JP2013133485A - High-strength hot-rolled steel sheet excellent in stretch flange formability, and manufacturing method therefor - Google Patents
High-strength hot-rolled steel sheet excellent in stretch flange formability, and manufacturing method therefor Download PDFInfo
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Abstract
Description
本発明は、自動車用部材の使途に有用な、引張強さ(TS)：850MPa以上1150MPa以下の高強度と優れた加工性（特に伸びフランジ性）を兼ね備えた高強度熱延鋼板およびその製造方法に関する。 The present invention is a high-strength hot-rolled steel sheet that has high tensile strength (TS): not less than 850 MPa and not more than 1150 MPa and excellent workability (particularly stretch flangeability), and a method for producing the same, useful for use in automobile parts About.
近年地球環境保全の観点から、CO2排出量の規制を目的として自動車業界全体で自動車の燃費改善が指向されている。自動車の燃費改善には、使用部材の薄肉化による自動車の軽量化が最も有効であるため、自動車部材用素材としての高強度熱延鋼板の使用量が増加しつつある。一方、鋼板を素材とする自動車部材の多くは、プレス加工やバーリング加工等によって成形されるため、自動車部材用鋼板には高強度に加えて優れた伸びフランジ性等の加工性を有することも要求される。 In recent years, from the viewpoint of global environmental conservation, the automobile industry as a whole has been aimed at improving the fuel efficiency of automobiles for the purpose of regulating CO 2 emissions. In order to improve the fuel efficiency of automobiles, it is most effective to reduce the weight of automobiles by reducing the thickness of the parts used. Therefore, the amount of high-strength hot-rolled steel sheets used as automobile material is increasing. On the other hand, since many automotive parts made of steel plates are formed by pressing or burring, etc., steel sheets for automotive parts are required to have workability such as excellent stretch flangeability in addition to high strength. Is done.
しかしながら、一般的に鉄鋼材料は高強度化に伴い延性が低下して加工性が劣化し、引張強さを850MPa以上にまで高強度化した鋼板の加工性は通常の軟鋼板よりもはるかに劣っている。そのため、高強度熱延鋼板を自動車部材等に適用するうえでは、伸びフランジ性などの加工性を兼ね備えた高強度熱延鋼板の開発が必須となり、現在までに様々な技術が提案されている。 However, in general, steel materials have a lower ductility and higher workability with higher strength, and the workability of steel sheets with higher tensile strength up to 850 MPa or higher is far inferior to that of ordinary mild steel sheets. ing. Therefore, in order to apply a high-strength hot-rolled steel sheet to automobile members and the like, it is essential to develop a high-strength hot-rolled steel sheet having workability such as stretch flangeability, and various techniques have been proposed so far.
例えば、特許文献１では、鋼板組成を質量％で、C：0.02〜0.08％、Si：0.01〜1.50％、Mn：0.1〜1.5％、Ti：0.03〜0.06％を含有し、P：0.1％以下、S：0.005％以下、Al：0.5％以下、N：0.009％以下に制限し、更に、Nb、Mo、Vの含有量の合計を0.01％以下に制限し、残部がFeおよび不可避的不純物からなり、C量に対するTi量の比がTi/C：0.375〜1.6である組成とし、結晶粒内のTiC析出物の平均直径を0.8〜3nm、平均個数密度を1×1017[個/cm3]以上とする技術が提案されている。そして、特許文献１で提案された技術によると、炭化物形成能が最も高いTiを効率的に析出強化に利用することで、合金元素の添加に起因する加工性低下が抑制された引張強度：540〜650MPaの省合金型高強度熱延鋼板が得られるとされている。 For example, in Patent Document 1, the steel plate composition is mass%, C: 0.02 to 0.08%, Si: 0.01 to 1.50%, Mn: 0.1 to 1.5%, Ti: 0.03 to 0.06%, and P: 0.1% or less. , S: 0.005% or less, Al: 0.5% or less, N: 0.009% or less, further limit the total content of Nb, Mo, V to 0.01% or less, the remainder from Fe and inevitable impurities The ratio of Ti amount to C amount is Ti / C: 0.375 to 1.6, the average diameter of TiC precipitates in the crystal grains is 0.8 to 3 nm, and the average number density is 1 × 10 17 [pieces / cm 3 The above technique has been proposed. And according to the technique proposed in Patent Document 1, the tensile strength: 540, in which Ti, which has the highest carbide forming ability, is efficiently used for precipitation strengthening to suppress the workability degradation due to the addition of alloy elements. It is said that an alloy-saving high-strength hot-rolled steel sheet of ~ 650 MPa can be obtained.
また、特許文献２では、鋼板組成を質量％で、C：0.015〜0.06％、Si：0.5％未満、Mn：0.1〜2.5％、P≦0.10％、S≦0.01％、Al：0.005〜0.3％、N≦0.01％、Ti：0.01〜0.30％、B：2〜50ppmを含有し、残部Fe及び不可避的不純物からなる組成とし、更に炭化物生成元素とCとの原子比を特定するとともに、鋼のγ／α変態温度を制御する元素であるSi、Mn、B、Moの含有量が所定の関係を満足するように規定し、フェライトとベイニティックフェライトの一方又は双方の面積率の合計が90％以上でありセメンタイトの面積率が5％以下である鋼板組織とする技術が提案されている。そして、特許文献２で提案された技術によると、炭化物析出により低下した粒界強度をB添加により向上することで打ち抜き端面の欠陥を抑え、伸びフランジ性に優れるとともに耐打ち抜き割れ性及び表面状態が良好であり、引張強度が690MPa以上という高強度の熱延鋼板を安価に、安定して製造することができるとされている。 Moreover, in patent document 2, a steel plate composition is the mass%, C: 0.015-0.06%, Si: Less than 0.5%, Mn: 0.1-2.5%, P <= 0.10%, S <= 0.01%, Al: 0.005-0.3% , N ≦ 0.01%, Ti: 0.01 to 0.30%, B: 2 to 50 ppm, the balance consisting of Fe and unavoidable impurities, and further specifying the atomic ratio of carbide-forming elements and C, The content of Si, Mn, B, and Mo, which are elements that control the γ / α transformation temperature, is specified so as to satisfy a predetermined relationship, and the total area ratio of one or both of ferrite and bainitic ferrite is 90 %, And a cementite structure with a cementite area ratio of 5% or less has been proposed. According to the technique proposed in Patent Document 2, the grain boundary strength reduced by carbide precipitation is improved by adding B to suppress defects in the punched end face, and the stretch flangeability is excellent and the punching cracking resistance and the surface condition are improved. It is said that a high-strength hot-rolled steel sheet having a good tensile strength of 690 MPa or more can be stably manufactured at low cost.
また、特許文献３では、鋼板の本質的成分をmass％で、C：0.01〜0.10％、Si：1.0％以下、Mn：2.5％以下、P：0.08％以下、S：0.005％以下、Al：0.015〜0.050％、Ti：0.10％〜0.30％および残部Feである成分とし、鋼板組織をフェライト主体とし、隣接する粒との方位差がすべて15°以上で囲まれた粒を単位粒とし、その平均粒径をｄμm としたとき、ｄが5μm 以下である組織とする技術が提案されている。そして、特許文献３で提案された技術によると、フェライト粒径とその形態をコントロールすることで、伸びフランジ性に優れた高強度熱延鋼板が得られるとされている。 Moreover, in patent document 3, the essential component of a steel plate is mass%, C: 0.01-0.10%, Si: 1.0% or less, Mn: 2.5% or less, P: 0.08% or less, S: 0.005% or less, Al: The component is 0.015 to 0.050%, Ti: 0.10% to 0.30% and the balance Fe, the steel sheet structure is mainly ferrite, and the grains whose orientation difference from adjacent grains are all surrounded by 15 ° or more are unit grains. There has been proposed a technique for forming a structure in which d is 5 μm or less when the average particle diameter is d μm. And according to the technique proposed by patent document 3, it is supposed that the high intensity | strength hot-rolled steel plate excellent in stretch flangeability will be obtained by controlling a ferrite particle size and its form.
しかしながら、特許文献１で提案された技術では、炭化物（TiC）による析出強化に利用するTiの含有量が0.03〜0.06％と少ないため、炭化物（TiC）を大量に析出させることができず、引張強さが650MPa程度の鋼板しか得られない。また、炭化物（TiC）析出量の増加を図るためにはTi含有量を多くすることが必要となる一方、Ti含有量が増加するにつれてTiCは粗大化する傾向にある。このような問題に対し、特許文献１で提案された技術では、炭化物（TiC）の粗大化を抑制する手段が十分に検討されていない。すなわち、特許文献１で提案された技術では、Ti含有量が0.06％を超えると炭化物（TiC）が粗大化し易くなるため、鋼板の引張強さを850MPa以上にすることは極めて困難である。 However, in the technique proposed in Patent Document 1, since the Ti content used for precipitation strengthening by carbide (TiC) is as low as 0.03 to 0.06%, a large amount of carbide (TiC) cannot be precipitated, and tensile Only steel sheets with a strength of about 650 MPa can be obtained. Further, in order to increase the amount of carbide (TiC) precipitation, it is necessary to increase the Ti content. On the other hand, TiC tends to become coarser as the Ti content increases. With respect to such a problem, in the technique proposed in Patent Document 1, means for suppressing the coarsening of carbide (TiC) has not been sufficiently studied. That is, with the technique proposed in Patent Document 1, if the Ti content exceeds 0.06%, carbide (TiC) tends to be coarsened, so it is extremely difficult to increase the tensile strength of the steel sheet to 850 MPa or more.
また、特許文献２で提案された技術では、Ti含有量が比較的少ない、或いはC、Ti含有量のバランスに問題があるため、引張強さ850MPa以上の鋼板を得ることはできていない。更に、特許文献２で提案された技術では、鋼板にBを含有させるとともに、固溶強化元素であり且つ析出強化に寄与する炭化物の析出を制御する元素でもあるMnを含有させることで、鋼板強度の向上を図っている。しかしながら、この技術では、その実施例が示すように、Mn含有量を1.0％としても引張強さが850MPa未満の鋼板しか得られていない。そして、鋼板強度の更なる向上を図る目的でMn含有量を1.0％よりも多くすると、鋼板にMnの中心偏析が生じ易くなる結果、鋼板の伸びフランジ性が低下してしまう。なお、特許文献２には、Mn含有量を0.5％とした鋼板の実施例も開示されているが、この鋼板の引張強さは750MPa未満と低く、しかも析出強化元素としてTiのほかに高価なNbも利用しているため、Tiを最大限活用することはできず、製品コストの高騰は避けられない。 Further, the technique proposed in Patent Document 2 cannot obtain a steel sheet having a tensile strength of 850 MPa or more because the Ti content is relatively small or there is a problem in the balance between C and Ti contents. Furthermore, in the technique proposed in Patent Document 2, the steel sheet strength is increased by including B in the steel sheet and Mn, which is a solid solution strengthening element and also an element that controls precipitation of carbides contributing to precipitation strengthening. We are trying to improve. However, in this technique, as shown in the Examples, only steel sheets having a tensile strength of less than 850 MPa can be obtained even when the Mn content is 1.0%. If the Mn content is more than 1.0% for the purpose of further improving the strength of the steel sheet, Mn center segregation is likely to occur in the steel sheet, resulting in a decrease in stretch flangeability of the steel sheet. Patent Document 2 discloses an example of a steel sheet having an Mn content of 0.5%. However, the tensile strength of this steel sheet is as low as less than 750 MPa, and it is expensive in addition to Ti as a precipitation strengthening element. Since Nb is also used, Ti cannot be utilized to the maximum, and the product cost is inevitably increased.
また、特許文献３で提案された技術では、固溶強化元素であり且つ変態を促進するとともに粒界形状に影響を及ぼすMnを含有させることで、鋼板強度の向上を図っている。しかしながら、特許文献３で提案された技術においても、その実施例が示すように、Mn含有量を1.5％としても引張強さが850MPa未満の鋼板しか得られず、しかもMnを1.0％を超えて多量に含有することから伸びフランジ性の低下も問題となる。なお、特許文献３には、Mn含有量を0.3％とした鋼板の実施例も開示されているが、この鋼板の引張強さは730MPaと低く、しかも析出強化元素としてTiのほかに高価なNbも利用しているため、特許文献２で提案された技術と同様、製品コストの高騰を招く。 In the technique proposed in Patent Document 3, the strength of the steel sheet is improved by containing Mn, which is a solid solution strengthening element and promotes transformation and affects the shape of the grain boundary. However, in the technique proposed in Patent Document 3, as shown in the examples, even when the Mn content is 1.5%, only a steel sheet having a tensile strength of less than 850 MPa can be obtained, and the Mn exceeds 1.0%. Since it is contained in a large amount, a decrease in stretch flangeability is also a problem. In addition, Patent Document 3 discloses an example of a steel plate with an Mn content of 0.3%. However, the tensile strength of this steel plate is as low as 730 MPa, and it is expensive Nb as a precipitation strengthening element in addition to Ti. Therefore, as in the technique proposed in Patent Document 2, the product cost increases.
以上のように、従来技術ではいずれも、引張強さ850MPa以上の鋼板強度を確保しようとする場合、固溶強化元素であるMnを多量に含有させることが必要となるため、Mnの中心偏析に起因する伸びフランジ性の低下を回避することができない。すなわち、従来技術では、引張強さが850MPa以上であり且つ伸びフランジ性にも優れた高強度熱延鋼板を得ることは極めて困難である。
本発明はかかる事情に鑑みてなされたものであって、850MPa以上の引張強さを有し、加工性、特に伸びフランジ性にも優れた高強度熱延鋼板を提供することを目的とする。
As described above, in all of the conventional techniques, it is necessary to contain a large amount of the solid solution strengthening element Mn when securing a steel sheet strength of 850 MPa or more in tensile strength. It is impossible to avoid the deterioration of stretch flangeability due to the above. That is, it is extremely difficult to obtain a high-strength hot-rolled steel sheet having a tensile strength of 850 MPa or more and excellent stretch flangeability by the conventional technology.
The present invention has been made in view of such circumstances, and an object thereof is to provide a high-strength hot-rolled steel sheet having a tensile strength of 850 MPa or more and excellent workability, particularly stretch flangeability.
上記課題を解決すべく、本発明者らは、加工性が良好なフェライト単相組織である熱延鋼板に着目し、該熱延鋼板の高強度化と加工性、特に伸びフランジ性に及ぼす各種要因について鋭意検討した。その結果、従来、固溶強化元素として鋼板の高強度化に極めて有効であるとされ、高強度熱延鋼板に積極的に含有させていたMnが、鋼板の加工性、特に伸びフランジ性に悪影響を及ぼすことを知見した。また、鋼板に含まれるMnの中心偏析が伸びフランジ性の劣化の要因となっていること、更にこれらの偏析を抑制するうえではMn含有量を低減する必要があることを知見した。 In order to solve the above-mentioned problems, the present inventors have focused on hot-rolled steel sheets that are ferrite single-phase structures with good workability, and various effects that affect the strength and workability of the hot-rolled steel sheets, particularly stretch flangeability. The factors were studied earnestly. As a result, Mn, which has been considered to be extremely effective as a solid solution strengthening element for increasing the strength of steel sheets, has been adversely affected by the workability of steel sheets, especially stretch flangeability, which was positively incorporated in high-strength hot-rolled steel sheets. I found out that It was also found that the center segregation of Mn contained in the steel sheet is a cause of deterioration of stretch flangeability, and that it is necessary to reduce the Mn content in order to suppress these segregations.
一方、固溶強化元素であるMn含有量の抑制に伴う鋼板強度の低下は避けられない。そこで、本発明者らは、Mnによる固溶強化に代わる強化機構として、炭化物による析出強化を採用し、該炭化物を鋼板のマトリックス（主相）であるフェライト相に微細析出させることで、所望の鋼板強度（引張強さ：850MPa以上）とすることを試みた。また、フェライト相に析出させる炭化物が微細かつ析出量が多いほど鋼板強度の大幅な向上効果が期待できることから、炭化物の析出量を十分に確保する手段、および該炭化物の微細化を図る手段について模索した。 On the other hand, a decrease in steel sheet strength accompanying the suppression of the Mn content, which is a solid solution strengthening element, is inevitable. Therefore, the present inventors adopt precipitation strengthening by carbide as a strengthening mechanism instead of solid solution strengthening by Mn, and finely precipitate the carbide in a ferrite phase which is a matrix (main phase) of a steel plate, thereby obtaining a desired structure. Attempts were made to obtain steel plate strength (tensile strength: 850 MPa or more). In addition, the finer the carbide precipitated in the ferrite phase and the greater the amount of precipitation, the greater the effect of improving the strength of the steel sheet can be expected. Therefore, we sought for means to ensure a sufficient amount of carbide precipitation and to refine the carbide. did.
その結果、炭化物形成能の高いTiを活用すること、すなわち鋼板のマトリックスであるフェライト相に析出させる炭化物としてTi炭化物を用いることが、炭化物析出量を確保する有効な手段であることに想到した。
また、Ti炭化物は、鋼板製造時、熱間圧延終了後の冷却過程で析出するが、高温域で析出する炭化物は粗大化し易い一方、炭化物を鋼板巻取り温度域で析出させると微細な炭化物が得られることが明らかになった。そして、Ti炭化物は、熱間圧延終了後の冷却過程において鋼のオーステナイト→フェライト変態とほぼ同時に析出することから、鋼のオーステナイト→フェライト変態点を低温化し、該変態点を鋼板巻取り温度域に調整することで、微細な炭化物が得られることを知見した。
As a result, it was conceived that utilizing Ti with high carbide forming ability, that is, using Ti carbide as a carbide to be precipitated in the ferrite phase that is the matrix of the steel sheet is an effective means for securing the amount of carbide precipitation.
In addition, Ti carbide precipitates during the cooling process after hot rolling at the time of steel plate production, while carbide precipitated in the high temperature range is easy to coarsen, while fine carbide is formed when carbide is precipitated in the steel sheet winding temperature range. It became clear that it was obtained. And since Ti carbide precipitates almost simultaneously with the austenite → ferrite transformation of steel in the cooling process after the end of hot rolling, the temperature of the austenite → ferrite transformation of the steel is lowered, and the transformation point is brought into the steel sheet winding temperature range. It has been found that fine carbides can be obtained by adjusting.
ここで、Mnは、鋼のオーステナイト→フェライト変態点を低下させる効果を有する元素であるが、先述のとおり鋼板に多量のMnを含有させると伸びフランジ性が低下してしまう。そこで、本発明者らは、熱延鋼板の伸びフランジ性と高強度の両立を図る手段について更に検討を進めた結果、鋼板のMn含有量を1.00％以下に低減するとともに、C、Ti含有量を調整することで、伸びフランジ性の低下を抑制しつつTi炭化物を微細且つ大量に析出させることが可能となり、Tiの炭化物による析出強化を最大限活用できることを知見した。 Here, Mn is an element having an effect of lowering the austenite → ferrite transformation point of steel, but if a large amount of Mn is contained in the steel plate as described above, stretch flangeability is lowered. Therefore, as a result of further study on means for achieving both stretch flangeability and high strength of the hot-rolled steel sheet, the present inventors reduced the Mn content of the steel sheet to 1.00% or less, and the C and Ti contents. It was found that by adjusting the amount of Ti, it was possible to precipitate a large amount of Ti carbide while suppressing a decrease in stretch flangeability, and the precipitation strengthening by Ti carbide could be utilized to the maximum extent.
また、本発明者らは、熱延鋼板に所定量のBを含有させ、更に熱延鋼板の製造条件を規定することで、熱延鋼板の諸特性を阻害することなく鋼のオーステナイト→フェライト変態点の低温化が可能であることを突き止めた。そして、鋼板に所定量のBを含有させることで、Mn含有量を更に低減した場合であってもTi炭化物の微細化が可能となり、BとMnの含有量を所定の関係を満足するように規定すると一層優れた伸びフランジ性を示す高強度熱延鋼板が得られることを知見した。更に、上記に加えてフェライト相の平均結晶粒径を10μm以下にすると、強度と伸びフランジ性がより一層向上することを知見した。 Further, the present inventors include a predetermined amount of B in the hot-rolled steel sheet, and further specify the production conditions of the hot-rolled steel sheet, thereby preventing the austenite → ferrite transformation of the steel without impairing various properties of the hot-rolled steel sheet. I found out that the temperature of the spot could be lowered. And by including a predetermined amount of B in the steel plate, even if the Mn content is further reduced, it is possible to refine the Ti carbide, so that the B and Mn content satisfy a predetermined relationship. It has been found that a high-strength hot-rolled steel sheet exhibiting even better stretch flangeability can be obtained when it is specified. Furthermore, in addition to the above, it has been found that when the average grain size of the ferrite phase is 10 μm or less, the strength and stretch flangeability are further improved.
本発明は上記の知見に基づき完成されたものであり、その要旨は次のとおりである。
［１］ 質量％で、
C ：0.05％以上0.09％以下、 Si：0.3％以下、
Mn：1.00％以下、 P ：0.03％以下、
S ：0.007％以下、 Al：0.1％以下、
N ：0.01％以下、 Ti：0.18％以上0.25％以下
を、C、NおよびTiが下記(1)を満足するように含有し、残部がFeおよび不可避的不純物からなる組成と、フェライト相の面積率が95％以上、該フェライト相の平均結晶粒径が10μm以下であり、前記フェライト相の結晶粒内の炭化物平均粒子径が10nm未満である組織を有し、引張強さが850MPa以上1150MPa以下であることを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。
記
1.0 ≦ ([C]/12)/([Ti*]/48) ≦ 1.5 ・・・ (1)
但し、[Ti*]＝[Ti]−3.4×[N]
（[C]、[N]、[Ti]：各元素の含有量（質量％））
The present invention has been completed based on the above findings, and the gist thereof is as follows.
[1] By mass%
C: 0.05% or more and 0.09% or less, Si: 0.3% or less,
Mn: 1.00% or less, P: 0.03% or less,
S: 0.007% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.18% or more and 0.25% or less, so that C, N and Ti satisfy the following (1), with the balance consisting of Fe and inevitable impurities, and the area of the ferrite phase The ratio is 95% or more, the ferrite phase has an average crystal grain size of 10 μm or less, and has a structure in which the average grain size of carbide in the ferrite phase grains is less than 10 nm, and the tensile strength is 850 MPa or more and 1150 MPa or less. A high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized by being
Record
1.0 ≤ ([C] / 12) / ([Ti * ] / 48) ≤ 1.5 ... (1)
However, [Ti * ] = [Ti] −3.4 × [N]
([C], [N], [Ti]: Content of each element (mass%))
［２］ 前記［１］において、前記組成に加えてさらに、質量％でB ：0.003％以下を、下記(2)式を満足するように含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。
記
2.21×[Mn]＋1.05×log(104×([B]＋0.0001))≧1.44 ・・・ (2)
（[Mn]、[B]：各元素の含有量（質量％））
[2] In the above [1], in addition to the above composition, B: 0.003% or less by mass% is contained so as to satisfy the following formula (2), and has excellent stretch flangeability High strength hot rolled steel sheet.
Record
2.21 x [Mn] + 1.05 x log (10 4 x ([B] + 0.0001)) ≥ 1.44 (2)
([Mn], [B]: Content of each element (mass%))
［３］ 前記［１］または［２］において、前記組成に加えてさらに、質量％でV：0.005％以上0.3％以下を含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。 [3] In the above [1] or [2], in addition to the above composition, V: 0.005% or more and 0.3% or less in mass% is further included, and high strength hot rolling excellent in stretch flangeability steel sheet.
［４］ 前記［１］ないし［３］のいずれかにおいて、前記組成に加えてさらに、質量％でW ：0.01％以上1.0％以下、Mo：0.01％以上0.5％以下のいずれか１種以上を含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。 [4] In any one of the above [1] to [3], in addition to the above composition, at least one of W: 0.01% to 1.0% and Mo: 0.01% to 0.5% by mass% is added. A high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized by containing.
［５］ 前記［１］ないし［４］のいずれかにおいて、前記組成に加えてさらに、質量％で、Se、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、Ni、Cr、Sb、Cu、Sn、Mg、Caのうちの１種以上を合計で1.0％以下含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。 [5] In any one of the above [1] to [4], in addition to the above composition, Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd in addition to the composition. Hg, Ag, Au, Pd, Pt, Co, Rh, Ir, Ru, Os, Tc, Re, Ta, Be, Sr, REM, Ni, Cr, Sb, Cu, Sn, Mg, Ca A high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized by containing a total of 1.0% or less of seeds or more.
［６］ 前記［１］ないし［５］のいずれかにおいて、鋼板表面にめっき層を有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。 [6] The high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized by having a plated layer on the steel sheet surface in any one of [1] to [5].
［７］ 前記［６］において、前記めっき層が亜鉛めっき層であることを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。 [7] The high-strength hot-rolled steel sheet excellent in stretch flangeability, wherein the plating layer is a galvanization layer in [6].
［８］ 前記［６］において、前記めっき層が合金化亜鉛めっき層であることを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。 [8] A high-strength hot-rolled steel sheet having excellent stretch flangeability, wherein the plating layer is an alloyed galvanized layer in [6].
［９］ 鋼素材を加熱し、粗圧延と仕上げ圧延からなる熱間圧延を施し、仕上げ圧延終了後、冷却し、巻き取り、熱延鋼板とするにあたり、
前記鋼素材を、質量％で、
C ：0.05％以上0.09％以下、 Si：0.3％以下、
Mn：1.00％以下、 P ：0.03％以下、
S ：0.007％以下、 Al：0.1％以下、
N ：0.01％以下、 Ti：0.18％以上0.25％以下
を、C、NおよびTiが下記(1)を満足するように含有し、残部がFeおよび不可避的不純物からなる組成とし、
前記加熱の加熱温度を1150℃以上1350℃以下とし、前記仕上げ圧延の仕上げ圧延温度を850℃以上とし、前記冷却を仕上げ圧延終了後3秒以内に開始し、前記冷却の平均冷却速度を30℃/s以上とし、前記巻き取りの巻取り温度を500℃以上700℃以下とすることを特徴とする、伸びフランジ性に優れた高強度熱延鋼板の製造方法。
記
1.0 ≦ ([C]/12)/([Ti*]/48) ≦ 1.5 ・・・ (1)
但し、[Ti*]＝[Ti]−3.4×[N]
（[C]、[N]、[Ti]：各元素の含有量（質量％））
[9] Heating the steel material, subjecting it to hot rolling consisting of rough rolling and finish rolling, cooling after completion of finish rolling, winding, and hot rolling steel sheet,
The steel material in mass%,
C: 0.05% or more and 0.09% or less, Si: 0.3% or less,
Mn: 1.00% or less, P: 0.03% or less,
S: 0.007% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.18% or more and 0.25% or less, so that C, N and Ti satisfy the following (1), the balance is composed of Fe and inevitable impurities,
The heating temperature of the heating is 1150 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the average cooling rate of the cooling is 30 ° C. The method for producing a high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized in that the winding temperature is 500 ° C. or higher and 700 ° C. or lower.
Record
1.0 ≤ ([C] / 12) / ([Ti * ] / 48) ≤ 1.5 ... (1)
However, [Ti * ] = [Ti] −3.4 × [N]
([C], [N], [Ti]: Content of each element (mass%))
［１０］ 前記［９］において、前記組成に加えてさらに、質量％でB ：0.003％以下を、下記(2)式を満足するように含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板の製造方法。
記
2.21×[Mn]＋1.05×log(104×([B]＋0.0001))≧1.44 ・・・ (2)
（[Mn]、[B]：各元素の含有量（質量％））
[10] In the above [9], in addition to the above composition, B: 0.003% or less by mass% is contained so as to satisfy the following formula (2), and has excellent stretch flangeability Manufacturing method of high-strength hot-rolled steel sheet.
Record
2.21 x [Mn] + 1.05 x log (10 4 x ([B] + 0.0001)) ≥ 1.44 (2)
([Mn], [B]: Content of each element (mass%))
［１１］ 前記［９］または［１０］において、前記組成に加えてさらに、質量％でV：0.005％以上0.3％以下を含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板の製造方法。 [11] In the above [9] or [10], in addition to the above composition, V: 0.005% or more and 0.3% or less in mass% is further included. A method of manufacturing a steel sheet.
［１２］ 前記［９］ないし［１１］のいずれかにおいて、前記組成に加えてさらに、質量％でW ：0.01％以上1.0％以下、Mo：0.01％以上0.5％以下のいずれか１種以上を含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板の製造方法。 [12] In any one of the above [9] to [11], in addition to the composition, W 1 is 0.01% or more and 1.0% or less, and Mo is 0.01% or more and 0.5% or less. A method for producing a high-strength hot-rolled steel sheet excellent in stretch flangeability, comprising:
［１３］ 前記［９］ないし［１２］のいずれかにおいて、前記組成に加えてさらに、質量％でSe、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、Ni、Cr、Sb、Cu、Sn、Mg、Caのうちの１種以上を合計で1.0％以下含有することを特徴とする、伸びフランジ性に優れた高強度熱延鋼板の製造方法。 [13] In any one of [9] to [12], in addition to the composition, Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd, One of Hg, Ag, Au, Pd, Pt, Co, Rh, Ir, Ru, Os, Tc, Re, Ta, Be, Sr, REM, Ni, Cr, Sb, Cu, Sn, Mg, Ca A method for producing a high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized by containing 1.0% or less in total.
本発明によると、自動車の構造部材等の使途に好適な、引張強さ：850MPa以上であり且つ伸びフランジ性に優れた高強度鋼板が得られ、自動車部材の軽量化や自動車部材成形を可能とする等、その効果は著しい。また、本発明によると、伸びフランジ性等の加工性を兼ね備えた引張強さ：850MPa以上の高強度熱延鋼板が得られることから、高強度熱延鋼板の更なる用途展開が可能となり、産業上格段の効果を奏する。 According to the present invention, it is possible to obtain a high-strength steel sheet having a tensile strength of 850 MPa or more and excellent stretch flangeability, which is suitable for the use of automobile structural members, etc., and can reduce the weight of automobile parts and form automobile parts. The effect is remarkable. In addition, according to the present invention, since a high strength hot rolled steel sheet having a tensile strength of 850 MPa or more that has workability such as stretch flangeability can be obtained, further application development of the high strength hot rolled steel sheet becomes possible. Has an exceptional effect.
以下、本発明について詳細に説明する。
まず、本発明鋼板の組織および炭化物の限定理由について説明する。
本発明の熱延鋼板は、フェライト相の面積率が95％以上、該フェライト相の平均結晶粒径が10μm以下であり、前記フェライト相の結晶粒内の炭化物平均粒子径が10nm未満である組織を有する。
Hereinafter, the present invention will be described in detail.
First, the structure of the steel sheet of the present invention and the reasons for limiting the carbide will be described.
The hot-rolled steel sheet of the present invention has a structure in which the ferrite phase area ratio is 95% or more, the average crystal grain size of the ferrite phase is 10 μm or less, and the average carbide grain size in the ferrite phase crystal grains is less than 10 nm. Have
フェライト相の面積率：95％以上
熱延鋼板の金属組織は、加工性に優れたフェライト単相組織とすることが好ましい。先述のとおり、本発明では熱延鋼板のマトリックスであるフェライトの結晶粒内に微細な炭化物を析出させることで所望の鋼板強度を確保する。そのため、本発明では、鋼素材に添加したCを微細な炭化物として析出させる必要があるところ、セメンタイトのような粗大な炭化物が存在すると微細な炭化物を形成するC量が減じ、鋼板強度が低下する。更に、マトリックス（フェライト）とセメンタイトのような粗大な炭化物との界面ではミクロボイドが発生し易くなるため、伸びフランジ性が低下する。
Area ratio of ferrite phase: 95% or more The metal structure of the hot-rolled steel sheet is preferably a ferrite single-phase structure excellent in workability. As described above, in the present invention, a desired strength of the steel sheet is ensured by precipitating fine carbides in the ferrite crystal grains which are the matrix of the hot-rolled steel sheet. Therefore, in the present invention, it is necessary to precipitate C added to the steel material as fine carbides. If coarse carbides such as cementite are present, the amount of C forming fine carbides is reduced, and the steel sheet strength is reduced. . Furthermore, since microvoids are likely to be generated at the interface between the matrix (ferrite) and a coarse carbide such as cementite, stretch flangeability deteriorates.
以上の理由により、フェライト相の面積率が95％を下回ると、ミクロボイド発生による悪影響が顕在化して伸びフランジ性が劣化する、或いはフェライト結晶粒内の微細な炭化物の析出量が不足して所望の鋼板強度（引張強さ：850MPa）が得られない。したがって、本発明の熱延鋼板組織は実質的にフェライト単相、すなわちフェライト相の面積率を95％以上とする必要がある。好ましくは98％以上である。 For the above reasons, when the area ratio of the ferrite phase is less than 95%, the adverse effect due to the generation of microvoids becomes obvious and the stretch flangeability deteriorates, or the precipitation amount of fine carbides in the ferrite crystal grains is insufficient and desired. Steel sheet strength (tensile strength: 850 MPa) cannot be obtained. Therefore, it is necessary that the hot rolled steel sheet structure of the present invention has a ferrite single phase, that is, a ferrite phase area ratio of 95% or more. Preferably it is 98% or more.
なお、本発明の熱延鋼板において、鋼板中に含有され得るフェライト相以外の組織としては、セメンタイト、パーライト、ベイナイト相、マルテンサイト相等が挙げられる。これらの組織が、多量に鋼板中に存在すると鋼板特性（伸びフランジ性等）が低下する。そのため、これらの組織は極力低減することが好ましいが、鋼板の金属組織全体に対する合計面積率が5％以下であれば許容される。好ましくは2％以下である。 In the hot-rolled steel sheet of the present invention, examples of the structure other than the ferrite phase that can be contained in the steel sheet include cementite, pearlite, bainite phase, and martensite phase. If these structures are present in a large amount in the steel sheet, the steel sheet properties (e.g. stretch flangeability) are deteriorated. Therefore, it is preferable to reduce these structures as much as possible, but it is acceptable if the total area ratio with respect to the entire metal structure of the steel sheet is 5% or less. Preferably it is 2% or less.
フェライト平均結晶粒径：10μm以下
フェライト平均結晶粒径が10μmを上回ると、結晶粒微細化強化による強化量が低下し、引張強さ850MPa以上の高強度熱延鋼板が得られなくなる。また、フェライト平均結晶粒径が10μmを上回ると、フェライト粒径が一定ではない混粒組織となるため、伸びフランジ成形時の変形挙動が鋼板内部で不均一となる。そして、この不均一な変形挙動はミクロボイド生成の原因となり、結果として鋼板の伸びフランジ性を低下させることとなる。そのため、フェライト平均結晶粒径を10μm以下とする。好ましくは6μm未満である。
Average ferrite grain size: 10 μm or less When the average ferrite grain size exceeds 10 μm, the amount of strengthening due to the strengthening of grain refinement decreases, and a high-strength hot-rolled steel sheet with a tensile strength of 850 MPa or more cannot be obtained. On the other hand, when the average ferrite grain size exceeds 10 μm, a mixed grain structure in which the ferrite grain size is not constant is formed, so that the deformation behavior at the time of stretch flange forming becomes non-uniform inside the steel sheet. And this non-uniform | heterogenous deformation | transformation behavior becomes a cause of micro void generation, As a result, the stretch flangeability of a steel plate will be reduced. Therefore, the ferrite average crystal grain size is set to 10 μm or less. Preferably it is less than 6 μm.
フェライト結晶粒内の炭化物
上記のとおり、本発明の熱延鋼板では、伸びフランジ性に悪影響を及ぼす板厚中央部のMn偏析を抑制する目的で固溶強化元素であるMn含有量を低減するため、固溶強化による鋼板強度の向上化は期待できない。そこで、本発明の熱延鋼板では、強度を確保する上でフェライト相の結晶粒内に炭化物を微細析出させることが必須となる。本発明においてフェライト相の結晶粒内に炭化物を微細析出させる炭化物としては、Ti炭化物、或いは更にV炭化物、TiとW、Mo、Vを含む複合炭化物等が挙げられる。なお、これらの炭化物の多くは、熱延鋼板製造工程における仕上げ圧延終了後の冷却過程で、オーステナイト→フェライト変態と同時に相界面析出する炭化物である。
Carbides in ferrite grains As described above, in the hot-rolled steel sheet of the present invention, in order to reduce the Mn content, which is a solid solution strengthening element, for the purpose of suppressing Mn segregation at the center of the sheet thickness that adversely affects stretch flangeability. The improvement of steel sheet strength by solid solution strengthening cannot be expected. Therefore, in the hot-rolled steel sheet of the present invention, it is essential to finely precipitate carbide in the ferrite phase crystal grains in order to ensure strength. In the present invention, examples of the carbide that finely precipitates carbide in the ferrite phase grains include Ti carbide, or further V carbide, composite carbide containing Ti and W, Mo, and V. Most of these carbides are carbides that precipitate at the interface at the same time as the austenite → ferrite transformation in the cooling process after finishing rolling in the hot rolled steel sheet manufacturing process.
フェライト結晶粒内の炭化物平均粒子径：10nm未満
本発明鋼では前記したTi等の炭化物を微細に分散させることで強化を図っている。炭化物が粗大化すると、鋼板に変形が加わった際に生じる転位の運動を阻害する炭化物数が減じることから、炭化物が微細化するほど鋼板は高強度化する。引張強さ850MPa以上の高強度熱延鋼板を得るには、上記炭化物の平均粒子径を10nm未満とする必要がある。好ましくは6nm以下である。
Carbide average particle diameter in ferrite crystal grains: less than 10 nm In the steel of the present invention, strengthening is achieved by finely dispersing carbides such as Ti described above. When the carbides are coarsened, the number of carbides that hinder the movement of dislocations that are generated when deformation is applied to the steel sheet decreases, so that the strength of the steel sheet increases as the carbides become finer. In order to obtain a high-strength hot-rolled steel sheet having a tensile strength of 850 MPa or more, it is necessary that the average particle diameter of the carbide is less than 10 nm. Preferably it is 6 nm or less.
次に、本発明熱延鋼板の成分組成の限定理由について説明する。なお、以下の成分組成を表す％は、特に断らない限り質量％（mass％）を意味するものとする。
C ：0.05％以上0.09％以下
Cは、Ti、或いは更にW、Mo、Vと結合し炭化物として鋼板中に微細分散する。すなわちCは、微細な炭化物を形成してフェライト組織を著しく強化させる元素であり、熱延鋼板を強化する上で必須の元素である。引張強さ850MPa以上の高強度熱延鋼板を得るには、C含有量を少なくとも0.05％以上とする必要がある。一方、後述するように本発明ではTiを多量に含有させるため、C含有量が0.09％を超えると、熱延鋼板を製造する工程の鋼素材（スラブ）再加熱時で粗大なTi炭化物を溶解しきれなくなり、最終的に得られる熱延鋼板に粗大なTi炭化物が残存してしまう。このように粗大なTi炭化物が残存すると、強度上昇に担う微細なTi炭化物の析出量が減少することで熱延鋼板の強度が急落するばかりか、粗大なTi炭化物とマトリックス界面でミクロボイドが生成し易くなり、鋼板の伸びフランジ性を劣化させることとなる。したがって、C含有量は0.05％以上0.09％以下とする。より望ましくは0.05％以上0.08％以下である。
Next, the reason for limiting the component composition of the hot-rolled steel sheet of the present invention will be described. In addition,% showing the following component composition shall mean the mass% (mass%) unless there is particular notice.
C: 0.05% or more and 0.09% or less
C combines with Ti, or even W, Mo, and V, and is finely dispersed in the steel sheet as a carbide. That is, C is an element that forms fine carbides and remarkably strengthens the ferrite structure, and is an essential element for strengthening the hot-rolled steel sheet. In order to obtain a high-strength hot-rolled steel sheet having a tensile strength of 850 MPa or more, the C content needs to be at least 0.05%. On the other hand, as described later, since a large amount of Ti is contained in the present invention, when the C content exceeds 0.09%, coarse Ti carbide is dissolved at the time of reheating the steel material (slab) in the process of manufacturing a hot-rolled steel sheet. It becomes impossible to squeeze and coarse Ti carbide remains in the finally obtained hot-rolled steel sheet. If coarse Ti carbide remains in this way, not only does the amount of fine Ti carbide precipitate that contributes to strength increase decrease, but the strength of the hot-rolled steel sheet drops sharply, and microvoids are generated at the coarse Ti carbide-matrix interface. It becomes easy to deteriorate the stretch flangeability of the steel sheet. Therefore, the C content is 0.05% or more and 0.09% or less. More desirably, it is 0.05% or more and 0.08% or less.
Si：0.3％以下
Siは、延性（伸び）低下をもたらすことなく鋼板強度を向上させる有効な元素として、従来の高強度鋼板では積極的に含有されている。しかしながら、Siは、鋼板表面に濃化し易く、この濃化により鋼板表面が部分的に硬化する。そして、このように表面が部分的に硬化した鋼板では、伸びフランジ成形時に表面での硬化部と硬化部以外との界面で応力集中を引き起こす結果、該界面でミクロボイドが発生し易くなり、伸びフランジ性が低下する。したがって、本発明ではSi含有量を極力低減することが望ましいが、0.3％までは許容できるため、Si含有量の上限を0.3％とする。好ましくは0.1％以下である。なお、Si含有量は不純物レベルまで低減してもよい。
Si: 0.3% or less
Si is actively contained in conventional high-strength steel sheets as an effective element for improving the steel sheet strength without reducing ductility (elongation). However, Si tends to be concentrated on the steel sheet surface, and the steel sheet surface is partially cured by this concentration. In such a steel sheet having a partially hardened surface, as a result of causing stress concentration at the interface between the hardened portion and the non-hardened portion on the surface during stretch flange molding, microvoids are likely to occur at the interface, and the stretch flange Sex is reduced. Therefore, in the present invention, it is desirable to reduce the Si content as much as possible, but up to 0.3% is acceptable, so the upper limit of the Si content is set to 0.3%. Preferably it is 0.1% or less. Note that the Si content may be reduced to the impurity level.
Mn：1.00％以下
Mn含有量は、本発明において、重要な要件のひとつである。Mnは、固溶強化元素であり、Siと同様、従来の高強度鋼板では積極的に含有されている。しかしながら、Mnは加工性を低下させる元素でもある。特に、Mnは、鋳造の際に不可避的に中心偏析を発生させ、この中心偏析部分は非常に硬質でありかつ延性に劣ることから鋼板の伸びフランジ成形時に偏析部と偏析しておらず硬化していない定常部分との界面で応力集中を引き起こし、伸びフランジ性を低下させる要因となる。そのため、本発明では、鋼板の伸びフランジ性を確保すべくMn含有量を低減し、1.00％以下とする必要がある。好ましくは0.85％以下である。
Mn: 1.00% or less
The Mn content is one of the important requirements in the present invention. Mn is a solid solution strengthening element and, like Si, is actively contained in conventional high-strength steel sheets. However, Mn is also an element that decreases workability. In particular, Mn inevitably generates center segregation during casting, and this center segregation part is very hard and inferior in ductility. This causes stress concentration at the interface with the unsteady stationary part, and causes a reduction in stretch flangeability. Therefore, in the present invention, the Mn content needs to be reduced to 1.00% or less in order to ensure stretch flangeability of the steel sheet. Preferably it is 0.85% or less.
但し、Mn含有量を極端に低減するとフェライト変態温度（変態点）が上昇し、熱延鋼板製造時、仕上げ圧延終了後の冷却過程でオーステナイト→フェライト変態と同時に析出する炭化物が高温に晒されることとなる。そして、このように炭化物が高温に晒されると炭化物は短時間で粗大化するため、最終的に得られる熱延鋼板の強度が低下する。このような観点から、Mn含有量は0.01％以上とすることが好ましい。特に、鋼のオーステナイト→フェライト変態点を低下させる効果を有するBを含有させない場合には、Mn含有量を0.66％以上とすることが好ましい。 However, if the Mn content is drastically reduced, the ferrite transformation temperature (transformation point) rises, and during the production of hot-rolled steel sheet, the carbide that precipitates simultaneously with the austenite → ferrite transformation is exposed to high temperatures during the cooling process after finish rolling. It becomes. And when a carbide | carbonized_material is exposed to high temperature in this way, since a carbide | carbonized_material will coarsen in a short time, the intensity | strength of the hot-rolled steel plate finally obtained will fall. From such a viewpoint, the Mn content is preferably set to 0.01% or more. In particular, when not containing B having the effect of lowering the austenite → ferrite transformation point of steel, the Mn content is preferably 0.66% or more.
P：0.03％以下
Pは、粒界に偏析して鋼板の伸びフランジ成形時に粒界割れの起点となり、伸びフランジ性を劣化させる有害な元素であるため、極力低減することが好ましい。そこで、本発明では上記問題点を回避すべく、P含有量を0.03％以下とする。好ましくは0.02％以下である。
P: 0.03% or less
P is a harmful element that segregates at the grain boundaries and becomes the starting point of grain boundary cracking during the stretch flange forming of the steel sheet, and deteriorates the stretch flangeability. Therefore, P is preferably reduced as much as possible. Therefore, in the present invention, in order to avoid the above problems, the P content is set to 0.03% or less. Preferably it is 0.02% or less.
S ：0.007％以下
Sは、鋼中で介在物として存在する。この介在物は、鋼板の打ち抜き時に楔状に伸び変形が不均一となるため、伸びフランジ性に著しい悪影響をもたらす。したがって、本発明では、S含有量を極力低減することが好ましく、0.007％以下とする。好ましくは0.004％以下である。
S: 0.007% or less
S exists as an inclusion in steel. This inclusion causes a remarkable adverse effect on the stretch flangeability because the deformation in the form of a wedge becomes uneven when the steel sheet is punched. Therefore, in the present invention, it is preferable to reduce the S content as much as possible, and the content is made 0.007% or less. Preferably it is 0.004% or less.
Al：0.1％以下
Alは、脱酸剤として作用する元素である。このような効果を得るためには0.02％以上含有することが望ましいが、Al含有量が0.1％を越えると介在物による伸びフランジ性への悪影響が顕在化する。したがって、Al含有量は0.1％以下とする。好ましくは0.08％以下である。
Al: 0.1% or less
Al is an element that acts as a deoxidizer. In order to obtain such an effect, it is desirable to contain 0.02% or more. However, if the Al content exceeds 0.1%, the adverse effect on the stretch flangeability due to inclusions becomes obvious. Therefore, the Al content is 0.1% or less. Preferably it is 0.08% or less.
N ：0.01％以下
Nは、製鋼の段階で炭化物形成元素であるTiと結合して粗大な窒化物を形成し、微細な炭化物の形成を阻害するため鋼板強度を著しく低下させる。また、粗大な窒化物は、ミクロボイド発生の原因となり、鋼板の伸びフランジ性も低下させる。したがってN含有量は極力低減することが好ましく、0.01％以下とする。好ましくは0.006％以下である。
N: 0.01% or less
N combines with Ti, which is a carbide-forming element, at the steelmaking stage to form coarse nitrides and inhibits the formation of fine carbides, so that the steel sheet strength is significantly reduced. In addition, coarse nitrides cause microvoids and reduce the stretch flangeability of the steel sheet. Therefore, the N content is preferably reduced as much as possible, and is 0.01% or less. Preferably it is 0.006% or less.
Ti：0.18％以上0.25％以下
Tiは、Cと炭化物を形成し、鋼板の高強度化に寄与する元素である。所望の熱延鋼板強度（引張強さ：850MPa以上）を確保するためには、Ti含有量を少なくとも0.18％以上とする必要がある。一方、Ti含有量が0.25％を超えると、熱延鋼板を製造する際、熱間圧延前の鋼素材（スラブ）加熱時に粗大なTi炭化物が完全に溶解せず、最終的に得られる（巻取り後の）熱延鋼板に粗大なTi炭化物が残存する。そして、この粗大なTi炭化物は、ミクロボイド発生の原因となるため、鋼板の伸びフランジ性を低下させる。したがって、Ti含有量は0.18％以上0.25％以下とする。好ましくは0.19％以上0.23％以下である。
Ti: 0.18% to 0.25%
Ti is an element that forms a carbide with C and contributes to an increase in strength of the steel sheet. In order to ensure the desired hot-rolled steel sheet strength (tensile strength: 850 MPa or more), the Ti content needs to be at least 0.18% or more. On the other hand, when the Ti content exceeds 0.25%, when manufacturing a hot-rolled steel sheet, coarse Ti carbide is not completely dissolved during heating of the steel material (slab) before hot rolling, and is finally obtained (winding) Coarse Ti carbide remains on the hot-rolled steel sheet after removal. And this coarse Ti carbide | carbonized_material causes a micro void generation, Therefore The stretch flangeability of a steel plate is reduced. Therefore, the Ti content is 0.18% or more and 0.25% or less. Preferably they are 0.19% or more and 0.23% or less.
本発明の熱延鋼板は、C、N、Tiを、上記した範囲で且つ(1)式を満足するように含有する。
1.0 ≦ ([C]/12)/([Ti*]/48) ≦ 1.5 ・・・ (1)
但し、[Ti*]＝[Ti]−3.4×[N]
（[C]、[N]、[Ti]：各元素の含有量（質量％））
上記(1)式は、熱延鋼板の引張強さを850MPa以上とするために満足すべき要件であり、本発明において重要な指標である。
The hot-rolled steel sheet of the present invention contains C, N, and Ti in the above-described range and satisfying the expression (1).
1.0 ≤ ([C] / 12) / ([Ti * ] / 48) ≤ 1.5 ... (1)
However, [Ti * ] = [Ti] −3.4 × [N]
([C], [N], [Ti]: Content of each element (mass%))
The above formula (1) is a requirement that must be satisfied in order to set the tensile strength of the hot-rolled steel sheet to 850 MPa or more, and is an important index in the present invention.
先述のとおり、本発明においては熱延鋼板中にTi炭化物を微細析出させることで所望の鋼板強度を確保する。ここで、Ti炭化物は、その平均粒子径が極めて小さい微細炭化物となる傾向が強いものの、鋼中に含まれるTiの原子濃度がCの原子濃度を超えると、Tiの炭化物析出能が急激に低下し、所望の熱延鋼板強度（引張強さ：850MPa以上）を確保することが困難となる。したがって、炭化物生成に寄与できるTiを「Ti*」とすると、本発明では、鋼素材に含まれるCの原子％（(Cの質量％)/12）を、Ti*の原子％（(Ti*の質量％)/48）以上にする必要がある。 As described above, in the present invention, the desired strength of the steel sheet is ensured by finely depositing Ti carbide in the hot-rolled steel sheet. Here, Ti carbide has a strong tendency to become fine carbide with a very small average particle diameter, but when the atomic concentration of Ti contained in the steel exceeds the atomic concentration of C, the Ti carbide precipitation ability rapidly decreases. In addition, it becomes difficult to ensure the desired hot-rolled steel sheet strength (tensile strength: 850 MPa or more). Therefore, when the Ti can contribute to the carbide generation and "Ti *" in the of the invention, C atomic percent contained in the steel material ((mass% of C) / 12) and, Ti * in atomic% ((Ti * %) / 48) or more.
また、後述のとおり本発明においては、鋼素材に所定量のTiを添加し、熱延前の加熱で鋼素材中の炭化物を溶解し、主に熱間圧延後の巻取り時にTi炭化物を析出させる。しかしながら、鋼素材に添加したTiの全量が炭化物生成に寄与するわけではなく、鋼素材に添加したTiの一部は、鋼板強度に寄与しない窒化物等の形成に消費される。巻取り温度よりも高温域では、Tiが炭化物よりも窒化物等を形成し易く、熱延鋼板の製造時、巻取り工程の前にTiが主に窒化物を形成するためである。そこで、上記組成を有する鋼素材について本発明者らが検討した結果、炭化物生成に寄与するTiであるTi*の量は、鋼素材に添加した全Ti量からTi窒化物の形成に消費されるTi量を差し引き、「[Ti]−3.4×[N]」で表現できることが明らかになった。 In addition, as described later, in the present invention, a predetermined amount of Ti is added to the steel material, the carbide in the steel material is dissolved by heating before hot rolling, and Ti carbide is precipitated mainly during winding after hot rolling. Let However, the total amount of Ti added to the steel material does not contribute to carbide generation, and a part of the Ti added to the steel material is consumed for forming nitrides and the like that do not contribute to the steel plate strength. This is because in the temperature range higher than the coiling temperature, Ti forms nitrides more easily than carbides, and Ti mainly forms nitrides before the winding process during the production of hot-rolled steel sheets. Therefore, as a result of investigations by the present inventors on a steel material having the above composition, the amount of Ti * , which is Ti that contributes to carbide formation, is consumed in the formation of Ti nitride from the total amount of Ti added to the steel material. It became clear that the amount of Ti can be subtracted and expressed as “[Ti] −3.4 × [N]”.
以上の理由により、本発明では、Cの原子％（[C]/12）を、Ti*（[Ti]−3.4×[N]）の原子％以上にする目的で、([C]/12)/([Ti*]/48)の値が1以上となるようにC、N、Tiの各元素を含有することとする。([C]/12)/([Ti*]/48)の値が1未満になると、フェライト結晶粒内に生成するTi炭化物の析出量が不十分となり、所望の熱延鋼板強度（引張強さ：850MPa以上）を得ることが困難となる。また、それに伴い炭化物の熱安定性が悪化し巻取工程時に炭化物が粗大化するため、所望の鋼板強度が得られなくなる。 For the above reasons, in the present invention, for the purpose of setting the atomic% of C ([C] / 12) to be equal to or higher than the atomic% of Ti * ([Ti] −3.4 × [N]), ([C] / 12 ) / ([Ti * ] / 48) contains C, N, and Ti elements so that the value is 1 or more. When the value of ([C] / 12) / ([Ti * ] / 48) is less than 1, the precipitation amount of Ti carbide generated in the ferrite crystal grains becomes insufficient, and the desired hot-rolled steel sheet strength (tensile strength) S: 850MPa or more) is difficult to obtain. In addition, the thermal stability of the carbide deteriorates accordingly, and the carbide becomes coarse during the winding process, so that a desired steel plate strength cannot be obtained.
一方、([C]/12)/([Ti*]/48)の値が1.5を超えると、鋼素材中に粗大なTi炭化物が生成し易くなる。そして、このような粗大なTi炭化物は、熱延鋼板の製造工程において、熱延前に鋼素材を加熱しても溶解せずに残存し、最終的に得られる熱延鋼板の強度低下の要因となる。
また、([C]/12)/([Ti*]/48)の値が1.5を超えると、鋼素材中のTi量に対するC量が過剰となる結果、セメンタイトが生成し易くなり、鋼板組織を実質的にフェライト単相組織とすることが困難となる。したがって、本発明では、([C]/12)/([Ti*]/48)の値が1.5以下となるようにC、N、Tiの各元素を含有することとする。好ましくは1.05以上1.48以下である。
On the other hand, if the value of ([C] / 12) / ([Ti * ] / 48) exceeds 1.5, coarse Ti carbides are easily generated in the steel material. Such coarse Ti carbide remains in the hot-rolled steel sheet manufacturing process without melting even if the steel material is heated before hot-rolling, and the cause of the strength reduction of the finally obtained hot-rolled steel sheet. It becomes.
In addition, when the value of ([C] / 12) / ([Ti * ] / 48) exceeds 1.5, the amount of C relative to the amount of Ti in the steel material becomes excessive, resulting in easy formation of cementite, and the steel sheet structure. Is substantially difficult to have a ferrite single-phase structure. Therefore, in the present invention, the elements of C, N, and Ti are contained so that the value of ([C] / 12) / ([Ti * ] / 48) is 1.5 or less. Preferably they are 1.05 or more and 1.48 or less.
以上が、本発明における基本成分であるが、上記した基本成分に加えてさらに(2)式を満足するようにB：0.003％以下を含有することができる。
B：0.003％以下
2.21×[Mn]＋1.05×log(104×([B]＋0.0001))≧1.44 ・・・ (2)
（[Mn]、[B]：各元素の含有量（質量％））
The above is the basic component in the present invention. In addition to the above basic component, B: 0.003% or less can be contained so as to satisfy the formula (2).
B: 0.003% or less
2.21 x [Mn] + 1.05 x log (10 4 x ([B] + 0.0001)) ≥ 1.44 (2)
([Mn], [B]: Content of each element (mass%))
Tiを含有する鋼素材を用いて熱延鋼板の製造する場合において、熱延前に鋼素材を加熱して鋼素材中の炭化物を溶解すると、Ti炭化物は通常、熱延後の冷却過程でオーステナイト→フェライト変態と同時に相界面析出する。ここで、鋼素材のオーステナイト→フェライト変態温度が高いと、熱間圧延後、Tiの拡散速度が速い高温域でTi炭化物が析出することになるため、Ti炭化物が粗大化し易くなる。一方、オーステナイト→フェライト変態の温度（Ar3変態点）を巻取り温度範囲（すなわち、Ti拡散速度が遅い温度域）まで低温化すれば、Ti炭化物の粗大化を効果的に抑制することができる。 When manufacturing a hot-rolled steel sheet using a Ti-containing steel material, if the steel material is heated before hot rolling to dissolve carbide in the steel material, Ti carbide is usually austenite in the cooling process after hot rolling. → Phase interface precipitates simultaneously with ferrite transformation. Here, when the austenite → ferrite transformation temperature of the steel material is high, Ti carbide precipitates in a high temperature range where the diffusion rate of Ti is high after hot rolling, and thus Ti carbide is easily coarsened. On the other hand, if the temperature of the austenite → ferrite transformation (Ar 3 transformation point) is lowered to the coiling temperature range (ie, the temperature range where the Ti diffusion rate is slow), the coarsening of Ti carbide can be effectively suppressed. .
そこで、本発明においては、鋼のオーステナイト→フェライト変態を遅延させ、Ti炭化物の析出温度（Ar3変態点）を後述する巻取り温度範囲まで安定的に低温化する目的で、上記した組成に加えてさらに、B ：0.003％以下を含有することができる。Bは、鋼のオーステナイト→フェライト変態温度を低下させる元素であり、本発明では、Bを添加して鋼のオーステナイト→フェライト変態温度を下げることによって、Ti炭化物の微細化を図ることができる。このような効果を得るためには、B含有量を0.0003％以上とすることが好ましい。一方、B含有量が0.003％を超えると、フェライト粒界の形状が複雑化し、粗大なBNが析出することで伸びフランジ性が低下する。したがって、B含有量は0.003％以下とすることが好ましい。より好ましくは0.0005％以上0.002％以下である。 Therefore, in the present invention, in order to delay the steel austenite → ferrite transformation and stably lower the Ti carbide precipitation temperature (Ar 3 transformation point) to the coiling temperature range described later, Furthermore, B: 0.003% or less can be contained. B is an element that lowers the austenite → ferrite transformation temperature of steel, and in the present invention, by adding B to lower the austenite → ferrite transformation temperature of steel, the Ti carbide can be refined. In order to obtain such an effect, the B content is preferably 0.0003% or more. On the other hand, if the B content exceeds 0.003%, the shape of the ferrite grain boundary becomes complicated, and coarse BN precipitates, so that stretch flangeability deteriorates. Therefore, the B content is preferably 0.003% or less. More preferably, it is 0.0005% or more and 0.002% or less.
先述のとおり、本発明では、鋼板の加工性（特に伸びフランジ性）に悪影響を及ぼすMnの含有量を低減することが好ましい。しかしながら、Mnは鋼のオーステナイト→フェライト変態温度を低下させるうえで有効な元素でもあり、鋼板の加工性を重視してMn含有量を大幅に低減すると、鋼のオーステナイト→フェライト変態を遅延させてTi炭化物の析出温度（Ar3変態点）を後述する巻取り温度範囲まで安定的に低温化することが極めて困難となる。したがって、本発明において、鋼板の加工性（特に伸びフランジ性）を確保する目的でMn含有量を低減する場合には、Bを積極的に含有させることが好ましい。 As described above, in the present invention, it is preferable to reduce the content of Mn that adversely affects the workability (particularly stretch flangeability) of the steel sheet. However, Mn is also an element effective in lowering the austenite → ferrite transformation temperature of steel, and if the Mn content is greatly reduced with emphasis on the workability of the steel sheet, the austenite → ferrite transformation of the steel is delayed and Ti is delayed. It is extremely difficult to stably lower the carbide precipitation temperature (Ar 3 transformation point) to the coiling temperature range described later. Therefore, in the present invention, in order to reduce the Mn content for the purpose of ensuring the workability (particularly stretch flangeability) of the steel sheet, it is preferable to positively contain B.
また、上記の組成を有する鋼素材を用いて熱延鋼板を製造するに際し、鋼のオーステナイト→フェライト変態を遅延させてTi炭化物の析出温度（Ar3変態点）を後述する巻取り温度範囲（700℃以下）まで安定的に低温化する手段について本発明者らが検討した結果、(2)式（2.21×[Mn]＋1.05×log(104×([B]＋0.0001))≧1.44、但し[Mn]、[B]は各元素の含有量（質量％））を満足するようにMnとBの含有量を規定することが好ましいことを知見した。 Moreover, when manufacturing a hot-rolled steel sheet using a steel material having the above composition, the austenite-to-ferrite transformation of the steel is delayed, and the Ti carbide precipitation temperature (Ar 3 transformation point) is described later in the coiling temperature range (700 As a result of the study by the present inventors on the means for stably lowering the temperature to below (° C.), the formula (2) (2.21 × [Mn] + 1.05 × log (10 4 × ([B] +0.0001)) ≧ 1.44 However, [Mn] and [B] were found to be preferable to define the contents of Mn and B so as to satisfy the content (mass%) of each element.
(2)式の左辺（2.21×[Mn]＋1.05×log(104×([B]＋0.0001))において、Mn、Bそれぞれの係数は、鋼のオーステナイト→フェライト変態点調整能を反映している。そして、(2)式の左辺の値が1.44以上であると、鋼のオーステナイト→フェライト変態点が後述する巻取り温度以下（700℃以下）となり、Ti炭化物の粗大化を抑制できることとなる。(2)式から明らかであるように、Bを含有しない場合には、Mn含有量を0.66％以上とすることが好ましい。なお、Mnのみ、もしくはBのみでもオーステナイト→フェライト変態点を所望の温度域に調整可能であるが、より優れた伸びフランジ性を有する高強度熱延鋼板を得るためには、Mn含有量を0.4％以下にまで低減するとともに、Bを複合添加することが好ましい。 In the left side of equation (2) (2.21 x [Mn] + 1.05 x log (10 4 x ([B] + 0.0001)), the coefficients of Mn and B are the ability to adjust the austenite to ferrite transformation point of steel. And, if the value on the left side of equation (2) is 1.44 or more, the austenite → ferrite transformation point of steel becomes below the coiling temperature (700 ° C or below), which will be described later, and the coarsening of Ti carbide is suppressed. As is clear from the formula (2), when B is not contained, the Mn content is preferably set to 0.66% or more, and austenite → ferrite transformation point with Mn alone or B alone. In order to obtain a high-strength hot-rolled steel sheet with better stretch flangeability, the Mn content should be reduced to 0.4% or less, and B should be added in combination. Is preferred.
また、上記した基本成分に加えてさらに、V：0.005％以上0.3％以下を含有することができる。Vは、Cと結合して炭化物を形成し、或いはTiおよびCと結合して複合炭化物を形成し、熱延鋼板の更なる強化に寄与する元素である。このような効果を得るためには、V含有量を0.005％以上とすることが好ましい。一方、V含有量が0.3％を超えると、炭化物粗大化の原因となり、鋼板強度が低下するおそれがある。したがって、V含有量は0.005％以上0.3％以下とすることが好ましい。また、0.03％以上0.2％以下とすることがより好ましい。 Further, in addition to the above basic components, V: 0.005% or more and 0.3% or less can be contained. V is an element that contributes to further strengthening of the hot-rolled steel sheet by combining with C to form a carbide, or combining with Ti and C to form a composite carbide. In order to obtain such an effect, the V content is preferably 0.005% or more. On the other hand, if the V content exceeds 0.3%, it may cause carbide coarsening and the steel sheet strength may be reduced. Therefore, the V content is preferably 0.005% or more and 0.3% or less. Further, it is more preferably 0.03% or more and 0.2% or less.
また、上記した基本組成に加えてさらに、W ：0.01％以上1.0％以下、Mo：0.01％以上0.5％以下のいずれか１種以上を含有することができる。Vと同様にW、Moは、TiおよびCと結合して複合炭化物を形成し、熱延鋼板の更なる強化に寄与する元素である。このような効果を得るためには、W ：0.01％以上、Mo：0.01％以上とすることが好ましい。一方、W含有量が1.0％、Mo含有量が0.5％を超えると、熱延鋼板製造時、熱間圧延終了後の冷却工程に続くコイル巻取り時に鋼のオーステナイト→フェライト変態が完了せず、実質的にフェライト単相組織の熱延鋼板が得られなくなり、伸びフランジ性が低下する。したがって、W含有量は0.01％以上1.0％以下とすることが好ましく、Mo含有量は0.01％以上0.5％以下とすることが好ましい。また、W含有量は0.02％以上0.2％以下、Mo含有量は0.02％以上0.1％以下とすることがより好ましい。 Further, in addition to the basic composition described above, any one or more of W: 0.01% to 1.0% and Mo: 0.01% to 0.5% can be contained. Like V, W and Mo are elements that combine with Ti and C to form a composite carbide and contribute to further strengthening of the hot-rolled steel sheet. In order to obtain such an effect, W: 0.01% or more and Mo: 0.01% or more are preferable. On the other hand, if the W content exceeds 1.0% and the Mo content exceeds 0.5%, the steel austenite → ferrite transformation is not completed during coil winding following the cooling process after hot rolling, A hot-rolled steel sheet having a ferrite single phase structure cannot be obtained substantially, and stretch flangeability is deteriorated. Therefore, the W content is preferably 0.01% or more and 1.0% or less, and the Mo content is preferably 0.01% or more and 0.5% or less. More preferably, the W content is 0.02% to 0.2%, and the Mo content is 0.02% to 0.1%.
また、前記基本成分に加えてさらに、Se、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、Ni、Cr、Sb、Cu、Sn、Mg、Caのいずれか１種以上を合計で1.0％以下含有してもよい。これら元素は、伸びフランジ性の観点から合計で1.0％までは許容できる。好ましくは合計で0.5％以下とする。上記以外の成分は、Feおよび不可避的不純物である。 In addition to the basic components, Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd, Hg, Ag, Au, Pd, Pt, Co, Rh, Ir, Any one or more of Ru, Os, Tc, Re, Ta, Be, Sr, REM, Ni, Cr, Sb, Cu, Sn, Mg, and Ca may be contained in a total of 1.0% or less. These elements are acceptable up to 1.0% in total from the viewpoint of stretch flangeability. Preferably, the total content is 0.5% or less. Components other than the above are Fe and inevitable impurities.
引張強さ：850MPa以上1150MPa以下
自動車車体の軽量化には鋼板強度は大きい方が望ましい一方、鋼板強度が過剰に高くなると鋼板の加工性が低下し、自動車部材を所定形状に成形する際の加工が困難になる等、様々な支障をきたす。したがって、熱延鋼板の引張強さは850MPa以上1150MPa以下とする。
Tensile strength: 850MPa or more and 1150MPa or less Higher steel sheet strength is desirable for reducing the weight of an automobile body. On the other hand, if the steel sheet strength becomes excessively high, the workability of the steel sheet deteriorates, and processing when forming automotive parts into a predetermined shape Causes various problems such as difficulty. Therefore, the tensile strength of the hot-rolled steel sheet is 850 MPa to 1150 MPa.
本発明の熱延鋼板は、後述する巻取り温度の上限である700℃までの加熱処理を施しても材質変動が小さい。そのため、鋼板に耐食性を付与する目的で、本発明の熱延鋼板にめっき処理を施し、その表面にめっき層を具えることができる。めっき処理における加熱温度は700℃以下でも製造可能であることから、本発明の熱延鋼板にめっき処理を施しても前記した本発明の効果を損なうことはない。めっき層の種類は特に問わず、電気めっき層、無電解めっき層のいずれも適用可能である。また、めっき層の合金成分も特に問わず、溶融亜鉛めっき層、合金化溶融亜鉛めっき層などが好適な例として挙げられるが、勿論、これらに限定されず従前公知のものがいずれも適用可能である。 The hot-rolled steel sheet of the present invention has little material fluctuation even when subjected to heat treatment up to 700 ° C., which is the upper limit of the coiling temperature described later. Therefore, for the purpose of imparting corrosion resistance to the steel sheet, the hot-rolled steel sheet of the present invention can be plated, and a plating layer can be provided on the surface thereof. Since it can be produced even at a heating temperature of 700 ° C. or less in the plating process, the effect of the present invention described above is not impaired even if the hot-rolled steel sheet of the present invention is plated. The type of the plating layer is not particularly limited, and any of an electroplating layer and an electroless plating layer can be applied. Further, the alloy component of the plating layer is not particularly limited, and a hot dip galvanized layer, an alloyed hot dip galvanized layer, and the like can be cited as suitable examples. is there.
次に、本発明の熱延鋼板の製造方法について説明する。
本発明は、上記した組成の鋼素材を加熱し、粗圧延と仕上げ圧延からなる熱間圧延を施し、仕上げ圧延終了後、冷却し、巻き取り、熱延鋼板とする。この際、前記加熱の加熱温度を1150℃以上1350℃以下とし、前記仕上げ圧延の仕上げ圧延温度を850℃以上とし、前記冷却を仕上げ圧延終了後3秒以内に開始し、前記冷却の平均冷却速度を30℃/s以上とし、前記巻き取りの巻取り温度を500℃以上700℃以下とすることを特徴とする。
Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.
In the present invention, the steel material having the above composition is heated, subjected to hot rolling consisting of rough rolling and finish rolling, cooled after completion of finish rolling, and wound into a hot-rolled steel sheet. At this time, the heating temperature of the heating is 1150 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the average cooling rate of the cooling Is 30 ° C./s or more, and the winding temperature of the winding is 500 ° C. or more and 700 ° C. or less.
本発明において、鋼の溶製方法は特に限定されず、転炉、電気炉等、公知の溶製方法を採用することができる。また、真空脱ガス炉にて2次精錬を行ってもよい。その後、生産性や品質上の問題から連続鋳造法により鋼素材とするのが好ましいが、造塊−分塊圧延法、薄スラブ連鋳法等、公知の鋳造方法でスラブとしても良い。なお、本発明においては、Mnを削減したことから700℃以上の絞り性（変形能）が良好であるため、連続鋳造による製造が容易となる。 In the present invention, the method for melting steel is not particularly limited, and a known melting method such as a converter or an electric furnace can be employed. Further, secondary refining may be performed in a vacuum degassing furnace. Then, it is preferable to use a steel material by a continuous casting method from the viewpoint of productivity and quality, but a slab may be formed by a known casting method such as an ingot-bundling rolling method or a thin slab continuous casting method. In the present invention, since Mn is reduced, the drawability (deformability) of 700 ° C. or higher is good, and therefore, the production by continuous casting becomes easy.
鋼素材の加熱温度：1150℃以上1350℃以下
上記の如く得られた鋼素材（鋼スラブ）に、粗圧延および仕上げ圧延を施すが、本発明においては、粗圧延に先立ち鋼素材を加熱して実質的に均質なオーステナイト相とし、粗大な炭化物を溶解する必要がある。鋼素材の加熱温度が1150℃を下回ると、粗大なTi炭化物が溶解しないため、熱間圧延終了後の冷却・巻取り工程で微細分散する炭化物の量が減じることとなり、最終的に得られる熱延鋼板の強度が著しく低下する。一方、上記加熱温度が1350℃を上回ると、スケールが噛み込み、鋼板表面性状を悪化させる。
Heating temperature of steel material: 1150 ° C or higher and 1350 ° C or lower The steel material (steel slab) obtained as described above is subjected to rough rolling and finish rolling. In the present invention, the steel material is heated prior to rough rolling. It is necessary to form a substantially homogeneous austenite phase and dissolve coarse carbides. When the heating temperature of the steel material is below 1150 ° C, coarse Ti carbide does not dissolve, so the amount of carbide finely dispersed in the cooling and winding process after hot rolling is reduced, and the heat finally obtained The strength of the rolled steel sheet is significantly reduced. On the other hand, when the heating temperature exceeds 1350 ° C., the scale bites and deteriorates the surface properties of the steel sheet.
以上の理由により、鋼素材の加熱温度は1150℃以上1350℃以下とする。好ましくは1200℃以上1320℃以下である。但し、鋼素材に熱間圧延を施すに際し、鋳造後の鋼素材が1150℃以上1350℃以下の温度域にある場合、或いは鋼素材の炭化物が溶解し、しかもオーステナイト単相域の温度となっている場合には、鋼素材を加熱することなく直送圧延してもよい。なお、粗圧延条件については特に限定されない。 For the above reasons, the heating temperature of the steel material is set to 1150 ° C or higher and 1350 ° C or lower. Preferably they are 1200 degreeC or more and 1320 degrees C or less. However, when hot rolling the steel material, if the steel material after casting is in the temperature range of 1150 ° C or higher and 1350 ° C or lower, or the carbide of the steel material is melted, and the temperature becomes the austenite single phase range. If it is, direct rolling may be performed without heating the steel material. The rough rolling conditions are not particularly limited.
仕上げ圧延温度：850℃以上
仕上げ圧延温度が850℃を下回ると、仕上げ圧延中にフェライト変態が開始してフェライト粒が伸展された組織となるうえ、部分的にフェライト粒が成長した混粒組織となるため、熱延鋼板の伸びフランジ性が低下する。したがって、仕上げ圧延温度は850℃以上とする。好ましくは890℃以上である。なお、仕上げ圧延温度の上限は特に定めないが、仕上げ圧延温度は熱間圧延前の加熱温度と通板速度、鋼板板厚により、自ずと決定される。この観点から仕上げ圧延温度は実質的に980℃以下である。
Finishing rolling temperature: 850 ° C or higher When the finishing rolling temperature is lower than 850 ° C, the ferrite transformation starts during finish rolling, and the ferrite grains are expanded, and the mixed grain structure in which the ferrite grains partially grow Therefore, the stretch flangeability of the hot-rolled steel sheet is lowered. Therefore, the finish rolling temperature is 850 ° C. or higher. Preferably it is 890 degreeC or more. The upper limit of the finish rolling temperature is not particularly defined, but the finish rolling temperature is naturally determined by the heating temperature before hot rolling, the sheet feeding speed, and the steel plate thickness. From this viewpoint, the finish rolling temperature is substantially 980 ° C. or lower.
仕上げ圧延終了後、強制冷却を開始するまでの時間：3s秒以内
仕上げ圧延直後の高温状態の鋼板においては、オーステナイト相に蓄積されたひずみエネルギーが大きいため、ひずみ誘起析出による炭化物が生じる。この炭化物は、高温で析出することから粗大化し易いため、ひずみ誘起析出が生じると微細な析出物が得られ難くなる。したがって、本発明では、ひずみ誘起析出を抑制する目的で熱間圧延終了後速やかに強制冷却を開始する必要があり、仕上げ圧延終了後、少なくとも3s以内に冷却を開始する。好ましくは2s以内である。
Time until start of forced cooling after finish rolling: within 3 s seconds In a high-temperature steel sheet immediately after finish rolling, carbides are generated due to strain-induced precipitation because the strain energy accumulated in the austenite phase is large. Since this carbide precipitates at a high temperature and is easy to coarsen, when strain-induced precipitation occurs, it becomes difficult to obtain a fine precipitate. Therefore, in the present invention, for the purpose of suppressing strain-induced precipitation, forced cooling needs to be started immediately after the end of hot rolling, and cooling is started within at least 3 seconds after the end of finish rolling. Preferably it is within 2 s.
平均冷却速度：30℃/s以上
上記のとおり、仕上げ圧延終了後の鋼板の高温に維持される時間が長いほど、ひずみ誘起析出による炭化物の粗大化が進行し易くなる。また、本発明においては、所定の鋼組成に規定することでオーステナイト→フェライト変態を抑制しているものの、冷却速度が小さいと高温でフェライト変態が開始し、炭化物が粗大化し易くなる。そのため、仕上げ圧延後は急冷する必要があり、上記問題を回避するには30℃/s以上の平均冷却速度で冷却する必要がある。好ましくは50℃/s以上である。但し、仕上げ圧延終了後の冷却速度が過剰に大きくなると、巻取温度の制御が困難となることが懸念されるため、150℃/s以下とすることが好ましい。なお、平均冷却速度は、強制冷却開始から冷却停止までの平均冷却速度である。強制冷却停止後は空冷だけで鋼板は冷却されるので、冷却停止後、鋼板の温度はわずかしか低下せずに鋼板は巻き取られる。通常、冷却停止温度は、巻取り温度＋5〜10℃程度に設定される。
Average cooling rate: 30 ° C./s or more As described above, the longer the time during which the steel sheet after finish rolling is maintained at a high temperature, the more easily the coarsening of the carbide by strain-induced precipitation proceeds. In the present invention, the austenite → ferrite transformation is suppressed by defining a predetermined steel composition. However, when the cooling rate is low, the ferrite transformation starts at a high temperature, and the carbide tends to be coarsened. Therefore, it is necessary to rapidly cool after finish rolling, and in order to avoid the above problem, it is necessary to cool at an average cooling rate of 30 ° C./s or more. Preferably, it is 50 ° C./s or more. However, if the cooling rate after finishing rolling is excessively increased, there is a concern that it is difficult to control the coiling temperature. Therefore, the temperature is preferably set to 150 ° C./s or less. The average cooling rate is the average cooling rate from the start of forced cooling to the stop of cooling. After the forced cooling is stopped, the steel sheet is cooled only by air cooling. Therefore, after the cooling is stopped, the temperature of the steel sheet is slightly decreased and the steel sheet is wound up. Usually, the cooling stop temperature is set to the coiling temperature +5 to 10 ° C.
巻取り温度：500℃以上700℃以下
巻取り温度が500℃を下回ると十分な量の炭化物が得られず、鋼板強度が低下する。一方、巻取り温度が700℃を超えると、析出した炭化物が粗大化するため鋼板強度が低下する。したがって、巻取温度の範囲は500℃以上700℃以下とする。好ましくは550℃以上660℃以下である。
Winding temperature: 500 ° C. or higher and 700 ° C. or lower When the winding temperature is lower than 500 ° C., a sufficient amount of carbide cannot be obtained, and the steel sheet strength decreases. On the other hand, when the coiling temperature exceeds 700 ° C., the precipitated carbide is coarsened, so that the steel sheet strength is lowered. Accordingly, the coiling temperature range is 500 ° C. or more and 700 ° C. or less. Preferably they are 550 degreeC or more and 660 degrees C or less.
なお、熱間圧延した巻き取り後の熱延鋼板は、表面にスケールが付着した状態であっても、酸洗を行うことによりスケールを除去した状態であっても、その特性が変わることはなく、いずれの状態においても前記した優れた特性を発現する。また、本発明では、巻き取り後の熱延鋼板にめっき処理を施して、熱延鋼板表面にめっき層を形成してもよい。
めっき処理の種類は特に問わず、電気めっき処理、無電解めっき処理のいずれも適用可能である。例えば、めっき処理として溶融亜鉛めっき処理を施して溶融亜鉛めっき層を形成することができる。或いは、上記溶融亜鉛めっき処理後、更に合金化処理を施して合金化溶融亜鉛めっき層を形成してもよい。また、溶融めっきには亜鉛の他に、アルミもしくはアルミ合金等、その他の金属や合金をめっきすることもできる。
Note that the hot-rolled steel sheet after being rolled by hot rolling does not change its properties even if the scale is attached to the surface or the scale is removed by pickling. The above-described excellent characteristics are exhibited in any state. In the present invention, the hot-rolled steel sheet after winding may be plated to form a plating layer on the surface of the hot-rolled steel sheet.
The type of the plating treatment is not particularly limited, and both electroplating treatment and electroless plating treatment are applicable. For example, a hot dip galvanizing process can be performed as a plating process to form a hot dip galvanized layer. Alternatively, after the hot dip galvanizing treatment, an alloying treatment may be further performed to form an alloyed hot dip galvanized layer. In addition to zinc, other metals and alloys such as aluminum or aluminum alloy can be plated for hot dipping.
本発明により得られる熱延鋼板は、700℃以下までの温度域であれば析出物の状態が変わることはない。そのため、例えば焼鈍温度を700℃以下とした連続めっきラインに通板させることができる。めっき層の付着方法としては、例えば、めっき浴に鋼板を浸漬して引き上げる方法などが挙げられる。合金化処理方法としては、例えば、めっき処理後にガス炉など、鋼板表面を加熱することができる炉内で行う方法などが挙げられる。 The hot-rolled steel sheet obtained by the present invention does not change the state of precipitates in the temperature range up to 700 ° C. or less. Therefore, for example, it can be passed through a continuous plating line with an annealing temperature of 700 ° C. or lower. Examples of the method for attaching the plating layer include a method in which a steel sheet is immersed in a plating bath and pulled up. Examples of the alloying method include a method performed in a furnace capable of heating the steel sheet surface such as a gas furnace after plating.
表１に示す組成を有する肉厚250mmのスラブ（鋼素材）に、表２に示す熱延条件で板厚1.2〜3.2mmの熱延鋼板とした。なお、表２に記載の冷却速度は、強制冷却開始から冷却停止までの平均冷却速度である。また、得られた熱延鋼板の一部に対しては、焼鈍温度700℃の溶融亜鉛めっきラインに通板し、その後、460℃のめっき浴（めっき組成：Zn-0.13mass％Al）に浸漬し、溶融亜鉛めっき材（GI材）とした。また、一部の溶融亜鉛めっき材（GI材）に対しては、溶融亜鉛めっきライン通板、めっき浴浸漬に次いで530℃で合金化処理を施してめっき材（GA材）とした。めっき付着量はGI材、GA材ともに片面当たり45g/m2とした。
なお、鋼板No.3〜7、20〜25を除き、巻き取りまでの冷却中にオーステナイトからフェライトへの変態は生じていないことを、別途確認している。
A 250 mm thick slab (steel material) having the composition shown in Table 1 was used as a hot rolled steel sheet having a thickness of 1.2 to 3.2 mm under the hot rolling conditions shown in Table 2. In addition, the cooling rate described in Table 2 is an average cooling rate from the start of forced cooling to the stop of cooling. In addition, a part of the obtained hot-rolled steel sheet is passed through a hot dip galvanizing line with an annealing temperature of 700 ° C and then immersed in a 460 ° C plating bath (plating composition: Zn-0.13 mass% Al). And a hot-dip galvanized material (GI material). Further, a part of the hot dip galvanized material (GI material) was subjected to alloying treatment at 530 ° C. after passing through the hot dip galvanizing line and plating bath so as to obtain a plated material (GA material). The plating adhesion amount was 45 g / m 2 per side for both GI and GA materials.
In addition, except steel plates Nos. 3 to 7 and 20 to 25, it was separately confirmed that transformation from austenite to ferrite did not occur during cooling to winding.
上記により得られた熱延鋼板（熱延鋼板、GI材、GA材）から試験片を採取し、組織観察、引張試験、穴拡げ試験を行い、フェライト相の面積率、フェライト相以外の組織の種類および面積率、フェライト相の平均結晶粒径、炭化物の平均粒子径、降伏強度、引張強さ、伸び、穴拡げ率（伸びフランジ性）を求めた。試験方法は次のとおりとした。 Samples are taken from the hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material) obtained as described above, and subjected to structure observation, tensile test, and hole expansion test. The type and area ratio, average grain size of ferrite phase, average grain size of carbide, yield strength, tensile strength, elongation, and hole expansion rate (stretch flangeability) were determined. The test method was as follows.
（i）組織観察
フェライト相の面積率は以下の手法により評価した。圧延方向に平行な断面の板厚中心部について、5％ナイタールによる腐食現出組織を走査型光学顕微鏡で400倍に拡大して10視野分撮影した。フェライト相は粒内に腐食痕やセメンタイトが観察されない形態を有する組織である。また、ポリゴナルフェライト、ベイニティックフェライト、アシキュラーフェライトおよびグラニュラーフェライトをフェライトとして面積率や粒径を求めた。フェライト相の面積率は、画像解析によりフェライト相とベイナイトやマルテンサイト等のフェライト相以外とを分離し、観察視野に対するフェライト相の面積率によって求めた。このとき、線状の形態として観察される粒界はフェライト相の一部として計上した。
フェライト相の平均結晶粒径は、上記の代表的な写真3枚について水平線および垂直線をそれぞれ3本ずつ引きASTM E 112-10に準拠した切断法によって求め、最終的に3枚の平均値を表３に記した。
フェライト相の結晶粒内の炭化物の平均粒子径は、得られた熱延鋼板の板厚中央部から薄膜法によってサンプルを作製し、透過型電子顕微鏡（倍率：135000倍）で観察を行い、100点以上の析出物粒子径の平均によって求めた。この析出物粒子径を算出する上で、粒子径が1.0μmより大きい粗大なセメンタイトや窒化物は含まないものとした。
(I) Microstructure observation The area ratio of the ferrite phase was evaluated by the following method. About the central part of the plate thickness with a cross section parallel to the rolling direction, the corrosion appearance structure by 5% nital was magnified 400 times with a scanning optical microscope and photographed for 10 fields of view. The ferrite phase is a structure having a form in which corrosion marks and cementite are not observed in the grains. Further, the area ratio and particle size were determined using polygonal ferrite, bainitic ferrite, acicular ferrite and granular ferrite as ferrite. The area ratio of the ferrite phase was determined by separating the ferrite phase from those other than the ferrite phase such as bainite and martensite by image analysis, and obtaining the area ratio of the ferrite phase with respect to the observation field. At this time, the grain boundary observed as a linear form was counted as a part of the ferrite phase.
The average crystal grain size of the ferrite phase is determined by drawing a horizontal line and a vertical line for each of the above three representative photographs by a cutting method in accordance with ASTM E 112-10. It was described in Table 3.
The average particle size of the carbides in the ferrite phase grains was measured using a transmission electron microscope (magnification: 135000 times) by preparing a sample from the center of the thickness of the obtained hot rolled steel sheet using a thin film method. It calculated | required by the average of the particle diameter of the precipitate more than a point. In calculating the particle diameter of the precipitate, coarse cementite and nitride having a particle diameter larger than 1.0 μm were not included.
（ii）引張試験
得られた熱延鋼板から圧延方向と垂直方向にJIS13号B引張試験片を作製し、JIS Z 2241(2011)の規定に準拠した引張試験を5回行い、平均の降伏強度(YS)、引張強さ(TS)、全伸び(El)を求めた。引張試験のクロスヘッドスピードは10mm/minとした。
(Ii) Tensile test JIS No. 13 B tensile test specimen was produced from the obtained hot-rolled steel sheet in the direction perpendicular to the rolling direction, and the tensile test was conducted five times in accordance with the provisions of JIS Z 2241 (2011). (YS), tensile strength (TS), and total elongation (El) were determined. The crosshead speed in the tensile test was 10 mm / min.
（iii）穴拡げ試験（伸びフランジ性評価）
得られた熱延鋼板からサンプルを採取して穴拡げ試験を行い、伸びフランジ性評価を行った。試験条件は日本鉄鋼連盟規格(T1001-1996)に準拠し、100W×100L mmのサンプル中央にクリアランス12％とした直径10mmの打抜加工を行い、頂角60°の円錐台のポンチを用いた。また、各サンプルについて5回繰り返し試験を行い、次式で算出される穴広げ率（λ）の平均値を求めた。なお、次式において「試験後孔径」は、打抜加工によって得られた初期孔径（直径10mm）に円錐台のポンチを挿入し、該孔を押し広げ、亀裂が熱延鋼板（試験片）を貫通したときの孔の径である。
（穴広げ率λ％）＝（試験後孔径−初期孔径）/（初期孔径）×100
得られた結果を表３に示す。
(Iii) Hole expansion test (stretch flangeability evaluation)
A sample was taken from the obtained hot-rolled steel sheet and subjected to a hole expansion test to evaluate stretch flangeability. The test conditions were in accordance with the Japan Iron and Steel Federation standard (T1001-1996), punched with a diameter of 10 mm with a clearance of 12% at the center of a 100 W x 100 L mm sample, and a punch with a truncated cone with a vertex angle of 60 ° was used. . Each sample was repeatedly tested 5 times, and the average value of the hole expansion ratio (λ) calculated by the following equation was obtained. In the following equation, the “post-test hole diameter” means that a truncated cone punch is inserted into the initial hole diameter (diameter 10 mm) obtained by punching, the hole is expanded, and cracks are generated in the hot-rolled steel sheet (test piece). It is the diameter of the hole when penetrating.
(Hole expansion ratio λ%) = (pore diameter after test−initial pore diameter) / (initial pore diameter) × 100
The obtained results are shown in Table 3.
本発明例（鋼板No.1,2,8〜19）はいずれも、引張強さTS：850MPa以上、1150MPa以下であり且つ優れた伸びフランジ性を有し、強度と加工性を兼備した熱延鋼板となっている。また、本発明例の熱延鋼板では、Mn含有量が少ないほど、より優れた伸びフランジ性を有することがわかる。一方、本発明の範囲を外れる比較例（鋼板No.3〜7,20〜25）は、所定の高強度が確保できていないか、十分な穴拡げ率が得られていない。 Examples of the present invention (steel plates Nos. 1, 2, 8 to 19) all have a tensile strength TS: 850 MPa or more and 1150 MPa or less, have excellent stretch flangeability, and have both strength and workability. It is a steel plate. Moreover, in the hot-rolled steel sheet of this invention example, it turns out that it has the more excellent stretch flangeability, so that Mn content is small. On the other hand, in the comparative examples (steel plates No. 3 to 7, 20 to 25) outside the scope of the present invention, a predetermined high strength is not ensured or a sufficient hole expansion rate is not obtained.
Claims (13)
C ：0.05％以上0.09％以下、 Si：0.3％以下、
Mn：1.00％以下、 P ：0.03％以下、
S ：0.007％以下、 Al：0.1％以下、
N ：0.01％以下、 Ti：0.18％以上0.25％以下
を、C、NおよびTiが下記(1)を満足するように含有し、残部がFeおよび不可避的不純物からなる組成と、フェライト相の面積率が95％以上、該フェライト相の平均結晶粒径が10μm以下であり、前記フェライト相の結晶粒内の炭化物平均粒子径が10nm未満である組織を有し、引張強さが850MPa以上1150MPa以下であることを特徴とする、伸びフランジ性に優れた高強度熱延鋼板。
記
1.0 ≦ ([C]/12)/([Ti*]/48) ≦ 1.5 ・・・ (1)
但し、[Ti*]＝[Ti]−3.4×[N]
（[C]、[N]、[Ti]：各元素の含有量（質量％）） % By mass
C: 0.05% or more and 0.09% or less, Si: 0.3% or less,
Mn: 1.00% or less, P: 0.03% or less,
S: 0.007% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.18% or more and 0.25% or less, so that C, N and Ti satisfy the following (1), with the balance consisting of Fe and inevitable impurities, and the area of the ferrite phase The ratio is 95% or more, the ferrite phase has an average crystal grain size of 10 μm or less, and has a structure in which the average grain size of carbide in the ferrite phase grains is less than 10 nm, and the tensile strength is 850 MPa or more and 1150 MPa or less. A high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized by being
Record
1.0 ≤ ([C] / 12) / ([Ti * ] / 48) ≤ 1.5 ... (1)
However, [Ti * ] = [Ti] −3.4 × [N]
([C], [N], [Ti]: Content of each element (mass%))
記
2.21×[Mn]＋1.05×log(104×([B]＋0.0001))≧1.44 ・・・ (2)
（[Mn]、[B]：各元素の含有量（質量％）） The high-strength heat excellent in stretch flangeability according to claim 1, further comprising, in addition to the composition, B: 0.003% or less by mass% so as to satisfy the following formula (2): Rolled steel sheet.
Record
2.21 x [Mn] + 1.05 x log (10 4 x ([B] + 0.0001)) ≥ 1.44 (2)
([Mn], [B]: Content of each element (mass%))
前記鋼素材を、質量％で、
C ：0.05％以上0.09％以下、 Si：0.3％以下、
Mn：1.00％以下、 P ：0.03％以下、
S ：0.007％以下、 Al：0.1％以下、
N ：0.01％以下、 Ti：0.18％以上0.25％以下
を、C、NおよびTiが下記(1)を満足するように含有し、残部がFeおよび不可避的不純物からなる組成とし、
前記加熱の加熱温度を1150℃以上1350℃以下とし、前記仕上げ圧延の仕上げ圧延温度を850℃以上とし、前記冷却を仕上げ圧延終了後3秒以内に開始し、前記冷却の平均冷却速度を30℃/s以上とし、前記巻き取りの巻取り温度を500℃以上700℃以下とすることを特徴とする、伸びフランジ性に優れた高強度熱延鋼板の製造方法。
記
1.0 ≦ ([C]/12)/([Ti*]/48) ≦ 1.5 ・・・ (1)
但し、[Ti*]＝[Ti]−3.4×[N]
（[C]、[N]、[Ti]：各元素の含有量（質量％）） The steel material is heated, subjected to hot rolling consisting of rough rolling and finish rolling, and after finishing rolling, cooled, wound, and hot rolled steel sheet,
The steel material in mass%,
C: 0.05% or more and 0.09% or less, Si: 0.3% or less,
Mn: 1.00% or less, P: 0.03% or less,
S: 0.007% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.18% or more and 0.25% or less, so that C, N and Ti satisfy the following (1), the balance is composed of Fe and inevitable impurities,
The heating temperature of the heating is 1150 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the average cooling rate of the cooling is 30 ° C. The method for producing a high-strength hot-rolled steel sheet excellent in stretch flangeability, characterized in that the winding temperature is 500 ° C. or higher and 700 ° C. or lower.
Record
1.0 ≤ ([C] / 12) / ([Ti * ] / 48) ≤ 1.5 ... (1)
However, [Ti * ] = [Ti] −3.4 × [N]
([C], [N], [Ti]: Content of each element (mass%))
記
2.21×[Mn]＋1.05×log(104×([B]＋0.0001))≧1.44 ・・・ (2)
（[Mn]、[B]：各元素の含有量（質量％）） The high-strength heat excellent in stretch flangeability according to claim 9, further comprising, in addition to the composition, B: 0.003% or less by mass% so as to satisfy the following formula (2): A method for producing rolled steel sheets.
Record
2.21 x [Mn] + 1.05 x log (10 4 x ([B] + 0.0001)) ≥ 1.44 (2)
([Mn], [B]: Content of each element (mass%))
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