JP4946092B2 - High-strength steel and manufacturing method thereof - Google Patents
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- 229910000831 Steel Inorganic materials 0.000 title claims description 69
- 239000010959 steel Substances 0.000 title claims description 69
- 238000004519 manufacturing process Methods 0.000 title claims description 13
- 238000001816 cooling Methods 0.000 claims description 73
- 229910052791 calcium Inorganic materials 0.000 claims description 16
- 229910052717 sulfur Inorganic materials 0.000 claims description 16
- 238000005098 hot rolling Methods 0.000 claims description 15
- 230000001186 cumulative effect Effects 0.000 claims description 12
- 229910052760 oxygen Inorganic materials 0.000 claims description 12
- 239000000203 mixture Substances 0.000 claims description 11
- 229910052757 nitrogen Inorganic materials 0.000 claims description 7
- 239000012535 impurity Substances 0.000 claims description 6
- 229910052748 manganese Inorganic materials 0.000 claims description 5
- 229910052750 molybdenum Inorganic materials 0.000 claims description 5
- 229910052759 nickel Inorganic materials 0.000 claims description 5
- 229910052720 vanadium Inorganic materials 0.000 claims description 5
- 229910052802 copper Inorganic materials 0.000 claims description 4
- 229910052698 phosphorus Inorganic materials 0.000 claims description 4
- 238000005496 tempering Methods 0.000 claims description 4
- 239000000463 material Substances 0.000 description 35
- 229910000859 α-Fe Inorganic materials 0.000 description 29
- 238000005096 rolling process Methods 0.000 description 22
- 230000000694 effects Effects 0.000 description 15
- 229910001566 austenite Inorganic materials 0.000 description 13
- 230000009466 transformation Effects 0.000 description 12
- 238000003466 welding Methods 0.000 description 12
- 238000000034 method Methods 0.000 description 10
- 230000015572 biosynthetic process Effects 0.000 description 9
- 239000002244 precipitate Substances 0.000 description 9
- 239000010953 base metal Substances 0.000 description 8
- 238000010438 heat treatment Methods 0.000 description 8
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 7
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 6
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 6
- 239000006104 solid solution Substances 0.000 description 6
- 229910001563 bainite Inorganic materials 0.000 description 5
- 238000005516 engineering process Methods 0.000 description 5
- 229910000734 martensite Inorganic materials 0.000 description 5
- 229910052758 niobium Inorganic materials 0.000 description 5
- 229910052799 carbon Inorganic materials 0.000 description 4
- 239000002131 composite material Substances 0.000 description 4
- 238000009863 impact test Methods 0.000 description 4
- 238000005728 strengthening Methods 0.000 description 4
- 230000002411 adverse Effects 0.000 description 3
- 229910045601 alloy Inorganic materials 0.000 description 3
- 239000000956 alloy Substances 0.000 description 3
- 230000007423 decrease Effects 0.000 description 3
- 230000007547 defect Effects 0.000 description 3
- 239000006185 dispersion Substances 0.000 description 3
- 229910052742 iron Inorganic materials 0.000 description 3
- 230000008018 melting Effects 0.000 description 3
- 238000002844 melting Methods 0.000 description 3
- 229910052751 metal Inorganic materials 0.000 description 3
- 239000002184 metal Substances 0.000 description 3
- 150000004767 nitrides Chemical class 0.000 description 3
- 230000006911 nucleation Effects 0.000 description 3
- 238000010899 nucleation Methods 0.000 description 3
- 229910001568 polygonal ferrite Inorganic materials 0.000 description 3
- 238000001556 precipitation Methods 0.000 description 3
- 229910052761 rare earth metal Inorganic materials 0.000 description 3
- 238000009864 tensile test Methods 0.000 description 3
- 230000001276 controlling effect Effects 0.000 description 2
- 238000002425 crystallisation Methods 0.000 description 2
- 230000008025 crystallization Effects 0.000 description 2
- 238000011156 evaluation Methods 0.000 description 2
- 229910001562 pearlite Inorganic materials 0.000 description 2
- 239000002994 raw material Substances 0.000 description 2
- 238000003303 reheating Methods 0.000 description 2
- 150000003568 thioethers Chemical class 0.000 description 2
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 1
- 229910000746 Structural steel Inorganic materials 0.000 description 1
- 230000002159 abnormal effect Effects 0.000 description 1
- 239000000654 additive Substances 0.000 description 1
- 230000000996 additive effect Effects 0.000 description 1
- QVGXLLKOCUKJST-UHFFFAOYSA-N atomic oxygen Chemical compound [O] QVGXLLKOCUKJST-UHFFFAOYSA-N 0.000 description 1
- 229910052796 boron Inorganic materials 0.000 description 1
- 238000005266 casting Methods 0.000 description 1
- 229910052804 chromium Inorganic materials 0.000 description 1
- 239000011362 coarse particle Substances 0.000 description 1
- 230000000052 comparative effect Effects 0.000 description 1
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Description
本発明は、海洋構造物やラインパイプ、圧力容器等に用いられる高張力鋼とその製造方法に関し、特に、降伏応力が355MPa以上で、母材の強度・靭性に優れるだけでなく溶接部の靭性(CTOD特性)にも優れる高張力鋼とその製造方法に関するものである。 The present invention relates to a high-strength steel used for offshore structures, line pipes, pressure vessels, and the like, and its manufacturing method. In particular, the yield stress is 355 MPa or more, and the strength and toughness of the base metal are not only excellent, but also the toughness of the welded portion. The present invention relates to a high-strength steel excellent in (CTOD characteristics) and a method for producing the same.
海洋構造物等に用いられる鋼は、溶接接合して所望の形状の構造物に仕上げられるのが普通である。そのため、これらの鋼には、構造物の安全性を確保する観点から、母材自体の強度や靭性に優れることは勿論のこと、溶接継手の溶接部(溶接金属や熱影響部)の靭性にも優れていることが要求される。 Steel used for offshore structures and the like is usually finished by welding to a desired shape. Therefore, from the viewpoint of ensuring the safety of the structure, these steels are not only excellent in the strength and toughness of the base metal itself, but also in the toughness of the welded joint (welded metal or heat affected zone) of the welded joint. It is required to be excellent.
鋼の靭性の評価基準としては、従来、主にシャルピー試験による吸収エネルギーが用いられてきた。しかし、近年では、より信頼性を高めるために、き裂開口変位試験(Crack tip opening displacement test、以降「CTOD試験」と略記する)が用いられることが多い。この試験は、疲労予き裂を靭性評価部に発生させた試験片を3点曲げし、破壊直前のき裂底の口開き量(塑性変形量)を測定し、脆性破壊の発生抵抗を評価するものである。 As an evaluation standard for the toughness of steel, conventionally, absorbed energy mainly by Charpy test has been used. However, in recent years, a crack opening displacement test (hereinafter abbreviated as “CTOD test”) is often used in order to increase reliability. In this test, a test piece in which a fatigue precrack was generated in the toughness evaluation part was bent at three points, and the amount of opening (plastic deformation) at the crack bottom just before the fracture was measured to evaluate the resistance to brittle fracture. To do.
ところで、上記用途に用いられるような板厚が厚い鋼は、一般に、多層溶接が行われるが、このような溶接では、熱影響部は複雑な熱履歴を受けるため、局所脆化域が発生し易く、特にボンド部(溶接金属と母材との境界)や2相域再熱部(溶接1サイクル目で粗粒となり、2サイクル目でαとγの2相域に加熱される領域)の靭性の低下が大きいという問題がある。ボンド部は、溶融点直下の高温に曝されるため、オーステナイト粒が粗大化し、引き続く冷却により、脆弱な上部ベイナイト組織に変態し易いからである。また、ボンド部には、ウッドマンステッテン組織や島状マルテンサイトといった脆化組織が生成するため、靭性はさらに低下する。 By the way, the steel having a large thickness as used in the above-mentioned applications is generally subjected to multi-layer welding. However, in such welding, the heat-affected zone receives a complicated heat history, so that a local embrittlement region occurs. Easy, especially in the bond part (boundary between the weld metal and the base metal) and the two-phase region reheat part (the region that becomes coarse in the first cycle of welding and heated to the two-phase region of α and γ in the second cycle) There is a problem that the reduction in toughness is large. This is because the bond portion is exposed to a high temperature just below the melting point, so that austenite grains are coarsened and are easily transformed into a fragile upper bainite structure by subsequent cooling. In addition, since a brittle structure such as a woodman-stetten structure or island martensite is generated in the bond portion, the toughness is further reduced.
上記問題に対する対策として、例えば、鋼中にTiNを微細分散させて、オーステナイト粒の粗大化を抑制したり、フェライト変態核として利用したりする技術が実用化されている。さらに、特許文献1や特許文献2には、希土類元素(REM)をTiと複合添加して鋼中に微細粒子を分散させることにより、オーステナイト粒成長を抑制し、溶接部の靭性を向上する技術が開示されている。その他に、Tiの酸化物を分散させる技術や、BNのフェライト核生成能と酸化物分散とを組み合わせる技術、さらには、CaやREMを添加して硫化物の形態を制御することにより高靭性を得る技術も提案されている。 As a countermeasure against the above problem, for example, a technique of finely dispersing TiN in steel to suppress coarsening of austenite grains or to use as a ferrite transformation nucleus has been put into practical use. Furthermore, Patent Document 1 and Patent Document 2 disclose a technique for suppressing the austenite grain growth and improving the toughness of the welded portion by adding a rare earth element (REM) in combination with Ti and dispersing fine particles in the steel. Is disclosed. In addition, technology to disperse Ti oxide, technology to combine ferrite nucleation ability of BN and oxide dispersion, and addition of Ca and REM to control the form of sulfides to achieve high toughness Obtaining technology has also been proposed.
一方、上記2相域再熱部、即ち最初の溶接で溶融点直下の高温に曝された領域が、続く重ね溶接時の再加熱によりフェライトとオーステナイトの2相域となる領域が、最も脆化する原因は、2パス目以降の再加熱により、オーステナイト領域に炭素が濃化し、これが冷却中に、島状マルテンサイトを含む脆弱なベイナイト組織を生成し、靭性を低下させるからである。そこで、この対策として、低C、低Si化することにより島状マルテンサイトの生成を抑制し、さらにCuを添加することにより母材強度を確保する技術が開示されている(例えば、特許文献3参照)。 On the other hand, the two-phase region reheat zone, that is, the region exposed to a high temperature immediately below the melting point in the first welding, the region that becomes the two-phase region of ferrite and austenite by reheating during the subsequent lap welding is most brittle. The reason for this is that carbon is concentrated in the austenite region by reheating after the second pass, and this generates a fragile bainite structure including island martensite during cooling, thereby reducing toughness. Therefore, as a countermeasure, a technique is disclosed in which generation of island martensite is suppressed by reducing C and Si, and the strength of the base material is ensured by adding Cu (for example, Patent Document 3). reference).
しかしながら、上述した熱影響部の靭性が劣るという問題は、上記従来技術によってある程度の改善がなされたものの、まだ幾つかの解決すべき問題点が残されている。例えば、TiNを利用する技術では、TiNが溶解する温度域まで加熱されるボンド部においてはその作用がなくなり、さらに、固溶Tiおよび固溶Nによる基地組織の脆化によって著しい靭性の低下が起こることがある。また、Tiの酸化物を利用する技術では、酸化物の微細分散が十分均質にできないという問題がある。さらに、近年、海洋構造物や船舶等が大型化していくのに伴って、それに用いられる鋼材は、より高強度化、厚肉化が進められている。それらの課題を達成するには、特許文献3の技術とは逆に、合金元素を多量に添加することが有効である。しかし、合金元素の多量添加は、熱影響部の靭性の低下を招くという問題点を有している。 However, although the above-described problem of poor toughness of the heat affected zone has been improved to some extent by the above prior art, some problems to be solved still remain. For example, in the technology using TiN, the action is lost in the bond portion heated to a temperature range where TiN dissolves, and further, the toughness is significantly lowered due to the embrittlement of the base structure due to solute Ti and solute N. Sometimes. Further, the technology using Ti oxide has a problem that fine dispersion of the oxide cannot be sufficiently homogeneous. Furthermore, in recent years, as marine structures, ships, and the like are increased in size, steel materials used therefor have been increased in strength and thickness. In order to achieve these problems, it is effective to add a large amount of alloy elements, contrary to the technique of Patent Document 3. However, the addition of a large amount of alloy elements has a problem in that the toughness of the heat-affected zone is reduced.
そこで、本発明の目的は、従来技術が抱える上記問題点を解決し、合金元素の添加量を増やすことなく、母材の強度・靭性に優れるとともに、溶接熱影響部の靭性にも優れる高張力鋼とその有利な製造方法を提案することにある。 Therefore, the object of the present invention is to solve the above-mentioned problems of the prior art, and to increase the tensile strength of the base metal and the toughness of the weld heat affected zone without increasing the amount of alloy elements added, It is to propose a steel and its advantageous production method.
発明者らは、高張力鋼の母材強度・靭性を向上すると共に、溶接熱影響部の靭性をも改善することができる方法について鋭意検討した。その結果、溶接熱影響部の靭性劣化は、脆化組織の生成に起因していることから、溶接熱影響部の靭性を向上させるためには、溶接時に高温加熱される領域におけるオーステナイト粒の粗大化を抑制したうえで、さらに、溶接後の冷却時にフェライト変態を促進させるための変態核を均一微細に分散させてやることが有効であることがわかった。 The inventors diligently studied a method capable of improving the base material strength and toughness of high-tensile steel and also improving the toughness of the weld heat affected zone. As a result, the deterioration of the toughness of the weld heat affected zone is caused by the formation of an embrittled structure. Therefore, in order to improve the toughness of the weld heat affected zone, the coarseness of austenite grains in the region heated at high temperature during welding Further, it was found that it is effective to uniformly disperse transformation nuclei for promoting ferrite transformation during cooling after welding.
そこで、発明者らは、上記脆化組織の生成を抑制する方法についてさらに検討した結果、硫化物の形態制御のために添加しているCaの添加量を適正範囲に制御することが有効であること、また、溶接熱影響部の靭性(CTOD特性)を向上するには、Niの添加が有効であることを見出した。 Thus, as a result of further study on the method for suppressing the formation of the embrittlement structure, the inventors are effective to control the addition amount of Ca added for the morphology control of the sulfide within an appropriate range. In addition, it has been found that the addition of Ni is effective for improving the toughness (CTOD characteristics) of the weld heat affected zone.
また、母材の強度・靭性に及ぼす圧延条件の影響について検討したところ、圧延後の冷却を、冷却速度が大きい前段冷却と小さい後段冷却とからなる2段冷却とし、それぞれの冷却速度を適正に制御すれば、鋼板組織がアシキュラーフェライト主体の組織となり、母材の強度・靭性に優れた高張力鋼を製造できることを見出した。さらに、母材の強度と靭性をより高めるには、オーステナイトの低温域で、未再結晶域を形成する効果が大きいNbを有効利用することが重要であり、そのためには、Nの含有量の上限を、従来よりも厳しく規制することが必要であることを見出し、本発明を完成させた。 In addition, when the influence of rolling conditions on the strength and toughness of the base metal was examined, the cooling after rolling was changed to two-stage cooling consisting of a pre-stage cooling with a large cooling rate and a small post-stage cooling, and each cooling rate was appropriately set. It was found that if controlled, the steel sheet structure becomes a structure mainly composed of acicular ferrite, and high strength steel excellent in the strength and toughness of the base material can be produced. Furthermore, in order to further increase the strength and toughness of the base material, it is important to effectively use Nb, which has a large effect of forming a non-recrystallized region in the low temperature region of austenite. The present inventors have found that it is necessary to regulate the upper limit more strictly than in the past, and completed the present invention.
すなわち、本発明は、C:0.05〜0.1mass%、Si:0.05〜0.5mass%、Mn:1〜2mass%、P:0.015mass%以下、S:0.005mass%以下、Al:0.005〜0.06mass%、Ni:0.3〜2mass%、Nb:0.004〜0.05mass%、Ti:0.005〜0.02mass%、N:0.003mass%未満、Ca:0.0005〜0.003mass%を含有し、
Ca,SおよびOが下記(1)式;
0<(Ca−(0.18+130×Ca)×O)/1.25/S<1 ・・・(1)
ここで、Ca,SおよびOは、各元素の含有量(mass%)
を満たして含有し、残部がFeおよび不可避的不純物からなる成分組成を有することを特徴とする高張力鋼である。
That is, the present invention is C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less. , Al: 0.005 to 0.06 mass%, Ni: 0.3 to 2 mass%, Nb: 0.004 to 0.05 mass%, Ti: 0.005 to 0.02 mass%, N: less than 0.003 mass% Ca: 0.0005 to 0.003 mass%,
Ca, S and O are represented by the following formula (1):
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%).
Is a high-strength steel characterized in that it has a component composition consisting of Fe and inevitable impurities .
本発明の高張力鋼は、上記成分組成に加えてさらに、B:0.0003〜0.0025mass%、Mo:0.7mass%以下、V:0.022mass%以下、Cu:1mass%以下およびCr:0.7mass%以下の中から選ばれる1種または2種以上を含有することを特徴とする。
In addition to the above component composition, the high-tensile steel of the present invention further includes B: 0.0003 to 0.0025 mass%, Mo: 0.7 mass% or less, V: 0.022 mass% or less, Cu: 1 mass% or less, and Cr: One or more selected from 0.7 mass% or less are contained.
また、本発明は、C:0.05〜0.1mass%、Si:0.05〜0.5mass%、Mn:1〜2mass%、P:0.015mass%以下、S:0.005mass%以下、Al:0.005〜0.06mass%、Ni:0.3〜2mass%、Nb:0.004〜0.05mass%、Ti:0.005〜0.02mass%、N:0.003mass%未満、Ca:0.0005〜0.003mass%を含有し、
Ca,SおよびOが下記(1)式;
0<(Ca−(0.18+130×Ca)×O)/1.25/S<1 ・・・(1)
ここで、Ca,SおよびOは、各元素の含有量(mass%)
を満たして含有し、残部がFeおよび不可避的不純物からなる成分組成を有する鋼スラブを1050〜1200℃に加熱後、950℃以上の温度域における累積圧下率が30%以上、950℃未満の温度域における累積圧下率が30〜70%となる熱間圧延を施し、その後、熱間圧延終了温度から600〜450℃間の冷却停止温度までを5〜20℃/sで冷却する前段冷却と、前段冷却停止温度から450℃未満〜200℃間の冷却停止温度までを1〜5℃未満/sで冷却する後段冷却を施すことを特徴とする高張力鋼の製造方法を提案する。
In the present invention, C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less , Al: 0.005 to 0.06 mass%, Ni: 0.3 to 2 mass%, Nb: 0.004 to 0.05 mass%, Ti: 0.005 to 0.02 mass%, N: less than 0.003 mass% Ca: 0.0005 to 0.003 mass%,
Ca, S and O are represented by the following formula (1):
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%).
After heating a steel slab having a composition composed of Fe and unavoidable impurities to 1050 to 1200 ° C., the cumulative rolling reduction in a temperature range of 950 ° C. or higher is 30% or more and less than 950 ° C. Pre-stage cooling in which the cumulative rolling reduction in the region is 30 to 70%, and thereafter cooling from 5 to 20 ° C./s from the hot rolling end temperature to the cooling stop temperature between 600 to 450 ° C., We propose a method for producing high-strength steel, characterized in that post-stage cooling is performed in which cooling is carried out at a temperature of less than 1 to 5 ° C./s from the temperature of the former stage of cooling to a cooling stop temperature of less than 450 ° C. to 200 ° C.
本発明の高張力鋼の製造方法における上記鋼スラブは、上記成分組成に加えてさらに、B:0.0003〜0.0025mass%、Mo:0.7mass%以下、V:0.022mass%以下、Cu:1mass%以下およびCr:0.7mass%以下の中から選ばれる1種または2種以上を含有することを特徴とする。
The steel slab in the method for producing high-tensile steel according to the present invention is further provided with B: 0.0003 to 0.0025 mass%, Mo: 0.7 mass% or less, V: 0.022 mass% or less in addition to the above component composition. Cu: 1% by mass or less and Cr: 0.7% by mass or less selected from 1 type or 2 types or more.
また、本発明の高張力鋼の製造方法は、上記後段冷却後の鋼に、450〜650℃で焼戻処理を施すことを特徴とする。 Moreover, the manufacturing method of the high-tensile steel of this invention is characterized by performing a tempering process at 450-650 degreeC to the steel after the said back | latter stage cooling.
本発明によれば、母材が、降伏応力が355MPa以上で靭性に優れると共に、溶接熱影響部の靭性(CTOD特性)にも優れる高強度鋼を安価に製造することができるので、海洋構造物や船舶等の大型化に大きく寄与する。 According to the present invention, since the base material can produce a high-strength steel having a yield stress of 355 MPa or more and excellent toughness and also excellent toughness (CTOD characteristics) of the weld heat affected zone, it can be manufactured at low cost. This greatly contributes to increasing the size of ships and ships.
本発明の基本的な技術思想について説明する。
本発明の第1の特徴は、溶接熱影響部の靭性を向上するために、硫化物の形態制御を目的として添加するCaの化合物(CaS)の晶出を有効利用するところにある。このCaSは、酸化物に比べて低温で晶出するため、均一に微細分散することができる。そして、CaSの添加量および添加時の溶鋼中の溶存酸素量を適性範囲に制御することによって、CaS晶出後でも固溶Sが確保され、CaSの表面上にMnSが析出して複合硫化物を形成する。このMnSには、フェライト核生成能があることが知られており、さらに、析出したMnSの周囲には、Mnの希薄帯が形成されるので、フェライト変態がより促進される。しかも、析出したMnS上には、TiN,BN,AlN等のフェライト生成核も析出するので、よりいっそうフェライト変態が促進される。
The basic technical idea of the present invention will be described.
The first feature of the present invention is to effectively use the crystallization of a Ca compound (CaS) added for the purpose of controlling the form of sulfides in order to improve the toughness of the weld heat affected zone. Since this CaS crystallizes at a lower temperature than the oxide, it can be uniformly finely dispersed. And by controlling the amount of CaS added and the amount of dissolved oxygen in the molten steel at the time of addition to an appropriate range, solid solution S is ensured even after crystallization of CaS, and MnS precipitates on the surface of CaS to form composite sulfide. Form. This MnS is known to have a ferrite nucleation ability. Further, since a Mn dilute band is formed around the precipitated MnS, the ferrite transformation is further promoted. Moreover, since ferrite-forming nuclei such as TiN, BN, and AlN also precipitate on the precipitated MnS, the ferrite transformation is further promoted.
上記技術によって、高温でも溶解しないフェライト変態生成核を微細に分散させることが可能となり、溶接熱影響部の組織を微細なフェライトパーライト化して、高靭性化を達成することができる。また、多層溶接時の熱サイクルにより2相域に再加熱される領域においても、最初の溶接による熱影響部の組織が微細化されるので、未変態の領域の靭性が向上し、さらに、再変態するオーステナイト粒も微細化するので、靭性の低下の度合いを小さく抑制することができる。 According to the above technique, it becomes possible to finely disperse the ferrite transformation formation nuclei that do not dissolve even at high temperatures, and the structure of the weld heat affected zone can be made into fine ferrite pearlite, thereby achieving high toughness. Even in the region that is reheated to the two-phase region due to the thermal cycle during multi-layer welding, the structure of the heat-affected zone by the first welding is refined, so that the toughness of the untransformed region is improved, and Since the transformed austenite grains are also refined, the degree of reduction in toughness can be suppressed small.
本発明の第2の特徴は、母材の強度・靭性を向上するために、素材成分として、N量を0.003mass%未満に制限する点にある。母材の強度・靭性を向上するためにはNbの添加が有効であるが、Nbは、窒化物中に容易に固溶するため、窒化物が多く存在する場合には、この窒化物中に多くのNbが固溶してしまい、強度や靭性向上に有効な固溶Nb量を確保できなくなってしまう。そこで、N量を0.003mass%未満に抑えて窒化物の生成を抑制することにより、Nbの効果を最大限に利用することができ、ひいては、母材の強度・靭性の向上を図ることができる。また、N量を低減することにより、連続鋳造鋳片の表面欠陥が減少し、製品歩留まりが向上するという効果も得ることが出来る。 The second feature of the present invention is that the amount of N is limited to less than 0.003 mass% as a raw material component in order to improve the strength and toughness of the base material. Addition of Nb is effective to improve the strength and toughness of the base metal. However, Nb is easily dissolved in the nitride, so if a large amount of nitride exists, Much Nb is dissolved, and it becomes impossible to secure a solid solution Nb amount effective for improving strength and toughness. Therefore, by suppressing the amount of N to less than 0.003 mass% and suppressing the formation of nitride, the effect of Nb can be utilized to the maximum, and consequently the strength and toughness of the base material can be improved. it can. Further, by reducing the N amount, the surface defects of the continuous cast slab are reduced, and the product yield can be improved.
本発明の第3の特徴は、鋼材圧延後の冷却を、前段冷却と後段冷却の2段階に分けて、後段冷却より前段冷却の冷却速度を大きく制御するところにある。この点について、実験結果を基に説明する。
C:0.08mass%、Si:0.2mass%、Mn:1.4mass%、Ni:0.4mass%を基本成分とする鋼スラブを、1150℃に加熱後、950℃以上の累積圧下率を40%、950℃未満での累積圧下率を50%、圧延終了温度を850℃とする熱間圧延を行い、その後、圧延終了温度から500℃までを冷却速度2〜25℃/sで冷却する前段冷却した後、さらに、350℃までを冷却速度3℃/sで冷却する後段冷却を行い、その後、空冷して板厚が10〜50mmtの厚鋼板とした。この厚鋼板について、アシキュラーフェライト組織の面積率、引張強度特性および−40℃における靭性(シャルピー吸収エネルギー)を測定した。
The third feature of the present invention lies in that the cooling after rolling the steel material is divided into two stages of the pre-stage cooling and the post-stage cooling, and the cooling rate of the pre-stage cooling is controlled to be greater than the post-stage cooling. This point will be described based on experimental results.
C: 0.08 mass%, Si: 0.2 mass%, Mn: 1.4 mass%, Ni: 0.4 mass% After heating the steel slab to 1150 ° C, the cumulative reduction ratio of 950 ° C or higher Hot rolling is performed at 40%, a cumulative reduction ratio of less than 950 ° C. is 50%, and a rolling end temperature is 850 ° C., and then the rolling end temperature to 500 ° C. is cooled at a cooling rate of 2 to 25 ° C./s. After the pre-stage cooling, further post-stage cooling was performed by cooling to 350 ° C. at a cooling rate of 3 ° C./s, and then air cooling was performed to obtain a thick steel plate having a plate thickness of 10 to 50 mmt. About this thick steel plate, the area ratio of the acicular ferrite structure, the tensile strength characteristics, and the toughness (Charpy absorbed energy) at −40 ° C. were measured.
一般に、フェライト−パーライト組織からなる高強度の組織に変化させる場合、島状マルテンサイトなどをラス間に含む比較的粗大な上部ベイナイト組織となり、靭性が大きく低下する。そこで、高強度と高靭性を両立させるためには、圧延条件の工夫などにより微細なアシキュラーフェライト組織とすることが必要となる。
図1は、母材強度およびアシキュラーフェライト面積率に及ぼす前段冷却の冷却速度の影響を示したものである。この図から、前段冷却の冷却速度が増すのに伴って強度が上昇し、靭性が低下する傾向があることがわかる。一方、アシキュラーフェライト組織の面積率は、冷却速度の増大とともに上昇するが、おおよそ10℃/s以上では上昇勾配が緩やかになること、すなわち、前段冷却の冷却速度をある速度以上に高めることにより、比較的高温で生成するポリゴナルフェライトの生成を抑制して、アシキュラーフェライト主体の組織とし、強度−靭性バランスに優れた鋼板を製造できることがわかった。
一方、後段冷却速度は、前段冷却速度より早いと島状マルテンサイトを生成し、母材の靭性を劣化させる。ただし、遅すぎると母材の強度が低下してしまうことから、適正な範囲に制御する必要があることもわかった。
In general, when the structure is changed to a high-strength structure composed of a ferrite-pearlite structure, a relatively coarse upper bainite structure including island martensite and the like between the laths is formed, and the toughness is greatly reduced. Therefore, in order to achieve both high strength and high toughness, it is necessary to obtain a fine acicular ferrite structure by devising rolling conditions.
FIG. 1 shows the influence of the cooling rate of the pre-cooling on the base material strength and the acicular ferrite area ratio. From this figure, it can be seen that the strength increases and the toughness tends to decrease as the cooling rate of the pre-stage cooling increases. On the other hand, the area ratio of the acicular ferrite structure increases with an increase in the cooling rate. However, when the cooling rate increases at about 10 ° C./s or more, the rising gradient becomes gentle, that is, by increasing the cooling rate of the preceding stage cooling to a certain rate or more. It was found that a steel sheet excellent in strength-toughness balance can be produced by suppressing the formation of polygonal ferrite produced at a relatively high temperature to form a structure mainly composed of acicular ferrite.
On the other hand, when the latter cooling rate is faster than the former cooling rate, island martensite is generated, and the toughness of the base material is deteriorated. However, it has also been found that if it is too slow, the strength of the base material will be reduced, so that it is necessary to control within an appropriate range.
次に、本発明に係る高張力鋼の成分組成を限定する理由について説明する。
C:0.05〜0.1mass%
Cは、鋼の強度に最も大きく影響する元素であり、構造用鋼として必要な強度(YS≧355MPa)を確保するためには0.05mass%以上含有させる必要がある。しかし、逆に、多過ぎると、溶接割れを引き起こすので、上限を0.1mass%とする。
Next, the reason for limiting the component composition of the high-strength steel according to the present invention will be described.
C: 0.05-0.1 mass%
C is an element that has the greatest influence on the strength of the steel. In order to ensure the strength necessary for structural steel (YS ≧ 355 MPa), it is necessary to contain 0.05 mass% or more. However, conversely, if too much, weld cracking is caused, so the upper limit is made 0.1 mass%.
Si:0.05〜0.5mass%
Siは、脱酸剤として添加される成分であり、0.05mass%以上添加する必要がある。一方、0.5mass%を超えると、母材の靭性を低下させるため0.5mass%以下とする必要がある。
Si: 0.05-0.5 mass%
Si is a component added as a deoxidizer, and it is necessary to add 0.05 mass% or more. On the other hand, when it exceeds 0.5 mass%, it is necessary to make it 0.5 mass% or less in order to reduce the toughness of the base material.
Mn:1〜2mass%
Mnは、母材の強度を確保するため、1mass%以上添加する必要がある。一方、2mass%を超えると、溶接部の靭性を著しく低下させるため、2mass%以下とする必要がある。好ましくは、1.2〜1.8mass%の範囲である。
Mn: 1 to 2 mass%
Mn needs to be added in an amount of 1 mass% or more in order to ensure the strength of the base material. On the other hand, if it exceeds 2 mass%, the toughness of the welded portion is remarkably lowered, so it is necessary to set it to 2 mass% or less. Preferably, it is in the range of 1.2 to 1.8 mass%.
P:0.015mass%以下
Pは、不可避的に混入する不純物であり、0.015mass%を超えると、溶接部の靭性を低下させるため、0.015mass%以下に制限する。好ましくは、0.012mass%以下である。
P: 0.015 mass% or less P is an impurity that is inevitably mixed. If it exceeds 0.015 mass%, the toughness of the welded portion is lowered, so that it is limited to 0.015 mass% or less. Preferably, it is 0.012 mass% or less.
S:0.005mass%以下
Sは、不可避的に混入する不純物であり、0.005mass%を超えて含有すると、母材および溶接部の靭性を低下させるため、0.005mass%以下とする。好ましくは、0.0035mass%以下である。
S: 0.005 mass% or less S is an impurity that is inevitably mixed, and if contained in excess of 0.005 mass%, the toughness of the base metal and the welded portion is lowered, so the content is made 0.005 mass% or less. Preferably, it is 0.0035 mass% or less.
Al:0.005〜0.06mass%
Alは、溶鋼を脱酸するために添加される元素であり、0.005mass%以上含有させる必要がある。一方、0.06mass%を超えて添加すると、母材の勒性を低下させるとともに、溶接による希釈によって溶接金属部に混入し、靭性を低下させるため、0.06mass%以下に制限する必要がある。
Al: 0.005-0.06 mass%
Al is an element added to deoxidize molten steel, and it is necessary to contain 0.005 mass% or more. On the other hand, if added over 0.06 mass%, the toughness of the base metal is reduced and mixed into the weld metal part by dilution by welding to reduce toughness. Therefore, it is necessary to limit to 0.06 mass% or less. .
Ni:0.3〜2mass%
Niは、鋼の強度および溶接熱影響部のCTOD特性の向上に有効な元素である。この効果は、0.3mass%以上の添加によって発現する。しかし、Niは、高価な元素であるため、上限を2mass%とする。
Ni: 0.3-2 mass%
Ni is an element effective for improving the strength of steel and the CTOD characteristics of the weld heat affected zone. This effect is manifested by addition of 0.3 mass% or more. However, since Ni is an expensive element, the upper limit is set to 2 mass%.
Nb:0.004〜0.05mass%
Nbは、オーステナイトの低温度域で、未再結晶域を形成するので、その温度域で圧延を施すことにより、母材組織の微細化、高靭性化を図ることができる。また、圧延・冷却後に焼戻処理を施すことにより、析出強化を図ることもできる。したがって、Nbは、鋼の強化の観点からは重要な添加元素である。上記効果を得るためには、Nbは0.004mass%以上添加する必要がある。しかし、Nbを0.05mass%を超えて過剰に添加した場合には、溶接部の靭性を劣化させるので、上限は0.05mass%とする。
Nb: 0.004 to 0.05 mass%
Since Nb forms a non-recrystallized region in the low temperature range of austenite, the base material structure can be refined and toughened by rolling in that temperature range. In addition, precipitation strengthening can be achieved by performing a tempering treatment after rolling and cooling. Therefore, Nb is an important additive element from the viewpoint of strengthening steel. In order to acquire the said effect, Nb needs to add 0.004 mass% or more. However, when Nb is added excessively in excess of 0.05 mass%, the toughness of the welded portion is deteriorated, so the upper limit is made 0.05 mass%.
Ti:0.005〜0.02mass%
Tiは、溶鋼が凝固する際にTiNとなって析出し、溶接部におけるオーステナイトの粗大化抑制やフェライト変態核となって、高靭性化に寄与する。0.005mass%未満ではその効果が小さく、一方、0.02mass%を超えると、TiN粒子の粗大化によって期待した効果が得られなくなる。よって、Tiの添加量は、0.005〜0.02mass%の範囲とする。
Ti: 0.005-0.02 mass%
Ti precipitates as TiN when the molten steel is solidified, and serves as a suppression of coarsening of austenite and ferrite transformation nuclei in the welded portion, thereby contributing to high toughness. If it is less than 0.005 mass%, the effect is small. On the other hand, if it exceeds 0.02 mass%, the effect expected by the coarsening of TiN particles cannot be obtained. Therefore, the amount of Ti added is in the range of 0.005 to 0.02 mass%.
N:0.003mass%未満
Nは、母材の強度と靭性の向上に必要な固溶Nb量を確保するために、0.003mass%未満とする必要がある。上述したように、Nbは、母材の組織微細化、高靭性化や析出強化に有効な元素である。これらの効果を得るためには、Nbが圧延前の加熱時に固溶状態である必要がある。しかし、NbとTiが同時に添加された場合、(Ti,Nb)(C,N)複合炭窒化物を形成し、この析出物は、NbCに比べて高温まで安定に存在するため、圧延時の加熱段階で、一部が溶解せずに残存して固溶Nbが減少する。さらに、N量が増加すると、(Ti,Nb)(C,N)は、より溶け難くなる傾向があるため、固溶Nb量はさらに減少する。そこで、固溶Nbの効果を最大限に有効利用するために、溶接部の靭性確保に必要な量のTiを添加した上で、N量の上限を制限し、本発明では、Nの上限を0.003mass%とする。
N: Less than 0.003 mass% N is required to be less than 0.003 mass% in order to ensure the amount of solute Nb necessary for improving the strength and toughness of the base material. As described above, Nb is an element effective for refining the structure of the base material, increasing the toughness, and precipitation strengthening. In order to obtain these effects, Nb needs to be in a solid solution state during heating before rolling. However, when Nb and Ti are added at the same time, a (Ti, Nb) (C, N) composite carbonitride is formed, and this precipitate exists stably up to a high temperature as compared with NbC. In the heating stage, a part of the solid solution remains without being dissolved, and the solid solution Nb is reduced. Furthermore, when the amount of N increases, (Ti, Nb) (C, N) tends to be more difficult to dissolve, so the amount of solid solution Nb further decreases. Therefore, in order to effectively use the effect of the solute Nb, the upper limit of the N amount is limited after adding an amount of Ti necessary for ensuring the toughness of the welded portion. 0.003 mass%.
Ca:0.0005〜0.003mass%
Caは、Sを固定して、靭性を向上する効果を有する。この効果を発現させるためには、少なくとも0.0005mass%添加する必要がある。しかし、0.003mass%以上含有しても、その効果が飽和するので、Caは、0.0005〜0.003mass%の範囲とする。
Ca: 0.0005 to 0.003 mass%
Ca has an effect of fixing S and improving toughness. In order to exhibit this effect, it is necessary to add at least 0.0005 mass%. However, even if it contains 0.003 mass% or more, since the effect is saturated, Ca is made into the range of 0.0005 to 0.003 mass%.
0<(Ca−(0.18+130×Ca)×O)/1.25/S<1
高温でも溶解しないフェライト変態生成核CaSを微細分散させるためには、Ca,SおよびOは、下記(1)式;
0<(Ca−(0.18+130×Ca)×O)/1.25/S<1 ・・・(1)
ここで、Ca,S,O:各元素の含有量(mass%)
の関係を満たして含有する必要がある。上記式中の、(Ca−(0.18+130×Ca)×O)/(1.25/S)は、硫化物形態制御に有効なCaとSの原子濃度の比を示す値であり、この価から、硫化物の形態を推定することができる(持田他、「鉄と鋼」、日本鉄鋼協会、第66年(1980)、第3号、P354〜362)。(Ca−(0.18+130×Ca)×O)/(1.25/S)≦0の場合には、CaSが晶出しないため、Sは、MnS単独の形態で析出するので、本発明の主眼である、溶接熱影響部でのフェライト生成核の微細分散を実現することができない。また、単独で析出したMnSは、鋼板圧延時に伸長されて、母材の靭性低下を引き起こす。また、(Ca−(0.18+130×Ca)×O)/(1.25/S)≧1の場合には、Sが完全にCaによって固定され、フェライト生成核として働くMnSがCaS上に析出しなくなるため、複合硫化物が、フェライト生成核として十分に機能することができなくなる。これに対して、Ca,S,Oが、上記(1)式を満たした場合には、CaS上にMnSが析出して複合硫化物を形成し、フェライト生成核として有効に機能することができる。なお、((Ca−(0.18+130×Ca)×O)/(1.25/S)は、好ましくは0.2〜0.8の範囲である。
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1
In order to finely disperse the ferrite transformation nuclei CaS that does not dissolve even at high temperatures, Ca, S and O are represented by the following formula (1):
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S, O: Content of each element (mass%)
It is necessary to contain and satisfy the relationship. In the above formula, (Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) is a value indicating the ratio of the atomic concentrations of Ca and S effective for sulfide morphology control, and this The form of sulfide can be estimated from the value (Mochida et al., “Iron and Steel”, Japan Iron and Steel Institute, 66th (1980), No. 3, P354-362). In the case of (Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) ≦ 0, since CaS does not crystallize, S precipitates in the form of MnS alone. It is not possible to achieve fine dispersion of ferrite nuclei at the weld heat affected zone, which is the main focus. Further, MnS precipitated alone is elongated during rolling of the steel sheet, causing a reduction in the toughness of the base material. In addition, when (Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) ≧ 1, S is completely fixed by Ca, and MnS acting as a ferrite nucleation precipitates on CaS. Therefore, the composite sulfide cannot sufficiently function as a ferrite formation nucleus. On the other hand, when Ca, S, and O satisfy the above formula (1), MnS precipitates on CaS to form a composite sulfide, and can effectively function as a ferrite nuclei. . In addition, ((Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) is preferably in the range of 0.2 to 0.8.
本発明の高張力鋼は、上記必須成分に加えてさらに、強度および靭性を高めるために、B,V,Cu,CrおよびMoから選ばれる1種または2種以上を含有させることができる。
B:0.0003〜0・0025mass%
Bは、オーステナイト粒界に偏析し、粒界から起こるフェライト変態を抑制してベイナイト組織の分率を高めることにより、鋼を高強度化する効果がある。その効果は、0.0003mass%以上の添加で得ることができる。しかし、0.0025mass%を超えて添加すると、逆に靭性が低下する。Bを添加する場合、より好ましい範囲は0.0005〜0.002mass%である。
In addition to the above essential components, the high strength steel of the present invention can further contain one or more selected from B, V, Cu, Cr and Mo in order to increase the strength and toughness.
B: 0.0003-0.0025 mass%
B segregates at the austenite grain boundaries and has the effect of increasing the strength of the steel by suppressing the ferrite transformation that occurs from the grain boundaries and increasing the fraction of the bainite structure. The effect can be obtained by adding 0.0003 mass% or more. However, if added over 0.0025 mass%, the toughness is conversely reduced. When adding B, a more preferable range is 0.0005 to 0.002 mass%.
V:0.2mass%以下
Vは、母材の強度・靭性の向上に有効な元素であり、また、VNとして析出してフェライト生成核としても働く元素でもある。しかし、添加量が0.2mass%を超えると、却って靭性の低下を招くので0.2mass%以下添加するのが好ましい。より好ましくは、0.15mass%以下である。
V: 0.2 mass% or less V is an element effective for improving the strength and toughness of the base material, and is also an element that precipitates as VN and also serves as a ferrite formation nucleus. However, if the addition amount exceeds 0.2 mass%, the toughness is deteriorated on the contrary. Therefore, it is preferable to add 0.2 mass% or less. More preferably, it is 0.15 mass% or less.
Cu:1mass%以下
Cuは、Niと同様の効果を有する元素であるが、1mass%を超えると、熱間脆性を引き起こして鋼板の表面性状を劣化させるため、1mass%以下の範囲で添加するのが好ましい。より好ましくは、0.8mass%以下である。
Cu: 1 mass% or less Cu is an element having the same effect as Ni. However, if it exceeds 1 mass%, it causes hot brittleness and deteriorates the surface properties of the steel sheet, so it is added in a range of 1 mass% or less. Is preferred. More preferably, it is 0.8 mass% or less.
Cr:0.7mass%以下
Crは、母材を高強度化するのに有効な元素であるが、多量に添加すると、逆に靭性に悪影響を与えるので、上限を0.7mass%とするのが好ましい。より好ましくは、0.5mass%以下である。
Cr: 0.7 mass% or less Cr is an element effective for increasing the strength of the base material. However, if added in a large amount, it adversely affects toughness, so the upper limit should be 0.7 mass%. preferable. More preferably, it is 0.5 mass% or less.
Mo:0.7mass%以下
Moは、Crと同様、母材を高強度化するのに有効な元素であるが、多量に添加すると、逆に靭性に悪影響を与えるので、上限を0.7mass%とするのが好ましい。より好ましくは、0.5mass%以下である。
Mo: 0.7 mass% or less Mo, like Cr, is an element effective for increasing the strength of the base material, but if added in a large amount, adversely affects toughness, so the upper limit is 0.7 mass%. Is preferable. More preferably, it is 0.5 mass% or less.
次に、本発明の高張力鋼の製造方法について説明する。
上述した本発明に適合する成分組成に調整した溶鋼を、転炉、電気炉、真空溶解炉等を用いて通常の方法で溶製し、連続鋳造または造塊−分塊圧延など通常の工程を経てスラブ等の鋼素材とする。この鋼素材を熱間圧延して厚肉の高張力鋼を得るが、この際、熱間圧延に先立って行う鋼素材の加熱温度は1050〜1200℃の範囲とする必要がある。1050℃以上に加熱するのは、鋼素材中に存在する鋳造欠陥を、熱間圧延によって確実に圧着させるためである。しかし、1200℃を超える温度に加熱すると、凝固時に析出したTiNが粗大化し、溶接部の靭性が低下するため、加熱温度は1200℃以下に規制する必要がある。
Next, the manufacturing method of the high strength steel of this invention is demonstrated.
The above-described molten steel adjusted to the component composition suitable for the present invention is melted by a normal method using a converter, an electric furnace, a vacuum melting furnace, etc., and a normal process such as continuous casting or ingot-bundling rolling is performed. After that, steel material such as slab is used. The steel material is hot-rolled to obtain a thick high-strength steel. At this time, the heating temperature of the steel material prior to the hot-rolling needs to be in the range of 1050 to 1200 ° C. The reason for heating to 1050 ° C. or higher is to ensure that the casting defects present in the steel material are crimped by hot rolling. However, when heated to a temperature exceeding 1200 ° C., TiN deposited during solidification becomes coarse and the toughness of the welded portion decreases, so the heating temperature needs to be regulated to 1200 ° C. or less.
上記温度に加熱した鋼素材は、その後、950℃以上の温度域における累積圧下率を30%以上とする熱間圧延と、950℃未満の温度域における累積圧下率を30〜70%とする熱間圧延を施し、所定の板厚を有する高張力鋼とする。950℃以上の温度域で累積圧下率30%以上の熱間圧延を施す理由は、この温度域で累積圧下率が30%以上の圧下を加えた場合には、オーステナイト粒が再結晶して組織を微細化できるのに対し、累積圧下率が30%未満では、加熱時に生成した異常粗大粒が残存して、母材の靭性に悪影響を及ぼすためである。 The steel material heated to the above temperature is then subjected to hot rolling in which the cumulative reduction ratio in the temperature range of 950 ° C. or higher is 30% or more, and heat in which the cumulative reduction ratio in the temperature range of less than 950 ° C. is 30 to 70%. Hot rolling is performed to obtain a high-strength steel having a predetermined thickness. The reason why hot rolling with a cumulative reduction ratio of 30% or more is performed in a temperature range of 950 ° C. or higher is that when a reduction with a cumulative reduction ratio of 30% or more is applied in this temperature range, the austenite grains recrystallize. This is because, when the cumulative rolling reduction is less than 30%, abnormal coarse particles generated during heating remain, which adversely affects the toughness of the base material.
また、950℃未満の温度域における累積圧下率を30〜70%とする熱間圧延を行う理由は、この温度域で圧延されたオーステナイト粒は十分再結晶しないため、圧延後のオーステナイト粒は、扁平に変形したままで、内部に変形帯などの欠陥に多量に含む内部歪の高いものとなる。そして、この蓄積された内部エネルギーが、その後のフェライト変態の駆動力として働き、フェライト変態を促進するからである。しかし、圧下率が30%未満では、上記の蓄積される内部エネルギーが十分ではないため、フェライト変態が起こりにくく、ベイナイト組織が生成する。一方、圧下率が70%を超えると、逆にポリゴナルフェライトの生成が促進され、アシキュラーフェライトの生成が抑制されるからである。 Also, the reason for performing hot rolling with a cumulative reduction ratio of 30 to 70% in a temperature range of less than 950 ° C. is that austenite grains rolled in this temperature range are not sufficiently recrystallized. It remains flat and has a high internal strain that is contained in a large amount in defects such as a deformation zone. This is because the accumulated internal energy acts as a driving force for the subsequent ferrite transformation and promotes the ferrite transformation. However, if the rolling reduction is less than 30%, the accumulated internal energy is not sufficient, so that ferrite transformation hardly occurs and a bainite structure is generated. On the other hand, if the rolling reduction exceeds 70%, the formation of polygonal ferrite is conversely accelerated and the formation of acicular ferrite is suppressed.
続く熱間圧延終了後の冷却は、前段冷却と後段冷却に分け、前者の冷却速度を後者のそれよりも相対的に大きくする、すなわち、前段冷却では、熱間圧延終了温度から600〜450℃間の冷却停止温度まで好ましくは熱間圧延終了温度から580〜480℃間の冷却停止温度までを、5〜20℃/s、好ましくは6〜16℃/sの冷却速度で冷却し、その後の後段冷却では、前段冷却の停止温度から450℃未満〜200℃間の後段冷却停止温度まで好ましくは前段冷却の停止温度から400〜300℃間の冷却停止温度までを、1〜5℃未満/s、好ましくは2〜4℃/sの冷却速度で冷却する必要がある。 Cooling after the end of the subsequent hot rolling is divided into pre-stage cooling and post-stage cooling, and the cooling rate of the former is made relatively larger than that of the latter, that is, in the pre-stage cooling, 600 to 450 ° C. from the hot rolling end temperature. Until the cooling stop temperature between, preferably from the hot rolling end temperature to the cooling stop temperature between 580 to 480 ° C. is cooled at a cooling rate of 5 to 20 ° C./s, preferably 6 to 16 ° C./s, and thereafter In the latter-stage cooling, from the first-stage cooling stop temperature to the second-stage cooling stop temperature between less than 450 ° C. and 200 ° C., preferably from the first-stage cooling stop temperature to the cooling stop temperature between 400 to 300 ° C., less than 1 to 5 ° C./s. It is necessary to cool at a cooling rate of preferably 2 to 4 ° C./s.
前段冷却における停止温度が上記温度域よりも高い場合には、強度の増加がほとんどなく、逆に、上記温度域よりも低い場合には靭性が劣化する。また、前段冷却速度が上記範囲の下限未満では、ポリゴナルフェライトが主体の組織となって強度の向上が得られず、逆に上記範囲の上限を超えると靭性が低化する。さらに、後段冷却における冷却停止温度が上記温度域の上限よりも高い場合には、強度の上昇が不十分となる。また、後段冷却速度が上記範囲の下限未満では、母材強度が不足し、逆に上記範囲の上限を超えると、母材の靭性が低下するからである。 When the stop temperature in the pre-cooling is higher than the above temperature range, there is almost no increase in strength. Further, if the cooling rate at the front stage is less than the lower limit of the above range, the structure is mainly composed of polygonal ferrite, and the strength cannot be improved. Conversely, if the upper limit of the above range is exceeded, the toughness is lowered. Furthermore, when the cooling stop temperature in the latter stage cooling is higher than the upper limit of the temperature range, the increase in strength is insufficient. Moreover, if the latter stage cooling rate is less than the lower limit of the above range, the base material strength is insufficient, and conversely if the upper limit of the above range is exceeded, the toughness of the base material is lowered.
なお、本発明では、残留する内部応力を低減する目的で、上記冷却後の鋼材に、450〜650℃の温度で焼戻処理を施すことが好ましい。焼戻処理温度が450℃未満では、残留応力の除去効果が小さく、一方、650℃を超えて高くなると、各種炭窒化物が析出して析出強化を起こし、靭性が低下するからである。 In the present invention, it is preferable to temper the steel material after cooling at a temperature of 450 to 650 ° C. for the purpose of reducing the residual internal stress. This is because if the tempering temperature is less than 450 ° C., the effect of removing residual stress is small, while if it exceeds 650 ° C., various carbonitrides precipitate to cause precipitation strengthening and lower toughness.
以上説明したように、本発明の高張力鋼の製造方法においては、熱間圧延における圧延温度に応じた適正な圧下率制御と、圧延終了後の2段冷却条件の適正な制御が重要であり、とくに前段冷却の冷却速度を後段冷却のそれより大きくすることにより、母材がアシキュラーフェライト主体の組織となり、強度・靭性に優れた鋼材を得ることができる。 As described above, in the method for producing high-strength steel according to the present invention, it is important to appropriately control the rolling reduction according to the rolling temperature in hot rolling and to properly control the two-stage cooling conditions after the end of rolling. In particular, when the cooling rate of the pre-stage cooling is made larger than that of the post-stage cooling, the base material becomes a structure mainly composed of acicular ferrite, and a steel material excellent in strength and toughness can be obtained.
表1に示す各種成分組成に調整した鋼スラブを素材とし、表2−1および表2−2に示す製造条件で、厚さが55mmまたは65mmの厚鋼板を製造した。かくして得られた各厚鋼板からサンプルを採取し、引張試験およびシャルピー試験に供した。引張試験は、各厚鋼板の板厚中央部から、圧延幅方向にJIS4号引張試験片を採取し、降伏応力(YS)、引張強度(TS)を求めた。また、シャルピー衝撃試験は、各鋼板の板厚中央部から、圧延幅方向にJIS4号衝撃試験片を採取し、−40℃での吸収エネルギー(vE−40℃)を求めた。 Steel slabs adjusted to various component compositions shown in Table 1 were used as raw materials, and thick steel plates having a thickness of 55 mm or 65 mm were manufactured under the manufacturing conditions shown in Tables 2-1 and 2-2. A sample was taken from each thick steel plate thus obtained and subjected to a tensile test and a Charpy test. In the tensile test, a JIS No. 4 tensile test piece was taken in the rolling width direction from the center of the thickness of each thick steel plate, and yield stress (YS) and tensile strength (TS) were determined. In the Charpy impact test, a JIS No. 4 impact test piece was taken in the rolling width direction from the center of the thickness of each steel plate, and the absorbed energy at -40 ° C (vE-40 ° C) was determined.
さらに、各鋼板から採取した試験板にレ開先(開先角度30°)を加工し、入熱量が45kJ/cmのサブマージアーク溶接を行って溶接継手を作製し、この溶接継手から、レ開先のストレートボンド部にノッチを施したCTOD試験片を採取し、−10℃でCTOD試験を行った。なお、CTOD試験片の作製および試験条件は、英国規格BS7448に準拠して行った。
また、切欠位置をボンド部とするJIS4号衝撃試験片を採取し、試験温度−40℃でシャルピー衝撃試験を実施し、吸収エネルギー(vE−40℃)を求めた。
Furthermore, a groove (
Moreover, the JIS4 impact test piece which makes a notch position a bond part was extract | collected, the Charpy impact test was implemented at test temperature -40 degreeC, and the absorbed energy (vE-40 degreeC) was calculated | required.
上記の試験結果を表2−1および表2−2に併記して示した。この結果から、本発明例の鋼板は、母材の降伏応力(YS)が355MPa以上で、シャルピー吸収エネルギー(vE−40℃)が200J以上を有しており、母材の強度、靭性が共に優れていること、さらに、サブマージアーク溶接継手ボンド部についても、vE−40℃が200J以上で、CTOD値が0.50mm以上であり、溶接熱影響部の靭性にも優れていることがわかる。これに対して、本発明の範囲を外れる比較例では、上記いずれか1つ以上の特性が劣った鋼板しか得られていない。 The test results are shown together in Table 2-1 and Table 2-2. From this result, the steel sheet of the example of the present invention has a yield stress (YS) of the base material of 355 MPa or more and a Charpy absorbed energy (vE-40 ° C.) of 200 J or more, and the strength and toughness of the base material are both. It can be seen that the submerged arc welded joint bond portion is excellent, and vE-40 ° C. is 200 J or more and the CTOD value is 0.50 mm or more, and the weld heat affected zone has excellent toughness. On the other hand, in the comparative example outside the scope of the present invention, only the steel sheet having any one or more of the above-described characteristics is obtained.
Claims (5)
Ca,SおよびOが下記(1)式を満たして含有し、残部がFeおよび不可避的不純物からなる成分組成を有することを特徴とする高張力鋼。
記
0<(Ca−(0.18+130×Ca)×O)/1.25/S<1 ・・・(1)
ここで、Ca,SおよびOは、各元素の含有量(mass%) C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less, Al: 0.005 ~0.06mass%, Ni: 0.3 ~2mass% , Nb: 0.004~0.05mass%, Ti: 0.005~0.02mass%, N: less than 0.003mass%, Ca: 0.0005 Contains 0.003 mass%,
A high-strength steel characterized in that Ca, S and O satisfy the following formula (1) , and the balance has a composition composed of Fe and inevitable impurities .
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%).
Ca,SおよびOが下記(1)式を満たして含有し、残部がFeおよび不可避的不純物からなる成分組成を有する鋼スラブを1050〜1200℃に加熱後、950℃以上の温度域における累積圧下率が30%以上、950℃未満の温度域における累積圧下率が30〜70%となる熱間圧延を施し、その後、熱間圧延終了温度から600〜450℃間の冷却停止温度までを5〜20℃/sで冷却する前段冷却と、前段冷却停止温度から450℃未満〜200℃間の冷却停止温度までを1〜5℃未満/sで冷却する後段冷却を施すことを特徴とする高張力鋼の製造方法。
記
0<(Ca−(0.18+130×Ca)×O)/1.25/S<1 ・・・(1)
ここで、Ca,SおよびOは、各元素の含有量(mass%) C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less, Al: 0.005 ~0.06mass%, Ni: 0.3 ~2mass% , Nb: 0.004~0.05mass%, Ti: 0.005~0.02mass%, N: less than 0.003mass%, Ca: 0.0005 Contains 0.003 mass%,
Ca, S and O satisfy the following formula (1) and the steel slab having a composition composed of Fe and unavoidable impurities in the balance is heated to 1050 to 1200 ° C., and then cumulative pressure in a temperature range of 950 ° C. or higher. The hot rolling is performed so that the cumulative reduction in the temperature range of 30% or more and less than 950 ° C. is 30 to 70%, and thereafter, from the hot rolling end temperature to the cooling stop temperature of 600 to 450 ° C. High tension characterized by performing pre-stage cooling that cools at 20 ° C./s, and post-stage cooling that cools from the pre-stage cooling stop temperature to a cooling stop temperature between 450 ° C. and 200 ° C. at less than 1 to 5 ° C./s. Steel manufacturing method.
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%).
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