying heat treatment of galvanized steel sheet Comprising the steps of manufacturing a lead-plated alloying steel, a method for producing a high-strength and high stretch galvanized alloyed steel.
本発明は、自動車鋼板の軽量化と衝突安定性の向上に寄与することができる、降伏強度が７００ＭＰａ以上、引張強度が９８０ＭＰａ以上、及び延伸率が２０％以上である、鋼板、熱延鋼板、冷延鋼板、亜鉛メッキ鋼板、及び亜鉛メッキ合金化鋼板を提供する。 The present invention can contribute to weight reduction of automobile steel plate and improvement of collision stability, steel plate, hot rolled steel plate, yield strength is 700 MPa or more, tensile strength is 980 MPa or more, and stretch ratio is 20% or more, A cold-rolled steel sheet, a galvanized steel sheet, and a galvanized alloyed steel sheet are provided.
以下、本発明の成分系を説明する。 Hereinafter, the component system of the present invention will be described.
Ｃ：０．５〜１．０重量％
Ｃの含量が０．５重量％未満である場合には強度の確保は可能であるが、安定な残留オーステナイトの安定性が劣位して２０％以上の延伸率の確保が困難となる恐れがある。一方、Ｃの含量が１．０重量％を超える場合には、ベイナイト変態が著しく遅くなりパーライトの生成が促進されることによってむしろ物性が劣化する恐れがある。よって、前記Ｃの含量は、０．５〜１．０重量％の範囲に限定することが好ましい。
C: 0.5 to 1.0% by weight
If the C content is less than 0.5% by weight, the strength can be ensured, but the stability of stable retained austenite is inferior, and it may be difficult to ensure a stretch ratio of 20% or more. . On the other hand, when the C content exceeds 1.0% by weight, the bainite transformation is remarkably slowed, and the formation of pearlite is promoted, so that the physical properties may be deteriorated. Therefore, the C content is preferably limited to a range of 0.5 to 1.0% by weight.
Ｍｎ：０．０１〜２．０重量％
Ｍｎは、鋼の製造工程のうちに不可避に含有されるＳとＦｅが結合したＦｅＳ形成によって赤熱脆性を防止するために添加されるが、Ｍｎの含量が０．０１重量％未満である場合には、その含量が少な過ぎて赤熱脆性が生じる恐れがある。一方、Ｍｎの含量が２．０重量％を超える場合には、ベイナイト変態速度を遅くして焼鈍工程の生産性を阻害する恐れがある。よって、前記Ｍｎの含量は、０．０１〜２．０重量％の範囲に限定することが好ましい。
Mn: 0.01 to 2.0% by weight
Mn is added in order to prevent red heat embrittlement by FeS formation in which S and Fe, which are inevitably contained in the steel manufacturing process, are combined, but the Mn content is less than 0.01% by weight. May be red brittle due to its low content. On the other hand, if the Mn content exceeds 2.0% by weight, the bainite transformation rate may be slowed to hinder the productivity of the annealing process. Therefore, the Mn content is preferably limited to a range of 0.01 to 2.0% by weight.
Ｓｉ：１．０〜２．０重量％
Ｓｉは、パーライト変態を促進させて未変態オーステナイト中のＣの含量を増加させて最終製品の残留オーステナイト分率を向上させるので、鋼中への積極的な添加を要する元素である。
Si: 1.0 to 2.0% by weight
Since Si promotes pearlite transformation and increases the content of C in untransformed austenite to improve the retained austenite fraction of the final product, it is an element that needs to be actively added to steel.
本発明のように残留オーステナイトを含有する変態組織鋼の生成のためには、Ｓｉは必然的に添加され、Ｓｉの含量が１．０重量％未満である場合には残留オーステナイトの安定化効果が大きくならない恐れがある。一方、Ｓｉの含量が２．０重量％を超える場合には熱間圧延及び冷間圧延の際にクラックが生じるなど圧延性が低下する恐れがある。よって、前記Ｓｉの含量は１．０〜２．０重量％の範囲に限定することが好ましい。 In order to produce a transformation structure steel containing retained austenite as in the present invention, Si is inevitably added, and when the Si content is less than 1.0% by weight, the effect of stabilizing retained austenite is obtained. There is a risk of not growing. On the other hand, if the Si content exceeds 2.0% by weight, the rollability may be deteriorated, for example, cracks may occur during hot rolling and cold rolling. Therefore, the Si content is preferably limited to a range of 1.0 to 2.0% by weight.
Ａｌ：０．０１〜２．０重量％
Ａｌは、二つの目的を有して添加されるが、その一つは鋼中に存在する酸素を除去して凝固の際に非金属介在物の形成を防止して、他の一つは本発明のように残留オーステナイトを安定化させるためのＣのオーステナイトへの拡散を促進するためである。
Al: 0.01 to 2.0% by weight
Al is added for two purposes, one of which removes the oxygen present in the steel to prevent the formation of non-metallic inclusions during solidification and the other is the main. This is to promote diffusion of C to austenite for stabilizing retained austenite as in the invention.
Ａｌの含量が０．０１重量％未満である場合には、前記添加目的を達することができず、２．０重量％を超える場合には製鋼の原単位の上昇が問題になる恐れがある。よって、前記Ａｌの含量は、０．０１〜２．０重量％の範囲に限定することが好ましい。 When the Al content is less than 0.01% by weight, the purpose of the addition cannot be achieved, and when it exceeds 2.0% by weight, an increase in the basic unit of steel production may become a problem. Therefore, the content of Al is preferably limited to a range of 0.01 to 2.0% by weight.
しかし、本発明で冷間圧延の後に焼鈍を行う冷延鋼板、メッキ鋼板、メッキ合金化鋼板の場合には、Ａｌの含量が過多になる場合にＡｃ３温度が高くなることによってＡｃ３以上の温度においての熱処理が実際的に困難になる恐れがあるので、冷延鋼板、亜鉛メッキ鋼板及び亜鉛メッキ合金化鋼板に対しては上限を１．０重量％に限定することが好ましい。 However, in the case of the cold rolled steel sheet, the plated steel sheet, and the plated alloyed steel sheet, which are annealed after cold rolling in the present invention, the Ac3 temperature is increased when the Al content is excessive, so that the temperature is higher than Ac3. Therefore, it is preferable to limit the upper limit to 1.0% by weight for cold-rolled steel sheets, galvanized steel sheets, and galvanized alloyed steel sheets.
Ｓ：０．０１２重量％以下
前記Ｓは、ＦｅＳ形成によって赤熱脆性の問題を誘発するので、Ｓの量を少なくするように管理することが必要であるので、０．０１２重量％以下に制限する。下限を規定しない理由は、前記の同じ理由でＳの量を下げれば下げるほど有利であるため制限はしない。
S: 0.012 wt% or less Since S induces a red hot brittleness problem due to the formation of FeS, it is necessary to control the amount of S to be small, so it is limited to 0.012 wt% or less. . The reason why the lower limit is not specified is not limited because the lower the amount of S for the same reason, the more advantageous it is.
Ｍｏ及びＣｒは、必要に応じて選択的に添加される成分である。
Ｍｏ：０．５重量％以下
前記Ｍｏは、パーライト相の生成を抑制するために添加されるが、０．５重量％を超える場合には、高価のＭｏに起因して製造原価が急上昇する。よって、Ｍｏの含量は、０．５重量％以下の範囲に限定することが好ましい。
Mo and Cr are components that are selectively added as necessary.
Mo: 0.5 wt% or less The Mo is added to suppress the formation of a pearlite phase, but when it exceeds 0.5 wt%, the manufacturing cost rapidly increases due to expensive Mo. Therefore, the Mo content is preferably limited to a range of 0.5% by weight or less.
Ｃｒ：１．０重量％以下
前記Ｃｒは、フェライト生成を抑制しベイナイト変態を促進するために添加されることができるが、１．０重量％を超える場合にＣｒ炭化物の形成が促進されることによって固溶Ｃ量が低減する恐れがある。よって、前記Ｃｒの含量は、１．０重量％以下の範囲に限定することが好ましい。
Cr: 1.0 wt% or less The Cr can be added to suppress ferrite formation and promote bainite transformation, but when it exceeds 1.0 wt%, formation of Cr carbide is promoted. Therefore, there is a risk that the amount of dissolved C may be reduced. Therefore, the Cr content is preferably limited to a range of 1.0% by weight or less.
Ｂ及びＴｉは、必要に応じて選択的に添加することができ、Ｂの効果を極大化するためには、不純物として存在するＮとＢが結合することを抑制するために、Ｔｉを共に添加することができる。 B and Ti can be selectively added as necessary, and in order to maximize the effect of B, both Ti and Ti are added to suppress the bonding of N and B present as impurities. can do.
Ｂ：０．００５重量％以下
前記Ｂは、フェライトの生成を抑制するためにさらに添加されるが、Ｂが０．００５重量％を超える場合には、Ｆｅ２３（Ｃ，Ｂ）６化合物の生成が促進され、フェライト生成がむしろ促進される恐れがある。よって、前記Ｂの含量は０．００５重量％以下の範囲に限定することが好ましい。
B: 0.005 wt% or less The B is further added to suppress the formation of ferrite, but when B exceeds 0.005 wt%, the formation of Fe23 (C, B) 6 compound is There is a risk that ferrite formation is rather accelerated. Therefore, the B content is preferably limited to a range of 0.005% by weight or less.
Ｔｉ：０．０６重量％以下
Ｔｉの含量が０．０６重量％を超える場合には、十分にＮ（通常０．０１％以下の濃度に不純物として存在）を除去してからも、剰余Ｔｉが多量に残ってＣと結合する問題がある。よって、前記Ｔｉの含量は、０．０６重量％以下の範囲に限定することが好ましい。
Ti: 0.06% by weight or less When the Ti content exceeds 0.06% by weight, the residual Ti remains even after sufficiently removing N (usually present as an impurity at a concentration of 0.01% or less). There is a problem that a large amount remains and bonds with C. Accordingly, the Ti content is preferably limited to a range of 0.06% by weight or less.
以下、本発明の鋼板の内部組織、降伏強度、引張強度及び延伸率に対して説明する。 Hereinafter, the internal structure, yield strength, tensile strength, and stretch ratio of the steel sheet of the present invention will be described.
前記内部組織は、ベイナイトと残留オーステナイトからなっており、前記ベイナイトは等軸晶の多角形フェライトがない組織であって、炭化物が存在しないベイナイトが好ましい。 The internal structure is composed of bainite and retained austenite, and the bainite is a structure free from equiaxed polygonal ferrite, and is preferably bainite free from carbides.
一般的に、ＴＲＩＰ鋼の微細組織は、残留オーステナイトを含有しなければならなく、技術的には、が０％を超える残留オーステナイト分率で含有することによって、延性の向上に残留オーステナイトが実質的な役割を果たす鋼を意味する。よって、本発明の実施形態として、ＴＲＩＰ鋼である場合には、微細組織は、残留オーステナイト分率が０％を超えて残部はベイナイトであることが好ましい。但し、本発明の微細組織は、残留オーステナイト分率が１０〜４０％の範囲で、残部はベイナイトであることがより好ましい。残留オーステナイトの分率が１０％未満である場合には、公知のＴＲＩＰ鋼と類似な水準であり、４０％を超える場合には残留オーステナイトの安定性が劣化されるためである。 Generally, the microstructure of TRIP steel must contain residual austenite, and technically, the residual austenite is substantially added to improve ductility by containing at a residual austenite fraction exceeding 0%. It means steel that plays an important role. Therefore, as an embodiment of the present invention, in the case of TRIP steel, it is preferable that the fine structure has a retained austenite fraction exceeding 0% and the balance is bainite. However, it is more preferable that the microstructure of the present invention has a retained austenite fraction in the range of 10 to 40% and the balance is bainite. This is because when the fraction of retained austenite is less than 10%, the level is similar to that of known TRIP steel, and when it exceeds 40%, the stability of retained austenite is deteriorated.
また、ベイナイト及び残留オーステナイトの平均ラス（ｌａｔｈ）幅を、各々０．３μｍ以下に限定する。平均ラス幅が０．３μｍを超える場合、組織が微細にできないため、強度の確保が困難である。よって、前記ベイナイトと残留オーステナイトの平均ラス幅は、０．３μｍ以下の範囲に限定することが好ましい。 Further, the average lath width of bainite and retained austenite is limited to 0.3 μm or less. When the average lath width exceeds 0.3 μm, it is difficult to secure strength because the structure cannot be made fine. Therefore, the average lath width of the bainite and retained austenite is preferably limited to a range of 0.3 μm or less.
また、前記鋼板は降伏強度が７００ＭＰａ以上で、引張強度は９８０ＭＰａ以上であり、延伸率は２０％以上である。降伏強度が７００ＭＰａ未満である場合には自動車の構造部材として使用することには不適合であって、引張強度が９８０ＭＰａ以上である場合を超高強度鋼とする。また、延伸率が２０％以上でないと延伸率が一定水準にないため、加工性向上の効果も期待することができない。 The steel sheet has a yield strength of 700 MPa or more, a tensile strength of 980 MPa or more, and a draw ratio of 20% or more. When the yield strength is less than 700 MPa, it is unsuitable for use as a structural member of an automobile, and the case where the tensile strength is 980 MPa or more is designated as ultra high strength steel. Moreover, since the stretch ratio is not at a certain level unless the stretch ratio is 20% or more, the effect of improving workability cannot be expected.
前記鋼板は、熱延鋼板、冷延鋼板、亜鉛メッキ鋼板及び亜鉛メッキ合金化鋼板のうち１種である。 The said steel plate is 1 type among a hot-rolled steel plate, a cold-rolled steel plate, a galvanized steel plate, and a galvanized alloyed steel plate.
以下、本発明の熱延鋼板、冷延鋼板、亜鉛メッキ鋼板、亜鉛メッキ合金化鋼板の製造方法に対して説明する。 Hereinafter, the manufacturing method of the hot rolled steel sheet, cold rolled steel sheet, galvanized steel sheet, and galvanized alloyed steel sheet of the present invention will be described.
本発明に適する熱延鋼板の製造方法によるスラブを通常の方法で再加熱して熱間圧延を行う。 The slab by the manufacturing method of the hot rolled steel sheet suitable for this invention is reheated by a normal method, and hot rolling is performed.
熱延鋼板の場合、Ａｒ３変態点以上にて熱間圧延を仕上げて、Ｍｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で巻取りする。 In the case of a hot-rolled steel sheet, hot rolling is finished at an Ar3 transformation point or higher, and winding is performed in a range of Ms (martensitic transformation start temperature) to 450 ° C.
熱間圧延の仕上げ温度をＡｒ３変態点以上に限定することは、二相域圧延が行われることを防止するためである。Ａｒ３変態点未満である場合には、二相域圧延が行われて不均一結晶粒組織が発生する恐れがある。よって、熱間圧延の仕上げ温度は、Ａｒ３変態点以上に限定することが好ましい。 The reason for limiting the hot rolling finishing temperature to the Ar3 transformation point or higher is to prevent the two-phase region rolling from being performed. When the temperature is less than the Ar3 transformation point, two-phase rolling may be performed to generate a non-uniform grain structure. Therefore, it is preferable to limit the hot rolling finishing temperature to the Ar3 transformation point or higher.
また、巻取り温度がＭｓ未満である場合には、マルテンサイトに変態することによって、ベイナイトと残留オーステナイト組織を得ることが困難になる恐れがある。一方、巻取り温度が４５０℃を超えて巻取りを行う場合、パーライト変態が生じる恐れがある。よって、巻取り温度は、Ｍｓ〜４５０℃の範囲に限定することが好ましい。 When the coiling temperature is lower than Ms, transformation to martensite may make it difficult to obtain a bainite and residual austenite structure. On the other hand, when the winding temperature exceeds 450 ° C., pearlite transformation may occur. Therefore, the winding temperature is preferably limited to a range of Ms to 450 ° C.
冷延鋼板の場合、重量％で、Ｃ：０．５〜１．０％、Ｍｎ：０．０１〜２．０％、Ｓｉ：１．０〜２．０％、Ａｌ：０．０１〜１．０％，Ｓ：０．０１２％以下、残部Ｆｅ、及びその他の不可避な不純物を含む鋼を熱間圧延して、Ａｒ３変態点以上にて熱間圧延を仕上げて６５０℃以下にて巻取りを行い、冷間圧延した後、Ａｃ３温度以上にて焼鈍を行った後に冷却した後、Ｍｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で過時効処理を行う。 In the case of a cold-rolled steel sheet, by weight, C: 0.5 to 1.0%, Mn: 0.01 to 2.0%, Si: 1.0 to 2.0%, Al: 0.01 to 1 0.0%, S: 0.012% or less, remaining Fe, and other steels containing other inevitable impurities are hot-rolled and hot-rolled at the Ar3 transformation point or higher and wound at 650 ° C or lower. After performing cold rolling, annealing is performed at the Ac3 temperature or higher, cooling is performed, and then an overaging treatment is performed in a range of Ms (martensitic transformation start temperature) to 450 ° C.
熱間圧延の仕上げ温度をＡｒ３変態点以上に限定するのは、二相域圧延が行われることを防止するためである。Ａｒ３変態点未満である場合には、二相域圧延が行われて不均一結晶粒組織が発生する恐れがある。よって、熱間圧延の仕上げ温度は、Ａｒ３変態点以上に限定することが好ましい。 The reason why the finishing temperature of hot rolling is limited to the Ar3 transformation point or more is to prevent the two-phase region rolling from being performed. When the temperature is less than the Ar3 transformation point, two-phase rolling may be performed to generate a non-uniform grain structure. Therefore, it is preferable to limit the hot rolling finishing temperature to the Ar3 transformation point or higher.
巻取り温度が６５０℃を超える場合、表面脱炭が非常に激しく生じる恐れがある。但し、下限を制限しない理由は、冷間圧延の後に焼鈍によって新たな組織を具現するので冷間圧延の前の微細組織は最終組織に大きな影響を及ぼすことができないためである。 When the coiling temperature exceeds 650 ° C., surface decarburization may occur very vigorously. However, the reason why the lower limit is not limited is that since a new structure is realized by annealing after cold rolling, the fine structure before cold rolling cannot greatly affect the final structure.
過時効処理温度がＭｓ未満である場合にはマルテンサイトに変態することによって、ベイナイトと残留オーステナイト組織を得ることが困難になる恐れがある。一方、４５０℃を超えて巻取りを行う場合、パーライト変態が生じる恐れがある。よって、時効処理温度は、Ｍｓ〜４５０℃に限定することが好ましい。 If the overaging temperature is less than Ms, transformation to martensite may make it difficult to obtain a bainite and residual austenite structure. On the other hand, when winding is performed at a temperature exceeding 450 ° C., pearlite transformation may occur. Therefore, the aging treatment temperature is preferably limited to Ms to 450 ° C.
そして、前記冷延鋼板をメッキ槽を通過させて亜鉛メッキ鋼板を製造して、前記亜鉛メッキ鋼板を合金化熱処理して亜鉛メッキ合金化鋼板を製造することができる。 Then, the cold-rolled steel sheet can be passed through a plating tank to produce a galvanized steel sheet, and the galvanized steel sheet can be alloyed and heat-treated to produce a galvanized alloyed steel sheet.
以下、実施例を介して本発明をより詳しく説明する。 Hereinafter, the present invention will be described in more detail through examples.
（実施例）
下記、表１に示す成分系を有する鋼塊を厚さ９０ｍｍ、幅１７５ｍｍに製造して１２００℃にて１時間再加熱した後、熱延して厚さが３ｍｍになるように熱間圧延を行った。一部は、熱延鋼板の試料として用いて、残りは冷延鋼板用の試料として、冷間圧延の後に焼鈍熱処理を行った。
(Example)
A steel ingot having the component system shown in Table 1 below is manufactured to a thickness of 90 mm and a width of 175 mm, reheated at 1200 ° C. for 1 hour, and then hot-rolled to a thickness of 3 mm. went. A part was used as a sample of a hot-rolled steel sheet, and the rest was used as a sample for a cold-rolled steel sheet, and an annealing heat treatment was performed after cold rolling.
熱延鋼板用の試料の熱間圧延仕上げ温度は、Ａｒ３変態点以上とし、冷却の後に４００℃に予め過熱された炉に装入して１時間維持した後、炉冷させて熱延巻取りを行った。 The hot rolling finish temperature of the sample for hot-rolled steel sheet is not less than the Ar3 transformation point, and after cooling, it is charged in a furnace preheated to 400 ° C. and maintained for 1 hour, and then cooled in the furnace and hot-rolled. Went.
一方、冷延鋼板用の試料は、下記表３に示すように、６００℃に予め加熱された炉に装入して１時間維持した後、炉冷を行うことによって熱延巻取りを行って、熱間圧延された板材をさらに６０％に冷間圧延して１．２ｍｍに仕上げた後、８４０〜９００℃にて焼鈍を行って、４００℃にてベイナイト変態を行った。 On the other hand, as shown in Table 3 below, the sample for cold-rolled steel sheet was charged into a furnace preheated to 600 ° C. and maintained for 1 hour, and then subjected to hot-rolling by performing furnace cooling. The hot-rolled plate material was further cold-rolled to 60% and finished to 1.2 mm, and then annealed at 840 to 900 ° C., and bainite transformation was performed at 400 ° C.
前記表１に示した成分系の鋼種を４００、６００℃にて巻取りした後、熱延鋼板の降伏強度と引張強度及び延伸率を測定した結果を表２に示し、前記表１に示した成分系の鋼種を６００℃にて巻取りした後に焼鈍し、時効処理した冷延鋼板の降伏強度、引張強度及び延伸率を測定した結果を、下記表３に示した。 Table 2 shows the results of measuring the yield strength, tensile strength, and draw ratio of the hot-rolled steel sheet after winding the component steel types shown in Table 1 at 400 ° C and 600 ° C. Table 3 below shows the results of measuring the yield strength, tensile strength, and stretch ratio of the cold-rolled steel sheet that was annealed after coiling the component steel types at 600 ° C. and then aged.
前記表２に示すように、発明例１〜８は、４００℃にて巻取り、降伏強度は７００ＭＰａ以上であって、引張強度は９８０ＭＰａ以上であり、延伸率は２０％以上であることを確認することができる。一方、比較例１〜１３は、６００℃にて巻取りを行って、降伏強度は７００ＭＰａ未満であるか、または引張強度は９８０ＭＰａ未満もしくは、延伸率が２０％未満であるように要求される物性の一部もしくは全部を満たさないことを確認することができる。 As shown in Table 2, Invention Examples 1 to 8 were wound at 400 ° C., yield strength was 700 MPa or more, tensile strength was 980 MPa or more, and stretch rate was 20% or more. can do. On the other hand, Comparative Examples 1 to 13 are wound up at 600 ° C., and yield strength is less than 700 MPa, or tensile strength is less than 980 MPa, or stretch properties are required to be less than 20%. It can be confirmed that some or all of the above are not satisfied.
前記表３に示すように、発明例９〜１７は、８４０〜９００℃にてオーステナイト単相域熱処理を行って、降伏強度は７００ＭＰａ以上であって、引張強度は９８０ＭＰａ以上であり、延伸率は２０％以上であることを確認することができる。 As shown in Table 3, Invention Examples 9 to 17 were subjected to austenite single-phase region heat treatment at 840 to 900 ° C., yield strength was 700 MPa or more, tensile strength was 980 MPa or more, and stretch ratio was It can be confirmed that it is 20% or more.
一方、Ａｌの含量が１．０重量％を超えてＣの含量が０．５重量％未満である比較例１５〜２０は、降伏強度が７００ＭＰａ未満であり、延伸率が２０％未満であることを確認することができる。これは、Ｃの含量が低いことによって適当な強度を確保することができず、Ｃが低くてＡｌが高いことに起因してＡｃ３温度が上昇することによる二相域熱処理が行われることによるフェライト生成が原因であることがわかる。 On the other hand, in Comparative Examples 15 to 20 in which the Al content exceeds 1.0% by weight and the C content is less than 0.5% by weight, the yield strength is less than 700 MPa and the stretch ratio is less than 20%. Can be confirmed. This is because ferrite cannot be secured due to the low content of C, and a two-phase heat treatment is performed by increasing the Ac3 temperature due to low C and high Al. It can be seen that the generation is the cause.
図１は、発明例１３の微細組織の写真であって、図１Ｂは、図１Ａを拡大した写真であり、相対的に暗い部分がベイナイト部分で明るい部分がオーステナイトである。各々の平均ラス幅が０．３μｍ以下に非常に微細な組織であることがわかる。 FIG. 1 is a photograph of the microstructure of Invention Example 13, and FIG. 1B is an enlarged view of FIG. 1A, in which a relatively dark portion is a bainite portion and a bright portion is austenite. It can be seen that each average lath width is a very fine structure of 0.3 μm or less.
一方に、図２は比較例１８の微細組織の写真であって、図２Ｂは、図２Ａを拡大した写真であり、相対的に暗い部分がベイナイト部分で、明るい部分がオーステナイトである。また、多角形フェライトが存在し、組織が相対的に非常に粗大であることがわかる。 On the other hand, FIG. 2 is a photograph of the microstructure of Comparative Example 18, and FIG. 2B is an enlarged photograph of FIG. 2A. The relatively dark part is the bainite part and the bright part is the austenite. Moreover, it turns out that polygonal ferrite exists and a structure | tissue is comparatively very coarse.
Claims (15)
Ａグループ：重量％で、０．５％以下のＭｏ及び１．０％以下のＣｒから選択される１種または２種、
Ｂグループ：重量％で、０．００５％以下のＢ及び０．０６％以下のＴｉから選択される１種または２種。 The high-strength and high-stretched steel sheet according to claim 1, wherein the steel sheet further includes at least one of group A and group B:
Group A: One or two selected from 0.5% or less of Mo and 1.0% or less of Cr by weight%,
Group B: One or two selected from B by 0.005% or less and Ti by 0.06% or less by weight%.
Ａｒ３変態点以上にて熱間圧延を仕上げて、
Ｍｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で巻取りして、
ベイナイト及び残留オーステナイトからなり、前記ベイナイトの平均ラス（ｌａｔｈ）幅が０．３μｍ以下であり、かつ前記残留オーステナイトの平均ラス幅が０．３μｍ以下である内部組織を有し、降伏強度が７００ＭＰａ以上であって、延伸率が２０％以上である鋼板を得る、
高強度高延伸熱延鋼板の製造方法。 C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: 0.01-2.0%, S: Hot rolling a slab composed of 0.012% or less, the remainder Fe, and other inevitable impurities,
Finish hot rolling above the Ar3 transformation point,
Winding in the range of Ms (Martensite transformation start temperature) to 450 ° C ,
It comprises bainite and retained austenite, and has an internal structure in which the average lath width of the bainite is 0.3 μm or less and the average lath width of the retained austenite is 0.3 μm or less, and the yield strength is 700 MPa or more. And obtaining a steel sheet having a draw ratio of 20% or more ,
A method for producing a high-strength, high-stretch hot-rolled steel sheet.
Ａグループ：重量％で、０．５％以下のＭｏ及び１．０％以下のＣｒから選択される１種または２種、
Ｂグループ：重量％で、０．００５％以下のＢ及び０．０６％以下のＴｉから選択される１種または２種。 The method for producing a high-strength, high-stretched hot-rolled steel sheet according to claim 8, wherein the hot-rolled steel sheet further includes at least one group of A group and B group:
Group A: One or two selected from 0.5% or less of Mo and 1.0% or less of Cr by weight%,
Group B: One or two selected from B by 0.005% or less and Ti by 0.06% or less by weight%.
前記の熱間圧延された熱延鋼板を６５０℃以下にて巻取りする巻取りステップと、
前記熱延鋼板を冷間圧延した後、Ａｃ３温度以上にて焼鈍した後に冷却するステップと、
前記冷延鋼板をＭｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で過時効処理する時効処理ステップと、
を含んで、ベイナイト及び残留オーステナイトからなり、前記ベイナイトの平均ラス（ｌａｔｈ）幅が０．３μｍ以下であり、かつ前記残留オーステナイトの平均ラス幅が０．３μｍ以下である内部組織を有し、降伏強度が７００ＭＰａ以上であって、延伸率が２０％以上である鋼板を得る、
高強度高延伸冷延鋼板の製造方法。 C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: 0.01-1.0%, S: 0.012% or less, and the balance Fe, and slabs made of other unavoidable impurities and hot rolling, and hot rolling step of finishing the hot rolling at least Ar3 transformation point,
A winding step for winding the hot-rolled hot-rolled steel sheet at 650 ° C. or lower;
After cold rolling the hot-rolled steel sheet, cooling after annealing at an Ac3 temperature or higher; and
An aging treatment step of over-aging the cold-rolled steel sheet in a range of Ms (martensitic transformation start temperature) to 450 ° C;
Nde-containing consists bainite and residual austenite, the average lath (LATH) width of the bainite is at 0.3μm or less, and an average lath width of the residual austenite has an internal tissue is 0.3μm or less, the yield Obtaining a steel plate having a strength of 700 MPa or more and a draw ratio of 20% or more;
A method for producing a high-strength, high-stretched cold-rolled steel sheet.
Ａグループ：重量％で、０．５％以下のＭｏ及び１．０％以下のＣｒから選択される１種または２種、
Ｂグループ：重量％で、０．００５％以下のＢ及び０．０６％以下のＴｉから選択される１種または２種。 The method for producing a high-strength, high-stretched cold-rolled steel sheet according to claim 10, wherein the cold-rolled steel sheet further includes at least one group of Group A and Group B:
Group A: One or two selected from 0.5% or less of Mo and 1.0% or less of Cr by weight%,
Group B: One or two selected from B by 0.005% or less and Ti by 0.06% or less by weight%.
前記の熱間圧延された熱延鋼板を６５０℃以下にて巻取りする巻取りステップと、
前記熱延鋼板を冷間圧延した後、Ａｃ３温度以上にて焼鈍した後に冷却するステップと、
前記冷延鋼板をＭｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で過時効処理する時効処理ステップと、
前記時効処理した冷延鋼板をメッキ槽を通過させて鋼板表面にメッキ層を形成させるステップと、
を含んで、ベイナイト及び残留オーステナイトからなり、前記ベイナイトの平均ラス（ｌａｔｈ）幅が０．３μｍ以下であり、かつ前記残留オーステナイトの平均ラス幅が０．３μｍ以下である内部組織を有し、降伏強度が７００ＭＰａ以上であって、延伸率が２０％以上である鋼板を得る、
高強度高延伸亜鉛メッキ鋼板の製造方法。 C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: 0.01-1.0%, S: 0.012% or less, and the balance Fe, and slabs made of other unavoidable impurities and hot rolling, and hot rolling step of finishing the hot rolling at least Ar3 transformation point,
A winding step for winding the hot-rolled hot-rolled steel sheet at 650 ° C. or lower;
After cold rolling the hot-rolled steel sheet, cooling after annealing at an Ac3 temperature or higher; and
An aging treatment step of over-aging the cold-rolled steel sheet in a range of Ms (martensitic transformation start temperature) to 450 ° C;
Passing the aged cold-rolled steel sheet through a plating tank to form a plating layer on the steel sheet surface; and
Nde-containing consists bainite and residual austenite, the average lath (LATH) width of the bainite is at 0.3μm or less, and an average lath width of the residual austenite has an internal tissue is 0.3μm or less, the yield Obtaining a steel plate having a strength of 700 MPa or more and a draw ratio of 20% or more;
A method for producing a high-strength, highly-stretched galvanized steel sheet.
Ａグループ：重量％で、０．５％以下のＭｏ及び１．０％以下のＣｒから選択される１種または２種、
Ｂグループ：重量％で、０．００５％以下のＢ及び０．０６％以下のＴｉから選択される１種または２種。 The method for producing a high-strength, high-stretched galvanized steel sheet according to claim 12, wherein the galvanized steel sheet further includes at least one group of Group A and Group B:
Group A: One or two selected from 0.5% or less of Mo and 1.0% or less of Cr by weight%,
Group B: One or two selected from B by 0.005% or less and Ti by 0.06% or less by weight%.
前記の熱間圧延された熱延鋼板を６５０℃以下にて巻取りする巻取りステップと、
前記熱延鋼板を冷間圧延した後、Ａｃ３温度以上にて焼鈍した後に冷却するステップと、
前記冷延鋼板をＭｓ（マルテンサイト変態開始温度）〜４５０℃の範囲で過時効処理する時効処理ステップと、
前記時効処理した冷延鋼板をメッキ槽を通過させて鋼板表面にメッキ層を形成させるステップと、
前記メッキ層を形成した亜鉛メッキ鋼板を合金化熱処理して亜鉛メッキ合金化鋼板を製造するステップと、
を含んで、ベイナイト及び残留オーステナイトからなり、前記ベイナイトの平均ラス（ｌａｔｈ）幅が０．３μｍ以下であり、かつ前記残留オーステナイトの平均ラス幅が０．３μｍ以下である内部組織を有し、降伏強度が７００ＭＰａ以上であって、延伸率が２０％以上である鋼板を得る、
高強度高延伸亜鉛メッキ合金化鋼板の製造方法。 C: 0.5-1.0%, Mn: 0.01-2.0%, Si: 1.0-2.0%, Al: 0.01-1.0%, S: 0.012% or less, and the balance Fe, and slabs made of other unavoidable impurities and hot rolling, and hot rolling step of finishing the hot rolling at least Ar3 transformation point,
A winding step for winding the hot-rolled hot-rolled steel sheet at 650 ° C. or lower;
After cold rolling the hot-rolled steel sheet, cooling after annealing at an Ac3 temperature or higher; and
An aging treatment step of over-aging the cold-rolled steel sheet in a range of Ms (martensitic transformation start temperature) to 450 ° C;
Passing the aged cold-rolled steel sheet through a plating tank to form a plating layer on the steel sheet surface; and
A step of producing a galvanized alloyed steel sheet by alloying heat treatment of the galvanized steel sheet on which the plated layer is formed;
Nde-containing consists bainite and residual austenite, the average lath (LATH) width of the bainite is at 0.3μm or less, and an average lath width of the residual austenite has an internal tissue is 0.3μm or less, the yield Obtaining a steel plate having a strength of 700 MPa or more and a draw ratio of 20% or more;
A method for producing high-strength, highly-stretched galvanized alloyed steel sheets.
Ａグループ：重量％で、０．５％以下のＭｏ及び１．０％以下のＣｒから選択される１種または２種、
Ｂグループ：重量％で、０．００５％以下のＢ及び０．０６％以下のＴｉから選択される１種または２種。 The method for producing a high-strength, high-stretched galvanized steel sheet according to claim 14, wherein the galvanized alloyed steel sheet further includes at least one of group A and group B:
Group A: One or two selected from 0.5% or less of Mo and 1.0% or less of Cr by weight%,
Group B: One or two selected from B by 0.005% or less and Ti by 0.06% or less by weight%.
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| KR10-2008-0133563 | 2008-12-24 | ||
| KR20080133563A KR101091294B1 (en) | 2008-12-24 | 2008-12-24 | Steel Sheet With High Strength And Elongation And Method For Manufacturing Hot-Rolled Steel Sheet, Cold-Rolled Steel Sheet, Galvanized Steel Sheet And Galvannealed Steel Sheet With High Strength And Elongation |
| PCT/KR2009/002098 WO2010074370A1 (en) | 2008-12-24 | 2009-04-22 | High-strength elongation steel sheet, hot-rolled steel sheet, cold-rolled steel sheet, zinc-coated steel sheet, and method for manufacturing alloyed zinc-coated steel sheet |
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| JP2011543380A Active JP5487215B2 (en) | 2008-12-24 | 2009-04-22 | Manufacturing method of high-strength, high-stretched steel sheet, hot-rolled steel sheet, cold-rolled steel sheet, galvanized steel sheet and galvanized alloyed steel sheet |
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| JP5756774B2 (en) * | 2012-03-09 | 2015-07-29 | 株式会社神戸製鋼所 | Steel sheet for hot pressing, press-formed product, and method for producing press-formed product |
| JP5756773B2 (en) * | 2012-03-09 | 2015-07-29 | 株式会社神戸製鋼所 | Steel sheet for hot pressing, press-formed product, and method for producing press-formed product |
| JP5764549B2 (en) * | 2012-03-29 | 2015-08-19 | 株式会社神戸製鋼所 | High-strength cold-rolled steel sheet, high-strength hot-dip galvanized steel sheet, high-strength galvannealed steel sheet excellent in formability and shape freezing property, and methods for producing them |
| CN104662181B (en) * | 2012-09-21 | 2016-10-19 | 日立金属株式会社 | The manufacture method of Maraging steel coiled material |
| PL2987884T3 (en) * | 2013-04-15 | 2019-07-31 | Nippon Steel & Sumitomo Metal Corporation | Hot-rolled steel sheet |
| KR101560940B1 (en) * | 2013-12-24 | 2015-10-15 | 주식회사 포스코 | Light weight steel sheet having excellent strength and ductility |
| WO2016108443A1 (en) * | 2014-12-30 | 2016-07-07 | 한국기계연구원 | High-strength steel plate having excellent combination of strength and ductility, and manufacturing method therefor |
| CN108396260B (en) * | 2017-02-05 | 2020-01-07 | 鞍钢股份有限公司 | High-strength high-hole-expansion-performance galvanized steel sheet and manufacturing method thereof |
| CN108396237B (en) * | 2017-02-05 | 2020-01-07 | 鞍钢股份有限公司 | High-plasticity cold-rolled sheet and production method thereof |
| KR101940919B1 (en) | 2017-08-08 | 2019-01-22 | 주식회사 포스코 | Hot rolled steel sheet having excellent strength and elongation and method of manufacturing the same |
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| JPS5842246B2 (en) * | 1979-04-28 | 1983-09-19 | 日新製鋼株式会社 | Method for manufacturing high-strength steel strip with composite structure |
| JPS60184634A (en) * | 1984-03-01 | 1985-09-20 | Nippon Steel Corp | Manufacture of high tension steel sheet having superior ductility and workability |
| JPS61217529A (en) * | 1985-03-22 | 1986-09-27 | Nippon Steel Corp | Manufacture of high strength steel sheet superior in ductility |
| JPS624832A (en) * | 1985-06-28 | 1987-01-10 | Nippon Steel Corp | Production of high-carbon steel strip having excellent ductility |
| US6159312A (en) * | 1997-12-19 | 2000-12-12 | Exxonmobil Upstream Research Company | Ultra-high strength triple phase steels with excellent cryogenic temperature toughness |
| JP2000144311A (en) * | 1998-11-13 | 2000-05-26 | Sumitomo Metal Ind Ltd | High carbon thin steel sheet |
| CN1295353C (en) * | 2000-04-07 | 2007-01-17 | 川崎制铁株式会社 | Production method of cold-rolled plate and hot-dip galvanized sheet steel with good strain-aged hardening characteristics |
| JP4188582B2 (en) * | 2001-02-09 | 2008-11-26 | 株式会社神戸製鋼所 | High-strength steel sheet with excellent workability and method for producing the same |
| CN1276987C (en) * | 2001-10-19 | 2006-09-27 | 住友金属工业株式会社 | Thin steel plate with excellent workability and forming precision and producing process thereof |
| JP4068950B2 (en) * | 2002-12-06 | 2008-03-26 | 株式会社神戸製鋼所 | High-strength steel sheet, warm-working method, and warm-worked high-strength member or parts |
| JP4412727B2 (en) * | 2004-01-09 | 2010-02-10 | 株式会社神戸製鋼所 | Super high strength steel sheet with excellent hydrogen embrittlement resistance and method for producing the same |
| JP5011846B2 (en) * | 2005-06-29 | 2012-08-29 | Ｊｆｅスチール株式会社 | High carbon hot rolled steel sheet and manufacturing method thereof |
| CN101346482B (en) * | 2005-12-26 | 2011-11-16 | Posco公司 | Carbon steel sheet superior in formability and manufacturing method thereof |
| CN101008066B (en) * | 2006-01-27 | 2010-05-12 | 宝山钢铁股份有限公司 | Hot rolling martensite steel plate with tensile strength higher than 1000Mpa and its production method |
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| CN102325916A (en) | 2012-01-18 |
| CN102325916B (en) | 2013-07-17 |
| WO2010074370A1 (en) | 2010-07-01 |
| JP2012514130A (en) | 2012-06-21 |
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| KR101091294B1 (en) | 2011-12-07 |
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