JP2015113504A - High strength hot-dip galvanized steel sheet excellent in processability and method for manufacturing the same - Google Patents
High strength hot-dip galvanized steel sheet excellent in processability and method for manufacturing the same Download PDFInfo
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- 229910001335 Galvanized steel Inorganic materials 0.000 title claims abstract description 30
- 239000008397 galvanized steel Substances 0.000 title claims abstract description 30
- 238000004519 manufacturing process Methods 0.000 title claims description 13
- 238000000034 method Methods 0.000 title description 27
- 229910000831 Steel Inorganic materials 0.000 claims abstract description 78
- 239000010959 steel Substances 0.000 claims abstract description 78
- 229910001566 austenite Inorganic materials 0.000 claims abstract description 61
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 53
- 230000000717 retained effect Effects 0.000 claims abstract description 27
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 21
- 239000000203 mixture Substances 0.000 claims abstract description 8
- 239000012535 impurity Substances 0.000 claims abstract description 7
- 238000001816 cooling Methods 0.000 claims description 37
- 238000000137 annealing Methods 0.000 claims description 33
- 238000010438 heat treatment Methods 0.000 claims description 20
- 238000005097 cold rolling Methods 0.000 claims description 15
- 238000011282 treatment Methods 0.000 claims description 15
- 238000005246 galvanizing Methods 0.000 claims description 14
- 238000003303 reheating Methods 0.000 claims description 12
- 229910052759 nickel Inorganic materials 0.000 claims description 8
- 229910052802 copper Inorganic materials 0.000 claims description 7
- 229910052750 molybdenum Inorganic materials 0.000 claims description 7
- 229910052758 niobium Inorganic materials 0.000 claims description 7
- 229910052720 vanadium Inorganic materials 0.000 claims description 7
- 230000009467 reduction Effects 0.000 claims description 6
- 238000005275 alloying Methods 0.000 claims description 4
- 238000012545 processing Methods 0.000 claims description 4
- 238000005096 rolling process Methods 0.000 description 28
- 230000000694 effects Effects 0.000 description 18
- 230000008569 process Effects 0.000 description 18
- 230000007423 decrease Effects 0.000 description 10
- 229910001562 pearlite Inorganic materials 0.000 description 9
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 8
- 238000005728 strengthening Methods 0.000 description 8
- 230000009466 transformation Effects 0.000 description 8
- 229910001563 bainite Inorganic materials 0.000 description 7
- 230000015572 biosynthetic process Effects 0.000 description 7
- 238000005098 hot rolling Methods 0.000 description 7
- 238000012360 testing method Methods 0.000 description 7
- 238000007747 plating Methods 0.000 description 6
- 238000001953 recrystallisation Methods 0.000 description 6
- 229910052804 chromium Inorganic materials 0.000 description 5
- 238000009749 continuous casting Methods 0.000 description 4
- 239000000463 material Substances 0.000 description 4
- 239000010960 cold rolled steel Substances 0.000 description 3
- 238000005461 lubrication Methods 0.000 description 3
- 229910052748 manganese Inorganic materials 0.000 description 3
- 150000001247 metal acetylides Chemical class 0.000 description 3
- 238000005498 polishing Methods 0.000 description 3
- LFQSCWFLJHTTHZ-UHFFFAOYSA-N Ethanol Chemical compound CCO LFQSCWFLJHTTHZ-UHFFFAOYSA-N 0.000 description 2
- 229910000794 TRIP steel Inorganic materials 0.000 description 2
- 230000002411 adverse Effects 0.000 description 2
- 229910052799 carbon Inorganic materials 0.000 description 2
- 230000008859 change Effects 0.000 description 2
- 230000003749 cleanliness Effects 0.000 description 2
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- 238000007796 conventional method Methods 0.000 description 2
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- 238000010191 image analysis Methods 0.000 description 2
- 229910052742 iron Inorganic materials 0.000 description 2
- 239000010410 layer Substances 0.000 description 2
- 229910052698 phosphorus Inorganic materials 0.000 description 2
- 238000005554 pickling Methods 0.000 description 2
- 229920006395 saturated elastomer Polymers 0.000 description 2
- 238000005204 segregation Methods 0.000 description 2
- 230000006641 stabilisation Effects 0.000 description 2
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- 238000004781 supercooling Methods 0.000 description 2
- 238000004804 winding Methods 0.000 description 2
- GRYLNZFGIOXLOG-UHFFFAOYSA-N Nitric acid Chemical compound O[N+]([O-])=O GRYLNZFGIOXLOG-UHFFFAOYSA-N 0.000 description 1
- 229910045601 alloy Inorganic materials 0.000 description 1
- 239000000956 alloy Substances 0.000 description 1
- 238000005266 casting Methods 0.000 description 1
- 239000011248 coating agent Substances 0.000 description 1
- 238000000576 coating method Methods 0.000 description 1
- 239000002131 composite material Substances 0.000 description 1
- 238000013329 compounding Methods 0.000 description 1
- 238000012937 correction Methods 0.000 description 1
- 230000007797 corrosion Effects 0.000 description 1
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Abstract
Description
本発明は、主に自動車、電気等の産業分野で使用される部材として好適な加工性に優れた高強度溶融亜鉛めっき鋼板およびその製造方法に関するものである。 The present invention relates to a high-strength hot-dip galvanized steel sheet excellent in workability suitable as a member mainly used in industrial fields such as automobiles and electricity, and a method for producing the same.
近年、地球環境保全の見地から、自動車の燃費向上が重要な課題となっている。これに伴い、車体材料の高強度化により薄肉化を図り、車体そのものを軽量化しようとする動きが活発となってきている。 In recent years, improving the fuel efficiency of automobiles has become an important issue from the viewpoint of global environmental conservation. Along with this, there is an active movement to reduce the thickness of the vehicle body by increasing the strength of the vehicle body material and to reduce the weight of the vehicle body itself.
一般的に、鋼板を高強度化した場合には、鋼板の延性が低下し、成形時における加工性の低下を招く。このため、高強度と高加工性を兼ね備えた材料の開発が求められている。さらに、最近では、自動車への耐食性向上の要求の高まりも加味して、溶融亜鉛めっきを施した高強度鋼板への要望も高まっている。 In general, when the strength of a steel plate is increased, the ductility of the steel plate is reduced, and the workability during forming is reduced. For this reason, development of the material which has high intensity | strength and high workability is calculated | required. Furthermore, recently, taking into account the increasing demand for improving corrosion resistance of automobiles, there is an increasing demand for high-strength steel sheets subjected to hot dip galvanization.
このような要求に対して、これまでにフェライト、マルテンサイト二相鋼（DP鋼）や残留オーステナイトの変態誘起塑性を利用したTRIP鋼など、種々の複合組織型高強度溶融亜鉛めっき鋼板が開発されてきた。 In response to these requirements, various composite-structured high-strength hot-dip galvanized steel sheets such as ferrite, martensite duplex steel (DP steel), and TRIP steel using transformation-induced plasticity of retained austenite have been developed. I came.
例えば、特許文献１には、質量％で、Ｃ:0.05〜0.15％、Si:0.3〜1.5％、Mn:1.5〜2.8％、Ｐ:0.03％以下、Ｓ:0.02％以下、Al:0.005〜0.5％、Ｎ:0.0060％以下、残部がFeおよび不可避的不純物からなり、さらに(Mn％)/(Ｃ％)≧15かつ(Si％)/(Ｃ％)≧4を満たし、フェライト相中に体積率で３〜20％のマルテンサイト相と残留オーステナイト相を含む成形性の良い高強度合金化溶融亜鉛めっき鋼板が提案されている。
すなわち、特許文献１は、多量のSiを添加することにより残留オーステナイト相（残留γ相）を確保して高延性を達成し、これによって、加工性に優れた合金化溶融亜鉛めっき鋼板を得ようとする技術である。
For example, Patent Document 1 includes mass%, C: 0.05 to 0.15%, Si: 0.3 to 1.5%, Mn: 1.5 to 2.8%, P: 0.03% or less, S: 0.02% or less, Al: 0.005 to 0.5. %, N: 0.0060% or less, the balance being Fe and inevitable impurities, further satisfying (Mn%) / (C%) ≧ 15 and (Si%) / (C%) ≧ 4, and volume in the ferrite phase A high-strength galvannealed steel sheet having a good formability containing a martensite phase and a retained austenite phase of 3 to 20% is proposed.
That is, Patent Document 1 secures a retained austenite phase (residual γ phase) by adding a large amount of Si to achieve high ductility, thereby obtaining an alloyed hot-dip galvanized steel sheet having excellent workability. Technology.
また、特許文献２には、伸びフランジ性に優れる溶融亜鉛めっき鋼板の製造方法として、焼鈍均熱後、溶融亜鉛めっき浴までの間にMs点以下まで強冷却し、生成したマルテンサイトを再加熱して焼き戻しマルテンサイトとすることにより、穴拡げ性を向上させる技術が開示されている。 Patent Document 2 discloses a method for producing a hot-dip galvanized steel sheet having excellent stretch flangeability, after annealing and soaking, to a hot dip galvanizing bath, strongly cooled below the Ms point, and reheats the martensite produced. Thus, a technique for improving hole expansibility by using tempered martensite is disclosed.
さらに、特許文献３には、質量％でＣ：0.05〜0.3％、Si：0.01〜2.5％、Mn：0.5〜3.5％、Ｐ：0.003〜0.100％以下、Ｓ：0.02％以下、Al：0.010〜1.5％を含有し、SiとAlの添加量の合計が0.5〜2.5％であり、残部が鉄および不可避的不純物からなり、組織は、面積率で、20％以上のフェライト相と10％以下（０％を含む）のマルテンサイト相と10％以上60％以下の焼戻しマルテンサイト相を有し、体積率で、３％以上10％以下の残留オーステナイト相を有し、かつ、残留オーステナイト相の平均結晶粒径を2.0μm以下とすることにより、優れた強度、延性および伸びフランジ性を有する高強度溶融亜鉛めっき鋼板が提案されている。 Further, in Patent Document 3, in mass%, C: 0.05 to 0.3%, Si: 0.01 to 2.5%, Mn: 0.5 to 3.5%, P: 0.003 to 0.100% or less, S: 0.02% or less, Al: 0.010 to Containing 1.5%, the total addition amount of Si and Al is 0.5-2.5%, the balance consists of iron and inevitable impurities, the structure is 20% or more ferrite phase and 10% or less (in area ratio) 0% (including 0%) martensite phase and 10% or more and 60% or less tempered martensite phase, 3% or more and 10% or less of retained austenite phase by volume, and average of retained austenite phase A high-strength hot-dip galvanized steel sheet having excellent strength, ductility, and stretch flangeability has been proposed by setting the crystal grain size to 2.0 μm or less.
しかしながら、特許文献１のようなDP鋼やTRIP鋼は、伸び特性には優れるものの、穴拡げ性が劣るという問題があった。ここに、穴拡げ性とは、加工穴部を拡張してフランジ成形させるときの加工性（伸びフランジ性）を示す指標で、伸び特性と共に高強度鋼板に要求される重要な特性である。 However, although DP steel and TRIP steel like patent document 1 are excellent in elongation characteristics, there was a problem that hole expansibility was inferior. Here, the hole expandability is an index indicating workability (stretch flangeability) when the processed hole portion is expanded and flange-formed, and is an important characteristic required for high-strength steel sheets together with the elongation characteristics.
また、特許文献２では、マルテンサイトを再加熱して焼戻しマルテンサイトにすることにより、穴拡げ性は向上するものの、延性が低下するという問題があった。 Moreover, in patent document 2, although the hole expansibility improved by reheating a martensite to tempered martensite, there existed a problem that ductility fell.
さらに、特許文献３では、良好な延性および伸びフランジ性が得られるものの、時として、鋼板の板厚方向の機械特性、特に硬さにばらつきが生じることがあり、これにより、所望の機械特性が安定して得られないという問題があった。 Furthermore, in Patent Document 3, although good ductility and stretch flangeability can be obtained, sometimes the mechanical properties in the thickness direction of the steel sheet, particularly the hardness, may vary. There was a problem that it could not be obtained stably.
本発明は、上記の問題を有利に解決するものであって、引張強さ（TS）：590MPa以上という高強度は言うまでもなく、延性および伸びフランジ性に優れ、かつ板厚方向の機械特性、特に硬さが均一な高強度溶融亜鉛めっき鋼板を、その製造方法と共に提供することを目的とする。 The present invention advantageously solves the above-mentioned problems, and it has excellent tensile strength (TS): 590 MPa or more, as well as excellent ductility and stretch flangeability, and mechanical properties in the thickness direction, particularly An object is to provide a high-strength hot-dip galvanized steel sheet having a uniform hardness together with its manufacturing method.
発明者らは、上記の問題を解決すべく、鋭意検討を重ねた。
その結果、冷間圧延に先立ち、熱延板を適正な条件で焼鈍することで、鋼板の板厚方向の組織が改質され、これによって、硬さを含む板厚方向の機械特性のばらつきが大幅に軽減されるとの知見を得た。
本発明は、上記の知見に立脚するものである。
The inventors made extensive studies to solve the above problems.
As a result, the structure in the thickness direction of the steel sheet is modified by annealing the hot-rolled sheet under appropriate conditions prior to cold rolling, thereby causing variations in mechanical properties in the thickness direction including hardness. The knowledge that it is greatly reduced was obtained.
The present invention is based on the above findings.
すなわち、本発明の要旨構成は次のとおりである。
１．質量％で、
Ｃ：0.05〜0.3％、
Si：0.01〜2.5％、
Mn：0.5〜3.5％、
Ｐ：0.003〜0.100％、
Ｓ：0.02％以下および
Al：0.010〜1.5％
を含有し、かつSiとAlの合計量が0.5〜2.5％であって、残部はFeおよび不可避的不純物からなり、
組織全体に対する面積率でフェライト相を20％以上、マルテンサイト相を10％以下（但し、０％を含む）、焼戻しマルテンサイト相を10〜60％含み、かつ組織全体に対する体積率で残留オーステナイト相を３〜10％含む組織を有し、
板厚方向の硬さばらつきΔHvが20以下であることを特徴とする加工性に優れた高強度溶融亜鉛めっき鋼板。
That is, the gist configuration of the present invention is as follows.
1. % By mass
C: 0.05-0.3%
Si: 0.01-2.5%
Mn: 0.5-3.5%
P: 0.003-0.100%
S: 0.02% or less and
Al: 0.010 to 1.5%
And the total amount of Si and Al is 0.5 to 2.5%, the balance consists of Fe and inevitable impurities,
20% or more of ferrite phase, 10% or less (including 0%) of martensite phase, 10-60% of tempered martensite phase, and retained austenite phase by volume ratio of the whole structure. Having a tissue containing 3 to 10%,
A high-strength hot-dip galvanized steel sheet excellent in workability, characterized by having a hardness variation ΔHv in the thickness direction of 20 or less.
２．前記鋼板が、さらに、質量％で、Cr：0.005〜2.00％、Mo：0.005〜2.00％、Ｖ：0.005〜2.00％、Ni：0.005〜2.00％およびCu：0.005〜2.00％のうちから選ばれる１種または２種以上を含有することを特徴とする前記１に記載の加工性に優れた高強度溶融亜鉛めっき鋼板。 2. The steel sheet is further selected by mass% from Cr: 0.005-2.00%, Mo: 0.005-2.00%, V: 0.005-2.00%, Ni: 0.005-2.00% and Cu: 0.005-2.00%. The high-strength hot-dip galvanized steel sheet having excellent workability as described in 1 above, comprising seeds or two or more kinds.
３．前記鋼板が、さらに、質量％で、Ti：0.01〜0.20％およびNb：0.01〜0.20％のうちから選ばれる１種または２種を含有することを特徴とする前記１または２に記載の加工性に優れた高強度溶融亜鉛めっき鋼板。 3. The workability according to 1 or 2 above, wherein the steel sheet further contains, by mass%, one or two selected from Ti: 0.01 to 0.20% and Nb: 0.01 to 0.20%. High strength hot-dip galvanized steel sheet.
４．前記鋼板が、さらに、質量％で、Ｂ：0.0002〜0.005％を含有することを特徴とする前記１〜３のいずれかに記載の加工性に優れた高強度溶融亜鉛めっき鋼板。 4). The high-strength hot-dip galvanized steel sheet having excellent workability as described in any one of 1 to 3 above, wherein the steel sheet further contains B: 0.0002 to 0.005% by mass%.
５．前記鋼板が、さらに、質量％で、Ca：0.001〜0.005％およびREM：0.001〜0.005％のうちから選ばれる１種または２種を含有することを特徴とする前記１〜４のいずれかに記載の加工性に優れた高強度溶融亜鉛めっき鋼板。 5. 5. The steel sheet according to any one of 1 to 4, wherein the steel sheet further contains one or two kinds selected from Ca: 0.001 to 0.005% and REM: 0.001 to 0.005% by mass%. High-strength hot-dip galvanized steel sheet with excellent workability.
６．前記溶融亜鉛めっきが、合金化溶融亜鉛めっきであることを特徴とする前記１〜５のいずれかに記載の加工性に優れた高強度溶融亜鉛めっき鋼板。 6). The high-strength hot-dip galvanized steel sheet having excellent workability according to any one of 1 to 5, wherein the hot-dip galvanizing is alloyed hot-dip galvanizing.
７．前記１〜５のいずれかに記載の加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法であって、
前記１〜５のいずれかに記載の成分組成を有するスラブを、熱間圧延し、500℃以上Ａc1点以下の温度域で１〜10時間保持する焼鈍処理を施したのち、圧下率が60％超となる冷間圧延を施し、
ついで、少なくとも500℃以上Ａc1点以下の温度域における平均加熱速度を10℃/s以上として750〜900℃の温度域まで加熱し、該温度域で10秒以上保持したのち、平均冷却速度を10℃/s以上として750℃から（Ms点−100℃）〜（Ms点−200℃）の温度域まで冷却する連続焼鈍を施し、
さらに、350〜600℃の温度域まで再加熱して10〜600秒保持したのち、鋼板表面に溶融亜鉛めっきを施すことを特徴とする加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法。
7). A method for producing a high-strength hot-dip galvanized steel sheet excellent in workability according to any one of 1 to 5,
The slab having the component composition described in any one of 1 to 5 above is hot-rolled and subjected to an annealing treatment that is held in a temperature range of 500 ° C. or higher and Ac 1 point or lower for 1 to 10 hours. % Cold rolling,
Next, the average heating rate in a temperature range of at least 500 ° C. to Ac 1 point is set to 10 ° C./s or more and heated to a temperature range of 750 to 900 ° C., held in the temperature range for 10 seconds or more, and then the average cooling rate is set. Apply continuous annealing to cool from 750 ° C to a temperature range of (Ms point-100 ° C) to (Ms point-200 ° C) as 10 ° C / s or more,
Furthermore, after reheating to the temperature range of 350-600 degreeC and hold | maintaining for 10-600 second, hot-dip galvanized steel plate excellent in workability characterized by performing hot dip galvanizing on the steel plate surface.
８．前記溶融亜鉛めっきを施した後、さらに、前記溶融亜鉛めっきの合金化処理を施すことを特徴とする前記７に記載の加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法。 8). 8. The method for producing a high-strength hot-dip galvanized steel sheet having excellent workability according to 7, wherein the hot-dip galvanizing is further performed, and then an alloying treatment of the hot-dip galvanizing is performed.
本発明によれば、引張強さ(TS)：590MPa以上であり、かつ高い延性と伸びフランジ性を有し、さらには板厚方向の硬さばらつきが小さい加工性に優れた高強度溶融亜鉛めっき鋼板を安定して得ることができる。
そして、本発明により得られる高強度溶融亜鉛めっき鋼板を、例えば自動車構造部材に適用することにより、自動車の軽量化と衝突安全性向上との両立が可能となり、自動車車体の高性能化に大きく寄与することができる。
According to the present invention, high-strength hot-dip galvanizing with a tensile strength (TS) of 590 MPa or more, high ductility and stretch flangeability, and excellent workability with little hardness variation in the thickness direction. A steel plate can be obtained stably.
And, by applying the high-strength hot-dip galvanized steel sheet obtained by the present invention to, for example, automobile structural members, it becomes possible to achieve both reduction in weight of automobiles and improvement in collision safety, which greatly contributes to improving the performance of automobile bodies. can do.
以下、本発明を具体的に説明する。
１）成分組成
まず、成分組成を前記の範囲に限定した理由について説明する。なお、各元素の含有量の単位は、特に断りがない限り質量％を意味するものとする。
Ｃ：0.05〜0.3％
Ｃは、オーステナイトを安定化させ、フェライト以外の相を生成しやすくするために必要な元素である。また、Ｃは、鋼板強度を上昇させるとともに、組織を複合化してTSとELのバランスを向上させる点でも、必要な元素である。Ｃ量が0.05％未満では、製造条件の最適化を図ったとしてもフェライト以外の相の確保が難しく、TSとELのバランスが低下する。一方、Ｃ量が0.3％を超えると、溶接部および熱影響部が硬化し、溶接部の機械特性が劣化する。従って、Ｃ量は0.05〜0.3％の範囲とする。好ましくは0.08〜0.15％の範囲である。
Hereinafter, the present invention will be specifically described.
1) Component composition First, the reason why the component composition is limited to the above range will be described. In addition, unless otherwise indicated, the unit of content of each element shall mean the mass%.
C: 0.05-0.3%
C is an element necessary for stabilizing austenite and facilitating generation of phases other than ferrite. C is a necessary element from the viewpoint of increasing the steel sheet strength and improving the balance between TS and EL by compounding the structure. If the C content is less than 0.05%, it is difficult to secure phases other than ferrite even if the production conditions are optimized, and the balance between TS and EL decreases. On the other hand, when the amount of C exceeds 0.3%, the welded portion and the heat affected zone are cured, and the mechanical properties of the welded portion are deteriorated. Accordingly, the C content is in the range of 0.05 to 0.3%. Preferably it is 0.08 to 0.15% of range.
Si：0.01〜2.5％
Siは、鋼の強化に有効な元素である。また、Siは、フェライト生成元素であり、オーステナイト相中へのＣの濃化を促進すると共に炭化物の生成を抑制することから、残留オーステナイトの生成を促進する働きを有する。このような効果を得るためには、Siを0.01％以上添加する必要がある。ただし、Siの過剰な添加は、延性や表面性状、溶接性を劣化させるので、上限は2.5％とする。好ましくは0.7％〜2.0％の範囲である。
Si: 0.01-2.5%
Si is an element effective for strengthening steel. Si is a ferrite-generating element and promotes the formation of retained austenite because it promotes the concentration of C in the austenite phase and suppresses the formation of carbides. In order to obtain such an effect, it is necessary to add Si by 0.01% or more. However, excessive addition of Si deteriorates ductility, surface properties, and weldability, so the upper limit is 2.5%. Preferably it is 0.7 to 2.0% of range.
Mn：0.5〜3.5％
Mnは、鋼の強化に有効な元素であり、焼戻しマルテンサイト相等の低温変態相の生成を促進する働きを有する。このような効果を得るためには、Mnを0.5％以上添加する必要がある。ただし、Mnが3.5％を超えて過剰に添加されると、第二相分率の過剰な増加や固溶強化によるフェライトの延性劣化が著しくなり、成形性が低下する。従って、Mn量は0.5〜3.5％の範囲とする。好ましくは1.5〜3.0％の範囲である。
Mn: 0.5-3.5%
Mn is an element effective for strengthening steel and has a function of promoting the generation of a low-temperature transformation phase such as a tempered martensite phase. In order to obtain such an effect, it is necessary to add 0.5% or more of Mn. However, if Mn is added excessively exceeding 3.5%, the ductile deterioration of ferrite due to excessive increase of the second phase fraction or solid solution strengthening becomes remarkable, and the formability deteriorates. Therefore, the Mn content is in the range of 0.5 to 3.5%. Preferably it is 1.5 to 3.0% of range.
Ｐ：0.003〜0.100％
Ｐは、鋼の強化に有効な元素であり、この効果は0.003％以上で得られる。しかしながら、Ｐが0.100％を超えて過剰に添加されると、粒界偏析により脆化を引き起こし、耐衝撃性を劣化させる。従って、Ｐ量は0.003〜0.100％の範囲とする。
P: 0.003-0.100%
P is an element effective for strengthening steel, and this effect is obtained at 0.003% or more. However, if P exceeds 0.100% and is added excessively, it causes embrittlement due to grain boundary segregation and degrades impact resistance. Therefore, the P content is in the range of 0.003 to 0.100%.
Ｓ：0.02％以下
Ｓは、MnSなどの介在物となって、耐衝撃特性の劣化や溶接部のメタルフローに沿った割れの原因になるので極力低い方が好ましいが、製造コストの面から0.02％以下とする。
S: 0.02% or less S is an inclusion such as MnS, which causes deterioration in impact resistance and cracks along the metal flow of the weld. % Or less.
Al：0.010〜1.5％
Alは、脱酸剤として作用して鋼の清浄度を改善する有用元素であり、通常、脱酸工程で添加される。このような効果を得るためには、Alは0.010％以上添加する必要がある。しかしながら、Alを過剰に添加すると、連続鋳造時における鋼片割れが発生する危険性が高まり、製造性が低下する。従って、Al量の上限は1.5％とする。
Al: 0.010 to 1.5%
Al is a useful element that acts as a deoxidizer to improve the cleanliness of steel and is usually added in the deoxidation step. In order to obtain such an effect, Al needs to be added by 0.010% or more. However, if Al is added excessively, the risk of steel piece cracking during continuous casting increases, and the productivity decreases. Therefore, the upper limit of the Al amount is 1.5%.
SiとAlの合計量：0.5〜2.5％
また、Alは、Siと同様にフェライト生成元素であり、オーステナイト相中へのＣの濃化を促進すると共に炭化物の生成を抑制することから、残留オーステナイト相の生成を促進する働きがある。このような効果は、AlとSiの合計量が0.5％未満では、十分に発現せず、所望の延性が得られない。一方、AlとSiの合計量が2.5％を超えると、鋼板中の介在物が増加し、延性を劣化させる。従って、AlとSiの合計量は0.5〜2.5％の範囲とする。好ましくは0.8〜2.0％の範囲である。
Total amount of Si and Al: 0.5-2.5%
Al, like Si, is a ferrite-forming element and promotes the formation of residual austenite phase because it promotes the concentration of C in the austenite phase and suppresses the formation of carbides. Such an effect is not sufficiently exhibited when the total amount of Al and Si is less than 0.5%, and desired ductility cannot be obtained. On the other hand, when the total amount of Al and Si exceeds 2.5%, inclusions in the steel sheet increase and ductility is deteriorated. Therefore, the total amount of Al and Si is in the range of 0.5 to 2.5%. Preferably it is 0.8 to 2.0% of range.
以上、基本成分について説明したが、本発明では、その他にも、以下に述べる成分を必要に応じて適宜含有させることができる。
Cr：0.005〜2.00％、Mo：0.005〜2.00％、Ｖ：0.005〜2.00％、Ni：0.005〜2.00％およびCu：0.005〜2.00％のうちから選ばれる１種または２種以上
Cr、Mo、Ｖ、NiおよびCuは、焼鈍温度からの冷却時にパーライト相の生成を抑制すると共に、低温変態相の生成を促進して鋼の強化に有効に働く。このような効果は、Cr、Mo、Ｖ、NiおよびCuのうちから選んだ少なくとも１種を0.005％以上含有させることで得ることができる。しかしながら、Cr、Mo、Ｖ、NiおよびCu量がそれぞれ2.00％を超えると、その効果は飽和し、コストアップの要因となる。従って、Cr、Mo、Ｖ、NiおよびCu量はそれぞれ0.005％〜2.00％の範囲とする。
The basic components have been described above, but in the present invention, other components described below can be appropriately contained as necessary.
One or more selected from Cr: 0.005-2.00%, Mo: 0.005-2.00%, V: 0.005-2.00%, Ni: 0.005-2.00% and Cu: 0.005-2.00%
Cr, Mo, V, Ni, and Cu effectively suppress the formation of a pearlite phase during cooling from the annealing temperature, and promote the formation of a low-temperature transformation phase to effectively work for strengthening the steel. Such an effect can be obtained by adding 0.005% or more of at least one selected from Cr, Mo, V, Ni and Cu. However, if the amount of Cr, Mo, V, Ni, and Cu exceeds 2.00%, the effect is saturated and causes an increase in cost. Accordingly, the Cr, Mo, V, Ni, and Cu contents are in the range of 0.005% to 2.00%, respectively.
Ti：0.01〜0.20％およびNb：0.01〜0.20％のうちから選ばれる１種または２種
TiおよびNbは炭窒化物を形成し、鋼を析出強化により高強度化する作用を有する。このような効果は、TiおよびNbをそれぞれ0.01％以上含有させることで得ることができる。一方、TiおよびNbをそれぞれ0.20％を超えて含有させても、過度に高強度化し、延性が低下する。従って、TiおよびNb量はそれぞれ0.01％〜0.20％の範囲とする。
One or two selected from Ti: 0.01-0.20% and Nb: 0.01-0.20%
Ti and Nb form carbonitrides and have the effect of strengthening steel by precipitation strengthening. Such an effect can be obtained by containing Ti and Nb in an amount of 0.01% or more. On the other hand, even if Ti and Nb are contained in amounts exceeding 0.20%, the strength is excessively increased and the ductility is lowered. Therefore, the Ti and Nb contents are in the range of 0.01% to 0.20%, respectively.
Ｂ：0.0002〜0.005％
Ｂは、オーステナイト相粒界からのフェライトの生成を抑制し、強度を上昇させる作用を有する。そのような効果は、Ｂを0.0002％以上含有させることで得ることができる。一方、Ｂ量が0.005％を超えると、その効果は飽和し、コストアップの要因となる。従って、Ｂ量は0.0002〜0.005％の範囲とする。
B: 0.0002 to 0.005%
B has the effect of suppressing the formation of ferrite from the austenite phase grain boundaries and increasing the strength. Such an effect can be obtained by containing B in an amount of 0.0002% or more. On the other hand, when the amount of B exceeds 0.005%, the effect is saturated and causes an increase in cost. Therefore, the B content is in the range of 0.0002 to 0.005%.
Ca：0.001〜0.005％およびREM：0.001〜0.005％のうちから選ばれる１種または２種
CaおよびREMはいずれも、硫化物の形態制御により加工性を改善する効果を有しており、必要に応じてCaおよびREMをそれぞれ0.001％以上含有させることができる。しかしながら、CaおよびREMの過剰な添加は清浄度に悪影響を及ぼす恐れがあるため、CaおよびREM量はそれぞれ0.005％以下とする。
One or two selected from Ca: 0.001 to 0.005% and REM: 0.001 to 0.005%
Both Ca and REM have the effect of improving workability by controlling the morphology of sulfides, and if necessary, Ca and REM can be contained in an amount of 0.001% or more, respectively. However, excessive addition of Ca and REM may adversely affect cleanliness, so the Ca and REM amounts are each 0.005% or less.
本発明の鋼板において、上記以外の成分は、Feおよび不可避的不純物である。ただし、Ｎは、鋼板の加工性等の劣化させない範囲として、0.01％以下の含有であれば許容できる。 In the steel sheet of the present invention, components other than those described above are Fe and inevitable impurities. However, N is acceptable as long as the content is 0.01% or less as a range in which the workability of the steel sheet is not deteriorated.
２）鋼組織
次に、本発明の鋼板における鋼組織を前記の範囲に限定した理由を説明する。
フェライト相：組織全体に対する面積率で20％以上
フェライト相が、組織全体に対する面積率で20％未満であると、TSとELのバランスが低下する。このため、フェライト相は組織全体に対する面積率で20％以上とする。好ましくは50％以上である。
2) Steel structure Next, the reason which limited the steel structure in the steel plate of this invention to the said range is demonstrated.
Ferrite phase: 20% or more in area ratio relative to the entire structure If the ferrite phase is less than 20% in area ratio relative to the entire structure, the balance between TS and EL decreases. For this reason, the ferrite phase is 20% or more in terms of the area ratio with respect to the entire structure. Preferably it is 50% or more.
マルテンサイト相：組織全体に対する面積率で10％以下（但し、０％を含む）
マルテンサイト相は、鋼の高強度化には有効に働くが、組織全体に対する面積率で10％を超えて過剰に存在すると、λ（穴拡げ率）を大幅に低下させる。従って、マルテンサイト相は組織全体に対する面積率で10％以下とする。なお、マルテンサイト相を全く含まない、すなわち、マルテンサイト相が組織全体に対する面積率で０％であっても、本発明の効果には影響を及ぼさず、問題はない。
Martensite phase: 10% or less in area ratio to the whole structure (including 0%)
The martensite phase works effectively to increase the strength of steel, but if it is present in excess of 10% in terms of the area ratio relative to the entire structure, λ (hole expansion rate) is greatly reduced. Therefore, the martensite phase is 10% or less in terms of the area ratio with respect to the entire structure. Even if the martensite phase is not included at all, that is, the area ratio of the martensite phase to the entire structure is 0%, the effect of the present invention is not affected and there is no problem.
焼戻しマルテンサイト相：組織全体に対する面積率で10〜60％
焼戻しマルテンサイト相は鋼の強化に有効に働く。また、焼戻しマルテンサイト相は、マルテンサイト相に比べて穴拡げ性への悪影響が小さく、穴拡げ性の大幅な低下を招くことなしに、強度を確保することができる有効な相である。ここに、焼戻しマルテンサイト相が組織全体に対する面積率で10％未満では、強度確保が困難となる。一方、60％を超えると、TSとELのバランスが低下する。従って、焼戻しマルテンサイト相の組織全体に対する面積率は10〜60％の範囲とする。好ましくは20〜50％の範囲である。
Tempered martensite phase: 10-60% in area ratio to the whole structure
The tempered martensite phase works effectively to strengthen the steel. Further, the tempered martensite phase is an effective phase that has less adverse effect on the hole expandability than the martensite phase and can ensure strength without causing a significant decrease in the hole expandability. Here, if the tempered martensite phase is less than 10% in terms of the area ratio relative to the entire structure, it is difficult to ensure strength. On the other hand, if it exceeds 60%, the balance between TS and EL decreases. Accordingly, the area ratio of the tempered martensite phase to the entire structure is in the range of 10 to 60%. Preferably it is 20 to 50% of range.
残留オーステナイト相：組織全体に対する体積率で３〜10％、
残留オーステナイト相は、鋼の強化に寄与するだけでなく、鋼のTSとELのバランスの向上に有効に働く。このような効果は、残留オーステナイト相が組織全体に対する体積率で３％以上の場合に得られる。一方、残留オーステナイト相は、加工時にマルテンサイトに変態し、穴拡げ性を低下させるが、組織全体に対する体積率で10％以下とすることにより、穴拡げ性の低下を抑制することができる。従って、残留オーステナイト相の組織全体に対する体積率は３〜10％の範囲とする。好ましくは５〜８％の範囲である。
Residual austenite phase: 3 to 10% by volume with respect to the entire structure,
The residual austenite phase not only contributes to strengthening of the steel, but also works effectively to improve the balance between steel TS and EL. Such an effect can be obtained when the retained austenite phase is 3% or more by volume ratio with respect to the entire structure. On the other hand, the retained austenite phase transforms into martensite during processing and decreases hole expansibility, but the decrease in hole expansibility can be suppressed by setting the volume ratio to 10% or less with respect to the entire structure. Therefore, the volume ratio with respect to the whole structure | tissue of a retained austenite phase shall be 3 to 10% of range. Preferably it is 5 to 8% of range.
なお、一般的に残留オーステナイトが存在すると、残留オーステナイトのTRIP効果により延性が向上する。一方、歪の付加により残留オーステナイトが変態して生成するマルテンサイトは、非常に硬質なものとなり、その結果、主相であるフェライトとの硬度差が大きくなり、伸びフランジ性（穴拡げ性）を低下させることが知られている。
本発明では、鋼板の成分組成および組織を厳密に制御することにより、高い延性と高い伸びフランジ性を両立しており、残留オーステナイトが存在するにもかかわらず、高い伸びフランジ性を確保している。ここに、残留オーステナイトが存在しても高い伸びフランジ性を得ることできる理由については必ずしも明らかではないが、発明者らは、本発明の鋼組織が残留オーステナイトと焼戻しマルテンサイトの複合組織となっているためと考えている。
In general, when retained austenite is present, ductility is improved by the TRIP effect of retained austenite. On the other hand, the martensite produced by the transformation of retained austenite due to the addition of strain becomes very hard. As a result, the hardness difference from the main phase ferrite increases, and stretch flangeability (hole expandability) increases. It is known to reduce.
In the present invention, by strictly controlling the composition and structure of the steel sheet, both high ductility and high stretch flangeability are achieved, and high stretch flangeability is ensured despite the presence of retained austenite. . Here, the reason why high stretch flangeability can be obtained even if residual austenite is present is not necessarily clear, but the inventors have made the steel structure of the present invention a composite structure of residual austenite and tempered martensite. I think it is because.
また、フェライト相、マルテンサイト相、焼戻しマルテンサイト相および残留オーステナイト相以外の相としては、パーライト相およびベイナイト相を含むことができる。ただし、延性および穴拡げ性確保の観点から、パーライト相やベイナイト相は３％以下とすることが望ましい。 Moreover, as phases other than a ferrite phase, a martensite phase, a tempered martensite phase, and a retained austenite phase, a pearlite phase and a bainite phase can be included. However, from the viewpoint of ensuring ductility and hole expandability, the pearlite phase and the bainite phase are preferably 3% or less.
なお、本発明におけるフェライト相、マルテンサイト相および焼戻しマルテンサイト相の面積率は、次のようにして求めることができる。
すなわち、鋼板の圧延方向に平行な板厚断面を研磨後、３％ナイタールで腐食し、SEM(走査電子顕微鏡)を用いて1000〜3000倍の倍率で10視野観察し、これら画像をMedia Cybernetics社製の画像解析ソフト“Image Pro Plus ver.4.0”で画像解析処理して、フェライト相、マルテンサイト相および焼戻しマルテンサイト相の面積を求め、その全観察面積に占める割合を、各相の面積率とする。
The area ratio of the ferrite phase, martensite phase and tempered martensite phase in the present invention can be determined as follows.
That is, after the plate thickness cross section parallel to the rolling direction of the steel plate was polished, it was corroded with 3% nital, and 10 fields of view were observed at a magnification of 1000 to 3000 using a scanning electron microscope (SEM). Image analysis software “Image Pro Plus ver.4.0” to analyze the area of the ferrite phase, martensite phase and tempered martensite phase. And
また、残留オーステナイト相の体積率は、鋼板を板厚方向の1/4面まで研磨し、この板厚1/4面の回折Ｘ線強度により求めた。入射Ｘ線にはMoKα線を使用し、残留オーステナイト相の｛111｝、｛200｝、｛220｝、｛311｝面とフェライト相の｛110｝、｛200｝、｛211｝面のピークの積分強度の全ての組み合わせについて強度比を求め、これらの平均値を残留オーステナイト相の体積率とした。 Further, the volume ratio of the retained austenite phase was determined by diffracting X-ray intensities on the 1/4 plane of the plate thickness after polishing the steel plate to 1/4 plane in the plate thickness direction. For incident X-rays, MoKα rays are used, and the peaks of {111}, {200}, {220}, {311} in the retained austenite phase and {110}, {200}, {211} in the ferrite phase Intensity ratios were obtained for all combinations of integrated intensities, and the average value of these ratios was taken as the volume ratio of the retained austenite phase.
３）機械特性
本発明で目標とする鋼板の機械特性は次の通りである。
引張強さTS：590MPa以上
強度−伸びバランスTS×EL：22000MPa・％以上
穴拡げ率λ：70％以上
板厚方向の硬さばらつきΔHv：20以下
3) Mechanical properties The mechanical properties of the steel sheet targeted in the present invention are as follows.
Tensile strength TS: 590 MPa or more Strength-elongation balance TS × EL: 22000 MPa ·% or more Hole expansion ratio λ: 70% or more Hardness variation in the plate thickness direction ΔHv: 20 or less
本発明では、上記の機械特性の中でも、板厚方向の硬さばらつきΔHvを20以下とすることが極めて重要である。というのは、板厚方向の硬さばらつきΔHvが20を超える場合、引張特性、穴拡げ性などの機械特性が安定して得られないからである。好ましくは板厚方向の硬さばらつきΔHvが15以下である。
ここに、この板厚方向の硬さばらつきΔHvを20以下とするには、冷間圧延に先立って行う焼鈍処理が特に重要である。この焼鈍処理については、後述する４）製造条件において、詳しく説明する。
なお、板厚方向の硬さばらつきΔHvは、鋼板の板厚方向に0.1mmピッチで全板厚にわたり断面硬さを測定し、得られた断面硬さの最大値と最小値の差として求めることができる。
In the present invention, among the above-mentioned mechanical properties, it is extremely important that the hardness variation ΔHv in the thickness direction is 20 or less. This is because when the hardness variation ΔHv in the thickness direction exceeds 20, mechanical properties such as tensile properties and hole expansibility cannot be stably obtained. The hardness variation ΔHv in the thickness direction is preferably 15 or less.
Here, in order to set the hardness variation ΔHv in the sheet thickness direction to 20 or less, an annealing process prior to cold rolling is particularly important. This annealing process will be described in detail in 4) Manufacturing conditions described later.
Note that the hardness variation ΔHv in the plate thickness direction is determined as the difference between the maximum and minimum cross-sectional hardness values obtained by measuring the cross-sectional hardness over the entire plate thickness at a pitch of 0.1 mm in the plate thickness direction of the steel plate. Can do.
４）製造条件
次に、本発明の製造方法について説明する。
まず、上記の成分組成に調整した鋼を転炉などで溶製し、連続鋳造法等でスラブとする。使用するスラブは、成分のマクロ偏析を防止するために連続鋳造法で製造するのが好ましいが、造塊法、薄スラブ鋳造法で製造してもよい。また、スラブを製造したのち、いったん室温まで冷却し、その後再度加熱する従来法に加え、室温まで冷却しないで、温片のままで加熱炉に挿入する、あるいはわずかの保熱をおこなった後に直ちに圧延する直送圧延・直接圧延などの省エネルギープロセスも問題なく適用できる。
4) Manufacturing conditions Next, the manufacturing method of this invention is demonstrated.
First, steel adjusted to the above component composition is melted in a converter or the like, and is made into a slab by a continuous casting method or the like. The slab to be used is preferably produced by a continuous casting method in order to prevent macro segregation of components, but may be produced by an ingot-making method or a thin slab casting method. Also, after manufacturing the slab, in addition to the conventional method of cooling to room temperature and then heating again, do not cool to room temperature, insert it into a heating furnace as it is, or immediately after a little heat retention Energy-saving processes such as direct rolling and direct rolling for rolling can be applied without problems.
スラブ加熱温度：1100℃以上
ついで、製造したスラブを加熱する。スラブ加熱温度は、低い方がエネルギー的には好ましいが、1100℃未満では、炭化物が十分に固溶できなかったり、圧延荷重の増大による熱間圧延時のトラブル発生の危険が増大する場合がある。このため、スラブ加熱温度は1100℃以上とすることが好ましい。なお、酸化重量の増加に伴うスケールロスの増大などの観点から、スラブ加熱温度は1300℃以下とすることが好ましい。
また、スラブ加熱温度を低くしても、シートバーを加熱することにより、熱間圧延時のトラブルを防止することができる、いわゆるシートバーヒーターを活用してもよい。
Slab heating temperature: 1100 ° C or higher Next, the manufactured slab is heated. A lower slab heating temperature is preferable in terms of energy, but if it is less than 1100 ° C, carbides may not be sufficiently dissolved, or there may be an increased risk of trouble during hot rolling due to an increase in rolling load. . For this reason, it is preferable that slab heating temperature shall be 1100 degreeC or more. Note that the slab heating temperature is preferably 1300 ° C. or less from the viewpoint of an increase in scale loss accompanying an increase in oxidized weight.
Moreover, even if the slab heating temperature is lowered, a so-called sheet bar heater that can prevent troubles during hot rolling by heating the sheet bar may be used.
上記のように加熱されたスラブに、粗圧延および仕上圧延からなる熱間圧延を施す。粗圧延の条件は特に規定する必要はなく、常法に従って行えばよい。また、仕上圧延は、次の条件を満足させることが好ましい。
仕上圧延終了温度：850℃以上
仕上圧延終了温度が850℃未満では、圧延中にα相とγ相が生成して、鋼板にバンド状組織が生成し易くなる。かかるバンド状組織は、冷間圧延後や焼鈍後にも残留し、材料特性に異方性を生じさせたり、加工性を低下させる原因となる場合がある。このため、仕上圧延温度は850℃以上とすることが好ましい。
The slab heated as described above is subjected to hot rolling consisting of rough rolling and finish rolling. The conditions for rough rolling need not be specified, and may be performed according to a conventional method. Moreover, it is preferable that finish rolling satisfies the following conditions.
Finish rolling end temperature: 850 ° C. or higher If the finish rolling end temperature is less than 850 ° C., an α phase and a γ phase are generated during rolling, and a band-like structure is easily generated on the steel sheet. Such a band-like structure remains even after cold rolling or annealing, and may cause anisotropy in material characteristics or cause a decrease in workability. For this reason, it is preferable that finishing rolling temperature shall be 850 degreeC or more.
巻取り温度：450〜700℃
巻取り温度が450℃未満であると、巻取り温度の制御が難しく温度ムラが生じやすくなり、その結果、冷間圧延性が低下するなどの問題が生じる場合がある。また、巻取り温度が700℃を超えると、地鉄表層で脱炭が生じるなどの問題が起こる場合がある。このため、巻取り温度は450〜700℃の範囲とすることが好ましい。
Winding temperature: 450-700 ° C
When the coiling temperature is less than 450 ° C., it is difficult to control the coiling temperature, and temperature unevenness is likely to occur. As a result, problems such as a decrease in cold rollability may occur. In addition, when the coiling temperature exceeds 700 ° C., problems such as decarburization may occur in the surface layer of the railway. For this reason, the winding temperature is preferably in the range of 450 to 700 ° C.
また、本発明における熱延工程では、熱間圧延時の圧延荷重を低減するために仕上圧延の一部または全部を潤滑圧延としてもよい。潤滑圧延を行うことは、鋼板形状の均一化、材質の均一化の観点からも有効である。
なお、潤滑圧延の際の摩擦係数は0.10〜0.25の範囲とすることが好ましい。また、相前後するシートバー同士を接合し、連続的に仕上圧延する連続圧延プロセスとすることが好ましい。連続圧延プロセスを適用することは、熱間圧延の操業安定性の観点からも望ましい。
Moreover, in the hot rolling process in this invention, in order to reduce the rolling load at the time of hot rolling, it is good also considering a part or all of finish rolling as lubrication rolling. Performing lubrication rolling is also effective from the viewpoint of uniform steel plate shape and uniform material.
In addition, it is preferable to make the friction coefficient in the case of lubrication rolling into the range of 0.10-0.25. Moreover, it is preferable to set it as the continuous rolling process which joins the sheet | seat bars which precede and follow, and finish-rolls continuously. The application of the continuous rolling process is also desirable from the viewpoint of the operational stability of hot rolling.
その後、通常は冷間圧延に供するのであるが、本発明では、この冷間圧延に先立ち、以下の条件で焼鈍処理を施す。本発明においては、この焼鈍処理が特に重要な工程である。
焼鈍処理条件：500℃以上Ａc1点以下の温度域で１〜10時間保持
冷間圧延に先立ち、適正な条件で焼鈍処理を行うことにより、鋼板の板厚方向の組織が改質されて、板厚方向の機械特性、特に硬さのばらつきΔHvを大幅に抑制することができる。
ここに、焼鈍処理における保持温度が500℃に満たない場合、または保持時間が１時間に満たない場合、熱延板組織中の硬質相（パーライト、ベイナイト）の軟化が不十分なため、冷間圧延後の連続焼鈍過程において再結晶が不均一に進行し、焼鈍後の組織も不均一となり、板厚方向の硬さばらつきΔHvを抑制することができない。一方、保持温度がＡc1点を超える場合、焼鈍後の冷却過程で硬質相（パーライト、ベイナイト）が再生してしまい、冷間圧延および連続焼鈍後の組織が不均一となり、板厚方向の硬さばらつきΔHvを抑制することができない。また、保持時間が10時間を超える場合、硬質相（パーライト、ベイナイト）中の炭化物が粗大化し過ぎてしまい、連続焼鈍過程での炭化物の溶解が不均一に進行するため、連続焼鈍後の組織が不均一となり、板厚方向の硬さばらつきΔHvを抑制することができない。従って、焼鈍処理における保持温度は500℃以上Ａc1点以下、保持時間は１〜10時間の範囲とする。好ましくは保持温度：600℃以上700℃以下、保持時間：３〜10時間の範囲である。
Then, although it normally uses for cold rolling, in this invention, prior to this cold rolling, an annealing process is performed on condition of the following. In the present invention, this annealing process is a particularly important process.
Annealing treatment conditions: Hold for 1 to 10 hours in a temperature range of 500 ° C. or more and Ac 1 point or less Prior to cold rolling, the structure in the thickness direction of the steel sheet is modified by performing annealing treatment under appropriate conditions, Mechanical properties in the plate thickness direction, particularly the hardness variation ΔHv can be greatly suppressed.
Here, when the holding temperature in annealing treatment is less than 500 ° C., or when the holding time is less than 1 hour, the soft phase of the hard phase (pearlite, bainite) in the hot rolled sheet structure is insufficiently softened. In the continuous annealing process after rolling, recrystallization proceeds non-uniformly, the structure after annealing becomes non-uniform, and the hardness variation ΔHv in the thickness direction cannot be suppressed. On the other hand, when the holding temperature exceeds Ac 1 point, the hard phase (pearlite, bainite) is regenerated in the cooling process after annealing, the structure after cold rolling and continuous annealing becomes non-uniform, and the hardness in the thickness direction The variation ΔHv cannot be suppressed. In addition, when the holding time exceeds 10 hours, the carbide in the hard phase (pearlite, bainite) becomes too coarse, and the dissolution of the carbide in the continuous annealing process proceeds non-uniformly. It becomes non-uniform, and the hardness variation ΔHv in the thickness direction cannot be suppressed. Accordingly, the holding temperature in the annealing treatment is 500 ° C. or more and Ac 1 point or less, and the holding time is in the range of 1 to 10 hours. Preferably, the holding temperature is 600 ° C. or higher and 700 ° C. or lower, and the holding time is 3 to 10 hours.
本発明では、上記したように、冷間圧延に先立って、適正な焼鈍処理を施すことにより、
熱延板組織中の硬質相（パーライト、ベイナイト）を軟質化させ、冷間圧延後の連続焼鈍過程において再結晶が均一に進行し、組織が均一化するように、鋼組織が改質され、その結果、板厚方向の機械特性のばらつきが解消されるのである。
In the present invention, as described above, by performing an appropriate annealing treatment prior to cold rolling,
The steel structure is modified so that the hard phase (pearlite, bainite) in the hot-rolled sheet structure is softened, and recrystallization proceeds uniformly in the continuous annealing process after cold rolling, so that the structure becomes uniform. As a result, variations in mechanical properties in the thickness direction are eliminated.
なお、この焼鈍処理は、熱延鋼板の表面の酸化スケールを酸洗により除去した後に行うことが好ましい。ここに、酸洗条件は特に制限されるものではなく、常法に従えば良い。 In addition, it is preferable to perform this annealing process after removing the oxidation scale of the surface of a hot-rolled steel plate by pickling. Here, the pickling conditions are not particularly limited, and may be according to ordinary methods.
冷間圧延の圧下率：60％超
上記のような焼鈍処理を経て得られた熱延板を、冷間圧延する。ここに、冷間圧延の圧下率は、再結晶の促進と再結晶粒の微細化のため、60％超とする。これ以外の条件は特に規定する必要はなく、常法に従って行えばよい。
Cold rolling reduction: over 60% The hot-rolled sheet obtained through the annealing treatment as described above is cold-rolled. Here, the rolling reduction of cold rolling is over 60% in order to promote recrystallization and refine the recrystallized grains. Conditions other than this need not be specified, and may be performed according to ordinary methods.
ついで、このようにして得られた冷延鋼板を以下の条件で連続焼鈍する。
少なくとも500℃以上Ａc1点以下の温度域における平均加熱速度：10℃/s以上
本発明の鋼における再結晶温度域である500℃からＡc1点の温度域において平均加熱速度を10℃/s以上とすることで、加熱昇温時の再結晶が抑制され、Ａc1点以上で生成するオーステナイト相（γ相）の微細化、ひいては焼鈍冷却後の残留オーステナイト相の微細化に有効に働く。また、平均加熱速度が10℃/s未満では、加熱昇温時にフェライト相（α相）の再結晶の進行が過度に進み、十分な微細化が達成できなくなる。従って、500℃以上Ａc1点以下の温度域における平均加熱速度は10℃/s以上とする。好ましくは15℃/s以上である。
Next, the cold-rolled steel sheet thus obtained is continuously annealed under the following conditions.
Average heating rate in a temperature range of at least 500 ° C. or more and Ac 1 point or less: 10 ° C./s or more The average heating rate in the temperature range from 500 ° C., which is the recrystallization temperature range of the steel of the present invention, to Ac 1 point By setting it as the above, the recrystallization at the time of heating temperature rise is suppressed, and it works effectively for refinement | miniaturization of the austenite phase (gamma phase) produced | generated in Ac 1 point or more, and also refinement | miniaturization of the retained austenite phase after annealing cooling. On the other hand, if the average heating rate is less than 10 ° C./s, the recrystallization of the ferrite phase (α phase) proceeds excessively at the time of heating and heating, and sufficient refinement cannot be achieved. Accordingly, the average heating rate in the temperature range of 500 ° C. or more and Ac 1 point or less is set to 10 ° C./s or more. Preferably it is 15 ° C./s or more.
750〜900℃の温度域で10秒以上保持
保持温度が750℃未満あるいは保持時間が10秒未満では、焼鈍時のオーステナイト相の生成が不十分となり、焼鈍冷却後に十分な量の低温変態相が確保できなくなる。一方、加熱温度が900℃を超えると、加熱時に生成するオーステナイト相が粗大化し、焼鈍後の残留オーステナイト相も粗大となる。従って、750〜900℃の温度域で10秒以上保持するものとする。
なお、保持時間の上限は特に規定しないが、600秒以上の保持は効果が飽和する上、コストアップにつながるので、保持時間は600秒未満とすることが好ましい。
Hold for 10 seconds or more in the temperature range of 750 to 900 ° C If the holding temperature is less than 750 ° C or the holding time is less than 10 seconds, the austenite phase is not sufficiently generated during annealing, and a sufficient amount of low-temperature transformation phase is present after annealing cooling. It cannot be secured. On the other hand, when the heating temperature exceeds 900 ° C., the austenite phase generated during heating becomes coarse, and the residual austenite phase after annealing becomes coarse. Therefore, it shall hold | maintain for 10 seconds or more in the temperature range of 750-900 degreeC.
The upper limit of the holding time is not particularly defined, but holding for 600 seconds or more saturates the effect and leads to an increase in cost, so the holding time is preferably less than 600 seconds.
冷却過程における750℃から（Ms点−100℃）〜（Ms点−200℃）の温度域までの平均冷却速度：10℃/s以上
冷却過程における750℃から（Ms点−100℃）〜（Ms点−200℃）の温度域までの平均冷却速度が10℃/s未満ではパーライトが生成し、TSとELのバランスおよび穴拡げ性が低下する。このため、冷却過程における750℃から（Ms点−100℃）〜（Ms点−200℃）の温度域までの平均冷却速度は、10℃/s以上とする。好ましくは30℃/s以上である。
なお、この平均冷却速度の上限は特に規定しないが、平均冷却速度が速すぎると鋼板形状が悪化したり、冷却停止温度の制御が困難となるため、200℃/s以下とすることが好ましい。
Average cooling rate from 750 ° C in the cooling process to the temperature range from (Ms point – 100 ° C) to (Ms point – 200 ° C): 10 ° C / s or more From 750 ° C in the cooling process (Ms point – 100 ° C) to ( When the average cooling rate up to the temperature range of (Ms point -200 ° C) is less than 10 ° C / s, pearlite is generated, and the balance between TS and EL and the hole expandability deteriorate. For this reason, the average cooling rate from 750 ° C. to the temperature range of (Ms point−100 ° C.) to (Ms point−200 ° C.) in the cooling process is set to 10 ° C./s or more. Preferably it is 30 ° C./s or more.
The upper limit of the average cooling rate is not particularly defined. However, if the average cooling rate is too high, the shape of the steel sheet is deteriorated and it is difficult to control the cooling stop temperature.
また、上記した冷却過程における冷却停止温度を（Ms点−100℃）〜（Ms点−200℃）の範囲とすることは、本発明において重要な条件の一つである。
すなわち、冷却停止時には、オーステナイト相の一部がマルテンサイトに変態し、残りは未変態のオーステナイト相となっている。そこから再加熱し、めっき処理、必要に応じて合金化処理を施した後、室温まで冷却することにより、マルテンサイト相は焼戻しマルテンサイト相となり、未変態オーステナイト相は残留オーステナイト相またはマルテンサイト相となる。このため、冷却停止温度が低い、すなわちMs点（Ms点：オーステナイトのマルテンサイト変態が開始する温度）からの過冷度が大きいほど、冷却中に生成するマルテンサイト量が増加し、未変態オーステナイト量が減少する。換言すれば、冷却停止温度の制御により、最終的なマルテンサイト相および残留オーステナイト相と焼戻しマルテンサイト相の面積率が決定されることになる。従って、本発明では、Ms点と冷却停止温度の差である過冷度を適切に制御することが必要となる。
In addition, it is one of the important conditions in the present invention that the cooling stop temperature in the cooling process described above is in the range of (Ms point−100 ° C.) to (Ms point−200 ° C.).
That is, when cooling is stopped, a part of the austenite phase is transformed into martensite, and the rest is an untransformed austenite phase. After reheating, plating treatment, alloying treatment if necessary, and cooling to room temperature, the martensite phase becomes tempered martensite phase, and the untransformed austenite phase is the retained austenite phase or martensite phase. It becomes. Therefore, the lower the cooling stop temperature, that is, the greater the degree of supercooling from the Ms point (Ms point: the temperature at which austenite martensite transformation starts), the greater the amount of martensite generated during cooling, and untransformed austenite. The amount decreases. In other words, the area ratio of the final martensite phase, residual austenite phase, and tempered martensite phase is determined by controlling the cooling stop temperature. Therefore, in the present invention, it is necessary to appropriately control the degree of supercooling, which is the difference between the Ms point and the cooling stop temperature.
ここに、冷却停止温度が（Ms点−100℃）より高い温度では、冷却停止時にマルテンサイト変態が不十分で未変態オーステナイト量が多くなる。これによって、最終的なマルテンサイト相または残留オーステナイト相が過剰に生成し、穴拡げ性が低下する。一方、冷却停止温度が（Ms点−200℃）より低くなると、冷却中にオーステナイト相がほとんどマルテンサイトに変態して未変態オーステナイト量が減少し、３％以上の残留オーステナイト相が得られない。従って、冷却停止温度は（Ms点−100℃）〜（Ms点−200℃）の範囲とする。好ましくは（Ms点−130℃）〜（Ms点−200℃）の範囲である。 Here, when the cooling stop temperature is higher than (Ms point −100 ° C.), the martensitic transformation is insufficient and the amount of untransformed austenite increases when the cooling is stopped. As a result, the final martensite phase or residual austenite phase is excessively generated, and the hole expandability is lowered. On the other hand, when the cooling stop temperature is lower than (Ms point −200 ° C.), the austenite phase is almost transformed into martensite during cooling, and the amount of untransformed austenite is reduced, and a residual austenite phase of 3% or more cannot be obtained. Accordingly, the cooling stop temperature is in the range of (Ms point−100 ° C.) to (Ms point−200 ° C.). Preferably, it is in the range of (Ms point−130 ° C.) to (Ms point−200 ° C.).
なお、Ms点は、オーステナイトからのマルテンサイト変態開始温度であり、焼鈍温度からの冷却過程における鋼板の線膨張係数の変化から求めることができる。 The Ms point is the martensitic transformation start temperature from austenite and can be obtained from the change in the linear expansion coefficient of the steel sheet in the cooling process from the annealing temperature.
350〜600℃の温度域まで再加熱して10〜600秒保持
また、本発明では、上記した連続焼鈍に引き続き、350〜600℃の温度域まで再加熱して10〜600秒保持する。これによって、前記冷却時に生成したマルテンサイト相が焼戻されて焼戻しマルテンサイト相となり、穴拡げ性が向上する。さらには、冷却時にマルテンサイトに変態しなかった未変態オーステナイト相が安定化され、最終的に３％以上の残留オーステナイト相が得られ、延性が向上する。
再加熱およびその後の保持による未変態オーステナイト相の安定化のメカニズムは、必ずしも明らかではないが、発明者らは、未変態オーステナイトへのＣの濃化が進むために、オーステナイト相が安定化されるものと考えている。
Reheating to a temperature range of 350 to 600 ° C. and holding for 10 to 600 seconds In the present invention, following the above-described continuous annealing, reheating to a temperature range of 350 to 600 ° C. and holding for 10 to 600 seconds. As a result, the martensite phase generated during the cooling is tempered to become a tempered martensite phase, and the hole expandability is improved. Furthermore, the untransformed austenite phase that did not transform into martensite during cooling is stabilized, and a residual austenite phase of 3% or more is finally obtained, thereby improving ductility.
Although the mechanism of stabilization of the untransformed austenite phase by reheating and subsequent holding is not necessarily clear, the inventors have stabilized the austenite phase as the concentration of C into untransformed austenite proceeds. I believe that.
ここに、再加熱温度が350℃未満ではマルテンサイト相の焼戻しおよびオーステナイト相の安定化が不十分となり、穴拡げ性および延性が低下する。一方、再加熱温度が600℃を超えると、冷却停止時の未変態オーステナイト相がパーライト相に変態し、最終的に３％以上残留オーステナイト相が得られなくなる。従って、再加熱温度は350〜600℃の範囲とする。好ましくは400〜550℃の範囲である。 Here, when the reheating temperature is less than 350 ° C., the tempering of the martensite phase and the stabilization of the austenite phase are insufficient, and the hole expansibility and ductility are lowered. On the other hand, when the reheating temperature exceeds 600 ° C., the untransformed austenite phase at the time of cooling stop is transformed into a pearlite phase, and finally a retained austenite phase of 3% or more cannot be obtained. Therefore, the reheating temperature is in the range of 350 to 600 ° C. Preferably it is the range of 400-550 degreeC.
また、再加熱時の保持時間が10秒未満では、オーステナイト相の安定化が不十分となる。一方、600秒を超えると冷却停止時の未変態オーステナイト相がベイナイトに変態し、最終的に３％以上の残留オーステナイト相が得られなくなる。従って、再加熱時の保持時間は10〜600秒の範囲とする。好ましくは10〜100秒の範囲である。 Further, if the holding time during reheating is less than 10 seconds, the austenite phase is not sufficiently stabilized. On the other hand, if it exceeds 600 seconds, the untransformed austenite phase at the time of cooling stop transforms into bainite, and finally a 3% or more retained austenite phase cannot be obtained. Accordingly, the holding time at the time of reheating is in the range of 10 to 600 seconds. The range is preferably 10 to 100 seconds.
さらに、上記のようにして得られた鋼板に、溶融亜鉛めっき処理を施し、鋼板表面の片側あるいは両側に溶融亜鉛めっき層を設ける。ここでは、0.12〜0.22％の溶解Al量のめっき浴に（浴温440〜500℃）鋼板を侵入させ、ガスワイピングなどで付着量を調整することで、鋼板表面に溶融亜鉛めっき層を設けることができる。
また、上記のめっき浴における溶解Al量を0.08〜0.18％として、その他は同じ条件でめっき処理を行い、めっき処理後、さらに450〜600℃まで加熱し、１〜30秒保持する合金化処理を施すことにより、鋼板表面の溶融亜鉛めっきを合金化溶融亜鉛めっきとすることができる。
Furthermore, the steel plate obtained as described above is subjected to a hot dip galvanizing treatment, and a hot dip galvanized layer is provided on one side or both sides of the steel plate surface. Here, a hot dip galvanized layer is provided on the surface of the steel sheet by allowing the steel sheet to penetrate into the plating bath with a dissolved Al content of 0.12 to 0.22% (bath temperature 440 to 500 ° C.) and adjusting the amount of adhesion by gas wiping or the like. Can do.
In addition, the amount of dissolved Al in the above plating bath is 0.08 to 0.18%, and the others are plated under the same conditions. After the plating, the alloy is further heated to 450 to 600 ° C. and held for 1 to 30 seconds. By applying, hot dip galvanization of the steel sheet surface can be made into alloyed hot dip galvanization.
なお、溶融亜鉛めっき処理後の鋼板（その後、さらに合金化処理を行った鋼板も含む）には、形状矯正、表面粗度等の調整のため調質圧延を加えてもよい。また、樹脂あるいは油脂コーティング、各種塗装等の処理を施しても何ら不都合はない。 In addition, you may add temper rolling to the steel plate after a hot-dip galvanization process (including the steel plate which performed alloying treatment after that) for adjustment of shape correction, surface roughness, etc. In addition, there is no inconvenience even if treatments such as resin or oil coating and various paintings are applied.
表１に示す成分組成を有し、残部がFeおよび不可避的不純物からなる鋼を転炉にて溶製し、連続鋳造法にて鋳片とした。得られた鋳片は、1200℃に加熱後、仕上圧延終了温度：900℃の条件で熱間圧延し、圧延終了後、冷却速度10℃/sで冷却し、600℃で巻き取った。ついで、得られた熱延鋼板を、酸洗後、還元ガス雰囲気中にて、表２に示す条件で焼鈍処理を行ったのち、冷間圧延を施して、板厚：1.2mmの冷延鋼板を製造した。
なお、冷間圧延は、圧下率を40〜80％とし、熱間圧延時の仕上げ板厚は、冷間圧延時の圧下率に応じて調整した。また、表２中のＡc1点は次式により、Ms点は連続焼鈍の冷却過程における鋼板の線膨張係数の変化により、それぞれ求めた。
Ａc1＝723-10.7[%Mn]-16.9[%Ni]+29.1[%Si]+16.9[%Cr]
ただし、[%Ｍ]は、Ｍ元素の含有量（質量％）を表す。
Steel having the composition shown in Table 1 and the balance being Fe and unavoidable impurities was melted in a converter and made into a slab by a continuous casting method. The obtained slab was heated to 1200 ° C., hot-rolled at a finish rolling finish temperature of 900 ° C., cooled at a cooling rate of 10 ° C./s after rolling, and wound up at 600 ° C. Next, the obtained hot-rolled steel sheet was pickled, annealed in a reducing gas atmosphere under the conditions shown in Table 2, and then cold-rolled to obtain a cold-rolled steel sheet having a thickness of 1.2 mm. Manufactured.
In cold rolling, the rolling reduction was 40 to 80%, and the finished sheet thickness during hot rolling was adjusted according to the rolling reduction during cold rolling. Further, the Ac 1 point in Table 2 was obtained from the following equation, and the Ms point was obtained from the change in the coefficient of linear expansion of the steel sheet during the cooling process of continuous annealing.
Ac 1 = 723-10.7 [% Mn] -16.9 [% Ni] +29.1 [% Si] +16.9 [% Cr]
However, [% M] represents the content (mass%) of the M element.
ついで、上記により得られた冷延鋼板に、連続溶融亜鉛めっきラインにて、表２に示す条件で連続焼鈍を施した。その後、表２に示す条件で再加熱後、保持し、ついで460℃で溶融亜鉛めっきを施したのち、平均冷却速度10℃/sの条件で室温まで冷却した。なお、一部の鋼板については、めっき処理後、さらに520℃に加熱して、合金化処理を行った。ここに、めっき付着量は、片面あたり35〜45g/m2であった。 Subsequently, the cold-rolled steel sheet obtained as described above was subjected to continuous annealing under the conditions shown in Table 2 in a continuous hot-dip galvanizing line. Thereafter, after being reheated under the conditions shown in Table 2, it was held, then hot dip galvanized at 460 ° C, and then cooled to room temperature at an average cooling rate of 10 ° C / s. Note that some of the steel sheets were further alloyed by heating to 520 ° C. after plating. Here, the plating adhesion amount was 35 to 45 g / m 2 per side.
かくして得られた溶融亜鉛めっき鋼板について、鋼組織、引張特性、穴拡げ性および板厚方向の硬さばらつきを調査した。得られた結果を表３に示す。 The hot-dip galvanized steel sheet thus obtained was investigated for steel structure, tensile properties, hole expandability, and hardness variation in the thickness direction. The obtained results are shown in Table 3.
なお、溶融亜鉛めっき鋼板の鋼組織、引張特性、穴拡げ性および板厚方向の硬さばらつきは、次のように測定した。
（１）鋼組織
鋼板の鋼組織は、３％ナイタール溶液（３％硝酸＋エタノール）で組織を現出し、走査型電子顕微鏡で深さ方向板厚1/4位置を10視野観察して、撮影した組織写真を用いて、画像解析処理を行ない、フェライト相、マルテンサイト相、および焼戻しマルテンサイト相の面積率を各相の分率として定量化した。
組織写真は、各組織の細かさに応じて1000〜3000倍の適切な倍率で撮影した。なお、画像解析処理については、Media Cybernetics社製の画像解析ソフト“Image Pro Plus ver.4.0”を用いて各相の面積を求め、全観察面積に占める割合（面積率）を求めた。
In addition, the steel structure of the hot dip galvanized steel sheet, tensile properties, hole expansibility, and hardness variation in the thickness direction were measured as follows.
(1) Steel structure The steel structure of the steel sheet was revealed with a 3% nital solution (3% nitric acid + ethanol), and 10 depths in the depth direction were observed with a scanning electron microscope. Image analysis processing was performed using the obtained structure photograph, and the area ratio of the ferrite phase, the martensite phase, and the tempered martensite phase was quantified as a fraction of each phase.
Tissue photographs were taken at an appropriate magnification of 1000 to 3000 times depending on the fineness of each tissue. For the image analysis processing, the area of each phase was determined using the image analysis software “Image Pro Plus ver. 4.0” manufactured by Media Cybernetics, and the ratio (area ratio) in the total observation area was determined.
また、残留オーステナイト相の体積率は、鋼板を板厚方向の1/4面まで研磨し、この板厚1/4面の回折Ｘ線強度により求めた。入射Ｘ線にはMoKα線を使用し、残留オーステナイト相の｛111｝、｛200｝、｛220｝、｛311｝面とフェライト相の｛110｝、｛200｝、｛211｝面のピークの積分強度の全ての組み合わせについて強度比を求め、これらの平均値を残留オーステナイト相の体積率とした。 Further, the volume ratio of the retained austenite phase was determined by diffracting X-ray intensities on the 1/4 plane of the plate thickness after polishing the steel plate to 1/4 plane in the plate thickness direction. For incident X-rays, MoKα rays are used, and the peaks of {111}, {200}, {220}, {311} in the retained austenite phase and {110}, {200}, {211} in the ferrite phase Intensity ratios were obtained for all combinations of integrated intensities, and the average value of these ratios was taken as the volume ratio of the retained austenite phase.
（２）引張特性
引張特性は、引張方向が鋼板の圧延方向と直角方向となるようサンプル採取したJIS5号試験片を用いて、JIS Z 2241（2011年）に準拠した引張試験を行い、TS（引張強さ）およびEL（伸び）を測定し、引張強さと伸びの積（TS×EL）で表される強度と伸びバランスの値を求めた。また、試験のＮ数は３とし、測定した値の平均値をTS、ELおよびTS×ELとして求めた。また、ELについて、測定した値の最大値と最小値の差（ΔEL）を求めた。
(2) Tensile properties Tensile properties are obtained by conducting a tensile test in accordance with JIS Z 2241 (2011) using JIS No. 5 test pieces sampled so that the tensile direction is perpendicular to the rolling direction of the steel sheet. Tensile strength) and EL (elongation) were measured, and the value of strength and elongation balance represented by the product of tensile strength and elongation (TS × EL) was determined. Moreover, N number of the test was set to 3, and the average value of the measured value was calculated | required as TS, EL, and TSxEL. Further, for EL, the difference (ΔEL) between the maximum value and the minimum value of the measured values was determined.
（３）穴拡げ性
穴拡げ性は、日本鉄鋼連盟規格JFST1001（1996年）に準じた穴拡げ試験を行い、穴拡げ率を測定した。試験のＮ数は５とし、測定した値の平均値を穴拡げ率λとして求めた。また、測定した値の最大値と最小値の差（Δλ）を求めた。
(3) Hole expandability The hole expandability was measured by performing a hole expansion test in accordance with the Japan Iron and Steel Federation Standard JFST1001 (1996). The N number of the test was 5, and the average value of the measured values was determined as the hole expansion rate λ. Further, the difference (Δλ) between the maximum value and the minimum value of the measured values was obtained.
（４）板厚方向の硬さばらつき（ΔHv）
板厚方向の硬さばらつきは、鋼板の板厚方向に0.1mmピッチで全板厚にわたり断面硬さを測定し、得られた断面硬さの最大値と最小値の差をΔHvとして評価した。ここに、断面硬さは、サンプルの切断面を研磨後、ビッカース硬さ試験機を用いて、JIS Z 2244（2009年）に準拠した硬さ試験を行うことにより求めた。なお、試験力は0.98Nとした。
(4) Hardness variation in thickness direction (ΔHv)
The hardness variation in the plate thickness direction was evaluated by measuring the cross-sectional hardness over the entire plate thickness at a pitch of 0.1 mm in the plate thickness direction of the steel plate, and evaluating the difference between the maximum value and the minimum value of the obtained cross-sectional hardness as ΔHv. Here, the cross-sectional hardness was determined by polishing a cut surface of the sample and then performing a hardness test in accordance with JIS Z 2244 (2009) using a Vickers hardness tester. The test force was 0.98N.
なお、上記の各特性試験においては、それぞれTS≧590MPa、TS×EL≧22000MPa・％、ΔEL≦２％、λ≧70％、Δλ≦20％およびΔHv≦20を満足する場合に良好と判定とした。 In each of the above characteristic tests, it was judged as good when TS ≧ 590 MPa, TS × EL ≧ 22000 MPa ·%, ΔEL ≦ 2%, λ ≧ 70%, Δλ ≦ 20%, and ΔHv ≦ 20, respectively. did.
表３より、本発明の鋼板ではいずれも、引張強さが590MPa以上で、強度と伸びのバランス（TS×EL）が22000MPa・％以上、穴拡げ率（λ）が70％以上であり、しかもΔELが２％以下、Δλが20％以下で、かつ板厚方向の硬さばらつき（ΔHv）が20以下であることから、板厚方向のばらつきなしに優れた強度、延性および伸びフランジ性が安定して得られていることがわかる。
一方、比較例となる鋼板はいずれも、引張強さ、強度と伸びのバランス（TS×EL）、穴拡げ率（λ）、板厚方向の硬さばらつき（ΔHv）のうちの少なくとも１つが目標とする特性を満足しておらず、また比較例No.6c、No.6d、No.9およびNo.20ではΔELが２％を超え、さらにΔλも20％を超える結果となった。
From Table 3, all the steel sheets of the present invention have a tensile strength of 590 MPa or more, a balance between strength and elongation (TS × EL) of 22000 MPa ·% or more, and a hole expansion ratio (λ) of 70% or more. Since ΔEL is 2% or less, Δλ is 20% or less, and hardness variation (ΔHv) in the thickness direction is 20 or less, excellent strength, ductility and stretch flangeability are stable without variation in the thickness direction. You can see that it is obtained.
On the other hand, each of the comparative steel plates has a target of at least one of tensile strength, balance between strength and elongation (TS × EL), hole expansion ratio (λ), and hardness variation in the thickness direction (ΔHv). In Comparative Examples No. 6c, No. 6d, No. 9 and No. 20, ΔEL exceeded 2%, and Δλ also exceeded 20%.
Claims (8)
Ｃ：0.05〜0.3％、
Si：0.01〜2.5％、
Mn：0.5〜3.5％、
Ｐ：0.003〜0.100％、
Ｓ：0.02％以下および
Al：0.010〜1.5％
を含有し、かつSiとAlの合計量が0.5〜2.5％であって、残部はFeおよび不可避的不純物からなり、
組織全体に対する面積率でフェライト相を20％以上、マルテンサイト相を10％以下（但し、０％を含む）、焼戻しマルテンサイト相を10〜60％含み、かつ組織全体に対する体積率で残留オーステナイト相を３〜10％含む組織を有し、
板厚方向の硬さばらつきΔHvが20以下であることを特徴とする加工性に優れた高強度溶融亜鉛めっき鋼板。 % By mass
C: 0.05-0.3%
Si: 0.01-2.5%
Mn: 0.5-3.5%
P: 0.003-0.100%
S: 0.02% or less and
Al: 0.010 to 1.5%
And the total amount of Si and Al is 0.5 to 2.5%, the balance consists of Fe and inevitable impurities,
20% or more of ferrite phase, 10% or less (including 0%) of martensite phase, 10-60% of tempered martensite phase, and retained austenite phase by volume ratio of the whole structure. Having a tissue containing 3 to 10%,
A high-strength hot-dip galvanized steel sheet excellent in workability, characterized by having a hardness variation ΔHv in the thickness direction of 20 or less.
請求項１〜５のいずれかに記載の成分組成を有するスラブを、熱間圧延し、500℃以上Ａc1点以下の温度域で１〜10時間保持する焼鈍処理を施したのち、圧下率が60％超となる冷間圧延を施し、
ついで、少なくとも500℃以上Ａc1点以下の温度域における平均加熱速度を10℃/s以上として750〜900℃の温度域まで加熱し、該温度域で10秒以上保持したのち、平均冷却速度を10℃/s以上として750℃から（Ms点−100℃）〜（Ms点−200℃）の温度域まで冷却する連続焼鈍を施し、
さらに、350〜600℃の温度域まで再加熱して10〜600秒保持したのち、鋼板表面に溶融亜鉛めっきを施すことを特徴とする加工性に優れた高強度溶融亜鉛めっき鋼板の製造方法。 A method for producing a high-strength hot-dip galvanized steel sheet excellent in workability according to any one of claims 1 to 5,
The slab having the component composition according to any one of claims 1 to 5 is hot-rolled and subjected to an annealing treatment for 1 to 10 hours in a temperature range of 500 ° C. or higher and Ac 1 point or lower, and then the reduction rate is Cold rolling over 60%,
Next, the average heating rate in a temperature range of at least 500 ° C. to Ac 1 point is set to 10 ° C./s or more and heated to a temperature range of 750 to 900 ° C., held in the temperature range for 10 seconds or more, and then the average cooling rate is set. Apply continuous annealing to cool from 750 ° C to a temperature range of (Ms point-100 ° C) to (Ms point-200 ° C) as 10 ° C / s or more,
Furthermore, after reheating to the temperature range of 350-600 degreeC and hold | maintaining for 10-600 second, hot-dip galvanized steel plate excellent in workability characterized by performing hot dip galvanizing on the steel plate surface.
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