JP4736617B2 - High-strength, high-strength cold-rolled steel sheet and method for producing the same - Google Patents
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本発明は、自動車のセンターピラーなどのコラム状の構造部材に用いられる590MPa以上の引張強度TSを有する高強度冷延鋼板、特に、剛性の高い高強度冷延鋼板、およびその製造方法に関する。 The present invention relates to a high-strength cold-rolled steel sheet having a tensile strength TS of 590 MPa or more used for a columnar structural member such as a center pillar of an automobile, in particular, a high-rigidity high-strength cold-rolled steel sheet and a method for producing the same.
近年、地球環境問題への関心の高まりを受けて、自動車でも排ガス規制が要請されており、自動車車体の軽量化が極めて重要な課題になっている。自動車車体を軽量化するには、車体に使用される鋼板を高強度化し、その板厚を減少させることが有効である。そのため、最近では、自動車のセンターピラー、ロッカー、サイドフレーム、クロスメンバーなどのコラム状の構造部材に対しても、590MPa以上のTSを有し、板厚が2.0mmを下回るような高強度冷延鋼板を積極的に適用しようという動きがある。 In recent years, in response to increasing interest in global environmental issues, there is a demand for exhaust gas regulations for automobiles, and weight reduction of automobile bodies has become an extremely important issue. In order to reduce the weight of an automobile body, it is effective to increase the strength of a steel plate used for the vehicle body and reduce the thickness of the steel plate. For this reason, recently, high-strength cold rolling with a TS of 590 MPa or more and a plate thickness of less than 2.0 mm is also used for columnar structural members such as automobile center pillars, rockers, side frames, and cross members. There is a movement to actively apply steel plates.
これまで、590MPa以上のTSを有する高強度冷延鋼板として、例えば、質量%で、C:0.04〜0.14%、Si:0.4〜2.2%、Mn:1.2〜2.4%、P:0.02%以下、S:0.01%以下、Al:0.002〜0.5%、Ti:0.005〜0.1%、N:0.006%以下、さらにNb、Mo、Vの1種以上を合計で0.005〜0.1%を含有し、Ti/S≧5を満足し、残部がFeおよび不可避的不純物からなる組成のスラブを、仕上温度(900+50×Si)以上で熱間圧延し、圧下率50〜85%で冷間圧延後、700〜900℃のフェライトとオーステナイトの二相共存温度域で10s〜5min焼鈍し、700℃から500℃までの間の平均冷却速度を1〜120℃/sとして250〜500℃に冷却し、必要に応じて再加熱した後250〜600℃の範囲の温度域に30s〜10min保持してから冷却し、マルテンサイトおよびオーステナイトの体積率を合計で6%以上とし、かつマルテンサイト、残留オーステナイトおよびベイナイトの硬質相の体積率α%を、α≦50000×(Ti/48+Nb/93+Mo/96+V/51)となるようにさせた穴広げ性に優れた低降伏比高強度冷延鋼板が開示されている(特許文献1)。 Up to now, as high-strength cold-rolled steel sheet having TS of 590 MPa or more, for example, in mass%, C: 0.04 to 0.14%, Si: 0.4 to 2.2%, Mn: 1.2 to 2.4%, P: 0.02% or less, S : 0.01% or less, Al: 0.002 to 0.5%, Ti: 0.005 to 0.1%, N: 0.006% or less, further containing one or more of Nb, Mo, V in a total amount of 0.005 to 0.1%, Ti / S ≧ 5 and a slab having a composition consisting of Fe and inevitable impurities in the balance is hot-rolled at a finishing temperature (900 + 50 × Si) or higher, and after cold rolling at a reduction rate of 50 to 85%, 700 to 900 Annealed for 10 s to 5 min in the two-phase coexistence temperature range of ferrite and austenite at ℃, cooled to 250 to 500 ℃ with an average cooling rate between 700 ℃ and 500 ℃ to 1 to 120 ℃ / s, as needed After reheating, hold in the temperature range of 250 to 600 ° C for 30 s to 10 min and then cool to make the total volume ratio of martensite and austenite 6% or more, and the hard phase of martensite, residual austenite and bainite Low yield ratio high-strength cold-rolled steel sheet with excellent hole expansibility, in which the volume ratio α% of α ≦ 50000 × (Ti / 48 + Nb / 93 + Mo / 96 + V / 51) is disclosed (Patent Document 1).
また、重量%で、C:0.02〜0.30%、Si:1.50%以下、Mn:0.60〜3.0%、P:0.20%以下、S:0.05%以下、Al:0.01〜0.10%、さらにMo:0.01〜1.0%、Cr:0.10〜1.5%、Nb:0.01〜0.05%、Ti:0.01〜0.5%、B:0.0005〜0.0030%、Ca:0.01%以下のうちの少なくとも1種を含有し、残部がFeおよび不可避的不純物からなる組成の鋼を、Ac3変態点以上で熱間圧延し、酸洗、冷間圧延後、連続焼鈍溶融亜鉛めっきラインにおいて、再結晶温度以上かつAc1変態点以上に加熱保持後、溶融亜鉛槽に至るまでの間において、Ms点以下に急冷して、鋼板全体あるいは部分的にマルテンサイトを生成させ、次いで、Ms点以上の温度であって少なくとも溶融亜鉛浴温度および合金化炉温度に加熱して全体あるいは部分的に焼戻マルテンサイトを生成させた伸びフランジ性に優れた高強度合金化溶融亜鉛めっき鋼板が開示されている(特許文献2)。 Also, by weight%, C: 0.02 to 0.30%, Si: 1.50% or less, Mn: 0.60 to 3.0%, P: 0.20% or less, S: 0.05% or less, Al: 0.01 to 0.10%, Mo: 0.01 to 1.0%, Cr: 0.10 to 1.5%, Nb: 0.01 to 0.05%, Ti: 0.01 to 0.5%, B: 0.0005 to 0.0030%, Ca: containing at least one of 0.01% or less, with the balance being Fe and Steel with an inevitable impurity composition is hot-rolled at the Ac 3 transformation point or higher, pickled and cold-rolled, and then heated and maintained above the recrystallization temperature and above the Ac 1 transformation point in a continuous annealing hot-dip galvanizing line After that, in the period up to the molten zinc bath, it is rapidly cooled below the Ms point to generate martensite as a whole or part of the steel sheet, and then at least the molten zinc bath temperature and alloying at a temperature above the Ms point. A high-strength alloyed hot-dip galvanized steel sheet having excellent stretch flangeability, which is heated to a furnace temperature to generate tempered martensite in whole or in part, is disclosed ( Patent Document 2).
なお、下記の特許文献3は、後述の[課題を解決するための手段]で述べる鋼のヤング率に関する。
しかしながら、特許文献1や特許文献2に記載の高強度冷延鋼板を、実際にこうしたコラム状の構造部材に適用すると、十分な剛性を確保できない、すなわち、撓んだり、ねじれたりする問題が起こる。 However, when the high-strength cold-rolled steel sheet described in Patent Document 1 or Patent Document 2 is actually applied to such a column-shaped structural member, sufficient rigidity cannot be secured, that is, problems such as bending or twisting occur. .
本発明は、剛性が高く、TSが590MPa以上の高強度冷延鋼板およびその製造方法を提供することを目的とする。 An object of the present invention is to provide a high-strength cold-rolled steel sheet having high rigidity and TS of 590 MPa or more and a method for producing the same.
一般に、鋼板の剛性を高めるにはヤング率を高めることが有効である。また、鋼のヤング率は、集合組織に大きく依存し、体心立方格子である普通鋼の場合は、原子の最密方向である<111>方向に高く、逆に原子密度の小さい<100>方向に低いので、{112}<110>方位を発達させれば鋼板の圧延方向と直角方向のヤング率を高めることができる。また、{112}<110>方位の発達した冷延鋼板とするには、熱間圧延後に{113}<110>方位を発達させることが有効である(例えば、特許文献3)。 In general, increasing the Young's modulus is effective for increasing the rigidity of the steel sheet. In addition, the Young's modulus of steel greatly depends on the texture, and in the case of plain steel with a body-centered cubic lattice, the atomic density is high in the <111> direction, and conversely, the atomic density is small <100> Since the direction is low, if the {112} <110> orientation is developed, the Young's modulus in the direction perpendicular to the rolling direction of the steel sheet can be increased. In order to obtain a cold rolled steel sheet with {112} <110> orientation, it is effective to develop the {113} <110> orientation after hot rolling (for example, Patent Document 3).
そこで、本発明者らは、成形性の観点から低温変態相で強化された高強度冷延鋼板の剛性を向上させるべく種々の検討を行ったところ、上記の{112}<110>方位を発達させてヤング率を高め、さらに下記の式(1)のXを230GPa以上にすることが有効であることを見出した;
X=0.005×TS/(ε2-ε1)[GPa] ・・・(1)
ここで、TSは鋼板の引張強度[MPa]、ε1は公称応力が0.40×TSのときの公称歪[%]、ε2は公称応力が0.45×TSのときの公称歪[%]を表す。
Therefore, the present inventors conducted various studies to improve the rigidity of the high-strength cold-rolled steel sheet reinforced with the low-temperature transformation phase from the viewpoint of formability, and developed the above {112} <110> orientation. It was found that it is effective to increase the Young's modulus and further set X in the following formula (1) to 230 GPa or more;
X = 0.005 × TS / (ε2-ε1) [GPa] (1)
Here, TS is the tensile strength [MPa] of the steel sheet, ε1 is the nominal strain [%] when the nominal stress is 0.40 × TS, and ε2 is the nominal strain [%] when the nominal stress is 0.45 × TS.
本発明は、このような知見に基づきなされたもので、質量%で、C:0.02〜0.20%、Si:1.5%以下、Mn:1.0〜3.5%、P:0.05%以下、S:0.01%以下、Al:1.5%以下、N:0.01%以下、Nb:0.01〜0.40%、残部Feおよび不可避的不純物からなり、[Nb]-(93/14)×[N]≧0.005を満足し、200〜500℃の温度域に60s以上滞留させて焼戻して得たフェライト相と低温変態相からなるミクロ組織を有し、かつ低温変態相の体積率が1〜60%であり、引張強度が590MPa以上で、かつ上記の式(1)のXが230GPa以上である剛性の高い高強度冷延鋼板を提供する。ここで、[M]は元素Mの含有量[質量%]を表す。 The present invention has been made on the basis of such knowledge, in mass%, C: 0.02 to 0.20%, Si: 1.5% or less, Mn: 1.0 to 3.5%, P: 0.05% or less, S: 0.01% or less Al: 1.5% or less, N: 0.01% or less, Nb: 0.01 to 0.40%, balance Fe and inevitable impurities, satisfying [Nb]-(93/14) × [N] ≧ 0.005, 200 to It has a microstructure composed of a ferrite phase and a low-temperature transformation phase obtained by tempering by staying in a temperature range of 500 ° C. for 60 s or more, a volume ratio of the low-temperature transformation phase is 1 to 60%, and a tensile strength is 590 MPa or more. In addition, a high-strength, high-strength cold-rolled steel sheet having X of 230 GPa or more in the above formula (1) is provided. Here, [M] represents the content [% by mass] of the element M.
本発明の高強度冷延鋼板では、Nb:0.03〜0.40%にすることが好ましい。 In the high-strength cold-rolled steel sheet of the present invention, Nb is preferably 0.03 to 0.40%.
本発明の高強度冷延鋼板は、さらに、質量%で、V:0.01〜0.5%、Ti:0.01〜0.2%のうち少なくとも1つの元素を含有することができる。 The high-strength cold-rolled steel sheet of the present invention can further contain at least one element of V: 0.01 to 0.5% and Ti: 0.01 to 0.2% by mass%.
また、本発明の高強度冷延鋼板は、さらに、質量%で、Mo:0.1〜1.0%、Cr:0.1〜1.0%、Ni:0.1〜1.0%、B:0.0005〜0.0020%のうち少なくとも1つの元素、あるいはCu:0.1〜2.0%を含有することもできる。 Further, the high-strength cold-rolled steel sheet of the present invention is further in mass%, Mo: 0.1 to 1.0%, Cr: 0.1 to 1.0%, Ni: 0.1 to 1.0%, B: 0.0005 to 0.0020%. It can also contain an element or Cu: 0.1-2.0%.
本発明の高強度冷延鋼板は、例えば、上記のような組成からなる鋼スラブを、950℃以下における総圧下量を30%以上とし、Ar3変態点〜900℃の仕上温度で熱間圧延し、650℃以下で巻取り、40%以上の圧下率で冷間圧延後、(875-400×[C]-30×[Mn])〜(915-400×[C]-30×[Mn])℃で1〜600s加熱し、次いで冷却して700〜500℃の温度域に20s以上滞留させるとともに200〜500℃の温度域に60s以上滞留させて焼戻すか、あるいは前記700〜500℃の温度域での滞留の後200℃以下に冷却し、再加熱して200〜500℃の温度域に60s以上滞留させて焼戻す条件で焼鈍する方法により製造可能である。ここで、[M]は元素Mの含有量[質量%]を表す。 The high-strength cold-rolled steel sheet of the present invention is, for example, a steel slab having the above composition, with a total reduction amount of 30% or more at 950 ° C. or less, and hot rolling at a finishing temperature of Ar 3 transformation point to 900 ° C. Then, after winding at 650 ° C. or less and cold rolling at a reduction rate of 40% or more, (875-400 × [C] -30 × [Mn]) to (915-400 × [C] -30 × [Mn] ]) Heated at 1 to 600 s for 1 to 600 s, then cooled and retained in the temperature range of 700 to 500 ° C. for 20 s or more and retained in the temperature range of 200 to 500 ° C. for 60 s or tempered, or 700 to 500 ° C. It can be manufactured by a method in which it is cooled to 200 ° C. or lower after being retained in the temperature range, reheated and retained in the temperature range of 200 to 500 ° C. for 60 seconds or more and annealed under the conditions of tempering. Here, [M] represents the content [% by mass] of the element M.
本発明により、自動車のセンターピラーなどのコラム状の構造部材に適用しても、剛性を確保できる590MPa以上のTSを有する高強度冷延鋼板を製造できるようになった。 The present invention makes it possible to produce a high-strength cold-rolled steel sheet having a TS of 590 MPa or more that can ensure rigidity even when applied to a columnar structural member such as a center pillar of an automobile.
以下に、本発明である高強度冷延鋼板およびその製造方法について詳細に説明する。 Below, the high-strength cold-rolled steel sheet and its manufacturing method which are this invention are demonstrated in detail.
1)成分
C:Cは、オーステナイト安定化元素なので、冷間圧延後の焼鈍時の冷却過程において焼入れ性を高め、低温変態相の生成を促進して高強度化に大きく寄与する。また、Ar3変態点を低下させるので、より低温域での熱間圧延を可能にし、未再結晶オーステナイトからのフェライト変態を促進して{113}<110>方位を発達させ、焼鈍時に{112}<110>方位を発達させてヤング率を向上させる。このような効果を得るためには、その量を0.02%以上とする必要がある。一方、その量が0.20%を超えると、硬質な低温変態相が増加して成形性が劣化するとともに、フェライト相が減少するためヤング率が低下する。したがって、C量は0.02〜0.20%、好ましくは0.02〜0.10%とする。
1) ingredients
Since C: C is an austenite stabilizing element, it enhances hardenability in the cooling process during annealing after cold rolling, promotes the formation of a low-temperature transformation phase, and greatly contributes to high strength. In addition, since the Ar 3 transformation point is lowered, it enables hot rolling in a lower temperature range, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and {112} } Develop <110> orientation to improve Young's modulus. In order to obtain such an effect, the amount needs to be 0.02% or more. On the other hand, if the amount exceeds 0.20%, the hard low-temperature transformation phase increases and the formability deteriorates, and the ferrite phase decreases, so the Young's modulus decreases. Therefore, the C content is 0.02 to 0.20%, preferably 0.02 to 0.10%.
Si:Siは、Ar3変態点を著しく上昇させ、熱間圧延中における加工オーステナイトの再結晶を促進するため、1.5%を超えて多量に含有すると高ヤング率化に必要な結晶方位を発達させることができなくなり、また、鋼板の溶接性、化成処理性、めっき性を劣化させたり、熱間圧延時に赤スケールと呼ばれる表面欠陥の発生を助長させる。したがって、Si量は1.5%以下とする。特に、本発明の鋼板が良好な表面性状を必要とする場合や本発明の鋼板に溶融亜鉛めっきを施す場合には、Si量を0.5%以下とすることが好ましい。なお、Siはフェライト安定化元素であり、焼鈍時の冷却過程においてフェライト変態を促進し、オーステナイト中にCを濃化させてオーステナイトを安定化させ、低温変態相の生成を促進する効果を有する。この効果を十分に得るためには、Si量を0.1%以上とすることが望ましい。 Si: Si significantly raises the Ar 3 transformation point and promotes recrystallization of processed austenite during hot rolling, so if it contains more than 1.5%, it develops the crystal orientation necessary for high Young's modulus In addition, it deteriorates the weldability, chemical conversion properties and plating properties of the steel sheet, and promotes the occurrence of surface defects called red scale during hot rolling. Therefore, the Si content is 1.5% or less. In particular, when the steel sheet of the present invention requires good surface properties, or when hot dip galvanizing is performed on the steel sheet of the present invention, the Si content is preferably 0.5% or less. Si is a ferrite stabilizing element and has the effect of promoting ferrite transformation in the cooling process during annealing, concentrating C in austenite to stabilize austenite, and promoting the formation of a low temperature transformation phase. In order to sufficiently obtain this effect, it is desirable that the Si content be 0.1% or more.
Mn:Mnは、本発明において重要な元素の1つであり、熱間圧延時に加工オーステナイトの再結晶を抑制するとともに、Ar3変態点を低下させるので、より低温域での熱間圧延を可能にし、未再結晶オーステナイトからのフェライト変態を促進して{113}<110>方位を発達させて、焼鈍後のヤング率を向上させる。また、Mnは、オーステナイト安定化元素でもあるので、焼鈍時の昇温過程においてAc1変態点を低下させ、未再結晶フェライトからのオーステナイト変態を促進し、その後の冷却過程において高ヤング率化に有利なフェライトの方位を発達させる。さらに、Mnは、固溶強化元素であるとともに、焼鈍時の冷却過程において焼入れ性を高め、低温変態相の生成を促進するので、高強度化に大きく寄与する。このような効果を得るためには、その量を1.0%以上とする必要がある。一方、その量が3.5%を超えると、焼鈍時の冷却過程で高ヤング率化に必要なフェライトの生成が著しく抑制され、また、低温変態相が増加して鋼が極端に高強度化して成形性が劣化するとともに、溶接性も劣化する。したがって、Mn量は1.0〜3.5%、好ましくは1.5〜2.5%とする。 Mn: Mn is one of the important elements in the present invention, and it suppresses recrystallization of processed austenite during hot rolling and lowers the Ar 3 transformation point, enabling hot rolling at lower temperatures. And promote the ferrite transformation from unrecrystallized austenite to develop the {113} <110> orientation and improve the Young's modulus after annealing. Mn is also an austenite stabilizing element, so it lowers the Ac 1 transformation point during the temperature rise process during annealing, promotes austenite transformation from unrecrystallized ferrite, and increases the Young's modulus in the subsequent cooling process. Develop advantageous ferrite orientation. Furthermore, Mn is a solid solution strengthening element, and enhances hardenability in the cooling process during annealing and promotes the formation of a low-temperature transformation phase, thus greatly contributing to an increase in strength. In order to obtain such an effect, the amount needs to be 1.0% or more. On the other hand, if the amount exceeds 3.5%, the formation of ferrite necessary for high Young's modulus during the cooling process during annealing is remarkably suppressed, and the low-temperature transformation phase increases, resulting in extremely high strength steel. As well as deterioration of weldability, weldability also deteriorates. Therefore, the Mn content is 1.0 to 3.5%, preferably 1.5 to 2.5%.
P:Pは、0.05%を超えて含有すると粒界に偏析して鋼板の延性や靭性を低下させるとともに、溶接性を劣化させる。また、本発明の鋼板に合金化溶融亜鉛めっきを施す場合には、Pは合金化速度を遅滞させる。したがって、P量は0.05%以下とする。なお、Pは固溶強化元素であり、フェライトを安定化してオーステナイト中へのC濃化を促進する作用や、Siを添加した鋼において赤スケールの発生を抑制する作用も有する。そのため、P量は0.01%以上とすることが好ましい。 When P: P is contained in an amount exceeding 0.05%, it segregates at the grain boundaries and lowers the ductility and toughness of the steel sheet and degrades the weldability. Moreover, when alloying hot dip galvanizing to the steel plate of this invention, P delays alloying speed | rate. Therefore, the P content is 0.05% or less. P is a solid solution strengthening element, and has the effect of stabilizing ferrite and promoting C concentration in austenite, and the effect of suppressing the generation of red scale in steel to which Si is added. Therefore, the P content is preferably 0.01% or more.
S:Sは、0.01%を超えて多量に含有すると熱間での延性を著しく低下させて熱間割れを誘起し、鋼板の表面性状を著しく劣化させる。また、強度にほとんど寄与しないばかりか、粗大なMnSとして析出し、穴広げ性などの延性を低下させる。したがって、S量は0.01%以下とする。なお、その量は少ないほど好ましいが、穴広げ性を向上させる観点からは0.005%以下とすることがより好ましい。 When S: S is contained in a large amount exceeding 0.01%, the hot ductility is remarkably lowered to induce hot cracking, and the surface properties of the steel sheet are remarkably deteriorated. Moreover, it not only contributes to the strength, but also precipitates as coarse MnS, reducing the ductility such as hole expandability. Therefore, the S content is 0.01% or less. The smaller the amount, the better. However, from the viewpoint of improving the hole expansion property, 0.005% or less is more preferable.
Al:Alは、フェライト安定化元素であり、鋼のAr3変態点を大きく上昇させるため加工オーステナイトの再結晶を促進して、高ヤング率化に必要な結晶方位の発達を抑制する。また、多量に含有されるとオーステナイト単相域が消失し、熱間圧延時にオーステナイト域で圧延を終了させることが困難となる。したがって、Al量は1.5%以下とする。なお、Alは、焼鈍時の冷却過程においてフェライト生成を促進し、オーステナイト中にCを濃化させてオーステナイトを安定化させ、低温変態相の生成を促進する効果を有するので、Al量は0.2%以上とすることが望ましい。 Al: Al is a ferrite stabilizing element, and promotes recrystallization of processed austenite to greatly increase the Ar 3 transformation point of steel, thereby suppressing the development of crystal orientation necessary for increasing the Young's modulus. Moreover, when it contains abundantly, an austenite single phase area | region will lose | disappear, and it will become difficult to complete | finish rolling in an austenite area | region at the time of hot rolling. Therefore, the Al content is 1.5% or less. In addition, Al has the effect of promoting ferrite formation in the cooling process during annealing, concentrating C in austenite to stabilize austenite, and promoting the formation of low-temperature transformation phase, so the Al amount is 0.2% It is desirable to set it above.
N:Nは、0.01%を超えて多量に含有されると熱間圧延中にスラブ割れを誘起し、鋼板の表面性状を劣化させる恐れがある。したがって、N量は0.01%以下とする。 If N: N is contained in a large amount exceeding 0.01%, slab cracking may be induced during hot rolling, and the surface properties of the steel sheet may be deteriorated. Therefore, the N content is 0.01% or less.
Nb:Nbは、本発明における最も重要な元素である。すなわち、Nbは、熱間圧延時に加工オーステナイトの再結晶を抑制し、未再結晶オーステナイトからのフェライト変態を促進して{113}<110>方位を発達させ、焼鈍後のヤング率を向上させる。また、焼鈍時の昇温過程において加工フェライトの再結晶を抑制し、未再結晶フェライトからのオーステナイト変態を促進して、その後の冷却過程において高ヤング率化に有利なフェライトの方位を発達させる。さらに、微細な炭窒化物として析出し、強度上昇にも寄与する。このような効果を得るためには、その量を0.01%以上、好ましくは0.03%以上とする必要がある。一方、その量が0.40%を超えると、スラブの再加熱時にその炭窒化物はすべて固溶することができず、粗大な炭窒化物が残るため熱間圧延時において加工オーステナイトの再結晶が抑制されず、また、冷間圧後の焼鈍時においても加工フェライトの再結晶が抑制されないので、高ヤング率化に有利なフェライトの方位を発達させることができない。また、連続鋳造後スラブを直接熱間圧延する場合においても、加工オーステナイトの再結晶が抑制されない。したがって、Nb量は0.01〜0.40%、好ましくは0.03〜0.40%とする。さらに、Nb窒化物は、Nb炭化物に比べ高温で析出するため熱間圧延時に粗大になりやすいので、加工オーステナイトの再結晶を抑制したり、焼鈍時に加工フェライトの再結晶を抑制したりする効果が小さく、高ヤング率化に対する寄与は少ない。したがって、炭化物として析出するNb量を確保する必要があるので、[Nb]-(93/14)×[N]≧0.005を満足させる必要がある。 Nb: Nb is the most important element in the present invention. That is, Nb suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves the Young's modulus after annealing. In addition, the recrystallization of the processed ferrite is suppressed during the temperature rising process during annealing, the austenite transformation from the unrecrystallized ferrite is promoted, and the orientation of the ferrite advantageous for increasing the Young's modulus is developed in the subsequent cooling process. Furthermore, it precipitates as fine carbonitride and contributes to an increase in strength. In order to obtain such an effect, the amount needs to be 0.01% or more, preferably 0.03% or more. On the other hand, if the amount exceeds 0.40%, all of the carbonitride cannot be dissolved at the time of reheating the slab, and coarse carbonitride remains, so recrystallization of processed austenite is suppressed during hot rolling. In addition, since recrystallization of the processed ferrite is not suppressed even during annealing after cold pressing, it is not possible to develop a ferrite orientation that is advantageous for increasing the Young's modulus. Even when the slab is directly hot-rolled after continuous casting, recrystallization of the processed austenite is not suppressed. Therefore, the Nb content is 0.01 to 0.40%, preferably 0.03 to 0.40%. In addition, Nb nitride precipitates at a higher temperature than Nb carbide, so it tends to be coarse during hot rolling, so it has the effect of suppressing recrystallization of processed austenite and suppressing recrystallization of processed ferrite during annealing. Small and has little contribution to increasing Young's modulus. Therefore, since it is necessary to secure the amount of Nb precipitated as carbide, it is necessary to satisfy [Nb] − (93/14) × [N] ≧ 0.005.
残部は、Feおよび不可避的不純物である。 The balance is Fe and inevitable impurities.
上記成分元素に加え、下記の理由により、質量%で、V:0.01〜0.5%、Ti:0.01〜0.2%のうち少なくとも1つの元素を含有させることが好ましい。 In addition to the above component elements, it is preferable to contain at least one element of V: 0.01 to 0.5% and Ti: 0.01 to 0.2% by mass% for the following reasons.
V:Vは、微細な炭窒化物として析出し、強度上昇に寄与する。そのためには、その量を0.01%とする必要がある。一方、その量が0.5%を超えても強度上昇効果は小さく、コスト増も招く。したがって、V量は0.01〜0.5%とする。 V: V precipitates as fine carbonitride and contributes to an increase in strength. For that purpose, the amount needs to be 0.01%. On the other hand, if the amount exceeds 0.5%, the effect of increasing the strength is small and the cost is increased. Therefore, the V amount is 0.01 to 0.5%.
Ti:Tiは、微細な炭窒化物として析出し、強度上昇に寄与する。また、熱間圧延時に加工オーステナイトの再結晶を抑制し、未再結晶オーステナイトからのフェライト変態を促進して高ヤング率化に寄与する。このような作用を有するためには、その量を0.01%以上とする必要がある。一方、その量が0.2%を超えると、スラブの再加熱時に炭窒化物はすべて固溶することができず、粗大な炭窒化物が残るため、強度を上昇させたり、加工オーステナイトの再結晶を抑制させる効果が得られない。また、連続鋳造後スラブを直接熱間圧延する場合においても、加工オーステナイトの再結晶を抑制する効果は小さく、コスト増を招く。したがって、Ti量は0.01〜0.2%とする。 Ti: Ti precipitates as fine carbonitrides and contributes to an increase in strength. Also, recrystallization of processed austenite during hot rolling is suppressed, and ferrite transformation from unrecrystallized austenite is promoted, thereby contributing to a higher Young's modulus. In order to have such an effect, the amount needs to be 0.01% or more. On the other hand, if the amount exceeds 0.2%, the carbonitride cannot completely dissolve when the slab is reheated, and coarse carbonitride remains, so that the strength is increased or recrystallization of the processed austenite is performed. The effect of suppressing cannot be obtained. Even when the slab is directly hot-rolled after continuous casting, the effect of suppressing recrystallization of the processed austenite is small, resulting in an increase in cost. Therefore, the Ti amount is set to 0.01 to 0.2%.
さらに、下記の理由により、質量%で、Mo:0.1〜1.0%、Cr:0.1〜1.0%、Ni:0.1〜1.0%、B:0.0005〜0.0020%のうち少なくとも1つの元素、あるいはCu:0.1〜2%を含有させることが好ましい。 Further, for the following reasons, at least one element of Mo: 0.1-1.0%, Cr: 0.1-1.0%, Ni: 0.1-1.0%, B: 0.0005-0.0020%, or Cu: 0.1- It is preferable to contain 2%.
Mo:Moは、界面の移動度を小さくすることにより焼入れ性を高める元素であり、焼鈍時の冷却過程において低温変態相の生成を促進して高強度化に大きく寄与する。また、熱間圧延時に加工オーステナイトの再結晶を抑制し、未再結晶オーステナイトからのフェライト変態を促進して{113}<110>方位を発達させ、焼鈍後のヤング率を向上させる。このような作用を得るためには、その量を0.1%以上とする必要がある。一方、その量が1.0%を超えると、その効果が飽和するだけでなく、コスト増を招く。したがって、Mo量は0.1〜1.0%とする。 Mo: Mo is an element that increases the hardenability by reducing the mobility of the interface, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to an increase in strength. It also suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves Young's modulus after annealing. In order to obtain such an action, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 1.0%, the effect is not only saturated but also the cost is increased. Therefore, the Mo amount is 0.1 to 1.0%.
Cr:Crは、セメンタイトの生成を抑制して焼入れ性を高める元素であり、焼鈍時の冷却過程において低温変態相の生成を促進して高強度化に大きく寄与する。また、熱間圧延時に加工オーステナイトの再結晶を抑制し、未再結晶オーステナイトからのフェライト変態を促進して{113}<110>方位を発達させ、焼鈍後のヤング率を向上させる。このような効果を得るには、その量を0.1%以上とする必要がある。一方、その量が1.0%を超えると、その効果が飽和するだけでなく、コスト増を招く。したがって、Cr量は0.1〜1.0%とする。なお、本発明の鋼板に溶融亜鉛めっきを施す場合には、表面に生成するCrの酸化物が不めっきを誘発するので、Cr量は0.5%以下とすることがより好ましい。 Cr: Cr is an element that suppresses the formation of cementite and enhances hardenability, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to an increase in strength. It also suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves Young's modulus after annealing. In order to obtain such an effect, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 1.0%, the effect is not only saturated but also the cost is increased. Therefore, the Cr content is 0.1 to 1.0%. When hot dip galvanizing is applied to the steel sheet of the present invention, the Cr oxide produced on the surface induces non-plating, so the Cr content is more preferably 0.5% or less.
Ni:Niは、オーステナイトを安定化することで焼入れ性を高める元素であり、焼鈍時の冷却過程において低温変態相の生成を促進して高強度化に大きく寄与する。また、焼鈍時の昇温過程においてAc1変態点を低下させ、未再結晶フェライトからのオーステナイト変態を促進し、冷却過程において高ヤング率化に有利なフェライトの方位を発達させる。さらに、熱間圧延時に加工オーステナイトの再結晶を抑制するとともに、Ar3変態点を低下させ、より低温域での熱間圧延を可能にすることで、未再結晶オーステナイトからのフェライト変態を促進して{113}<110>方位を発達させ、焼鈍後のヤング率を向上させる。さらにまた、Cu添加の場合に起こる熱間圧延時の割れを防止する。このような作用を得るためには、その量を0.1%以上とする必要がある。一方、その量が1.0%を超えると、焼鈍時の冷却過程で高ヤング率化に必要なフェライトの生成を抑制したり、低温変態相を増加させて鋼を極端に高強度化し、成形性を低下させる。したがって、Ni量は0.1〜1.0%とする。 Ni: Ni is an element that enhances hardenability by stabilizing austenite, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing and greatly contributes to high strength. In addition, it lowers the Ac 1 transformation point in the temperature raising process during annealing, promotes austenite transformation from unrecrystallized ferrite, and develops a ferrite orientation that is advantageous for increasing the Young's modulus in the cooling process. In addition, it suppresses recrystallization of processed austenite during hot rolling, lowers the Ar 3 transformation point, and enables hot rolling in a lower temperature range, thereby promoting ferrite transformation from unrecrystallized austenite. To develop the {113} <110> orientation and improve the Young's modulus after annealing. Furthermore, cracks during hot rolling that occur when Cu is added are prevented. In order to obtain such an action, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 1.0%, the formation of ferrite necessary for increasing the Young's modulus during the cooling process during annealing is suppressed, or the low-temperature transformation phase is increased to increase the strength of the steel and formability. Reduce. Therefore, the Ni content is 0.1 to 1.0%.
B:Bは、オーステナイトからフェライトへの変態を抑制し、焼入れ性を高める元素であり、焼鈍時の冷却過程において低温変態相の生成を促進して高強度化に大きく寄与する。また、熱間圧延時に加工オーステナイトの再結晶を抑制し、未再結晶オーステナイトからのフェライト変態を促進して{113}<110>方位を発達させ、焼鈍後のヤング率を向上させる。こうした効果を得るためには、その量を0.0005%以上とする必要がある。一方、その量が0.0020%を超えると、焼鈍時の冷却過程でフェライトの生成を著しく抑制してヤング率を低下させる。したがって、B量は0.0005〜0.0020%とする。 B: B is an element that suppresses the transformation from austenite to ferrite and improves the hardenability, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to an increase in strength. It also suppresses recrystallization of processed austenite during hot rolling, promotes ferrite transformation from unrecrystallized austenite, develops the {113} <110> orientation, and improves Young's modulus after annealing. In order to obtain such an effect, the amount needs to be 0.0005% or more. On the other hand, if the amount exceeds 0.0020%, the formation of ferrite is remarkably suppressed during the cooling process during annealing, and the Young's modulus is lowered. Therefore, the B amount is 0.0005 to 0.0020%.
Cu:Cuは、焼入れ性を高める元素であり、焼鈍時の冷却過程において低温変態相の生成を促進して高強度化に大きく寄与する。この効果を得るためには、その量を0.1%以上とする必要がある。一方、その量が2.0%を超えると、熱間での延性を低下させて割れを誘起するとともに、焼入れ性の効果も飽和する。したがって、Cu量は0.1〜2.0%とする。 Cu: Cu is an element that enhances hardenability, and promotes the formation of a low-temperature transformation phase in the cooling process during annealing, and greatly contributes to high strength. In order to obtain this effect, the amount needs to be 0.1% or more. On the other hand, if the amount exceeds 2.0%, the hot ductility is lowered to induce cracking, and the effect of hardenability is saturated. Therefore, the Cu content is 0.1 to 2.0%.
2)ミクロ組織
自動車のセンターピラーなどのコラム状の構造部材には、伸びなどの成形性に優れることが必要である。そのため、フェライト相とマルテンサイトやベイナイトなどの低温変態相とからなる複合組織によって高強度化を図る必要がある。このとき、590MPa以上のTSを得るには、低温変態相の体積率を1〜60%にする必要がある。一方、低温変態相の体積率が60%を超えると過度に高強度化し、伸びなどの成形性が著しく低下する。
2) Microstructure Column-shaped structural members such as automobile center pillars must have excellent formability such as elongation. Therefore, it is necessary to increase the strength by a composite structure including a ferrite phase and a low-temperature transformation phase such as martensite or bainite. At this time, in order to obtain TS of 590 MPa or more, the volume ratio of the low temperature transformation phase needs to be 1 to 60%. On the other hand, when the volume ratio of the low temperature transformation phase exceeds 60%, the strength is excessively increased and the moldability such as elongation is remarkably lowered.
ここで、低温変態相の体積率は、鋼板の圧延方向に平行な板厚断面の板厚1/4の位置を光学顕微鏡あるいは走査型電子顕微鏡により倍率1000で観察し、100mm四方の領域に存在する低温変態相の占有面積率を画像処理によって求め、体積率とした。 Here, the volume ratio of the low-temperature transformation phase is present in a 100 mm square region by observing the position of the plate thickness 1/4 of the plate thickness cross section parallel to the rolling direction of the steel plate with an optical microscope or a scanning electron microscope at a magnification of 1000. The area ratio occupied by the low-temperature transformation phase to be obtained was determined by image processing and used as the volume ratio.
3)X
一般に、ある部材の剛性は、素材のヤング率Eと断面2次モーメントの積で表され、断面2次モーメントは素材の板厚tのλ乗に比例するので、比例定数をAとすると下記の式(2)のように表せる。
剛性=A×E×tλ ・・・(2)
自動車のセンターピラーなどのコラム状の構造部材ではλは1に近い値となり、通常の鋼板のEは200〜210GPaなので、今、軽量化のために板厚を10%減少させても剛性を一定に保つには、Eを11%向上させて230GPaにする必要がある。
3) X
In general, the rigidity of a certain member is expressed by the product of the Young's modulus E of the material and the moment of inertia of the section, and the moment of inertia of the section is proportional to the λth power of the thickness t of the material. It can be expressed as equation (2).
Rigidity = A × E × t λ (2)
In column-shaped structural members such as automobile center pillars, λ is close to 1, and the normal steel sheet E is 200 to 210 GPa. Therefore, even if the sheet thickness is reduced by 10% to reduce weight, the rigidity is constant. In order to maintain this, E must be improved by 11% to 230 GPa.
Eは、図1に示す引張試験で求まる応力-歪曲線における低歪域における直線部分の傾きであるが、板厚を減少させると素材にはより高い応力が負荷されるようになるので、こうしたEで剛性を評価することは十分でなく、素材に実際に負荷される応力近辺での応力-歪曲線の傾きで評価する必要がある。特に、低温変態相を有する高強度冷延鋼板では、可動転位が多いためと思われるが、応力-歪曲線の傾きが低歪域から低下してくるので、上述したように十分な剛性を確保できず、撓んだり、ねじれたりする問題が生じる。 E is the slope of the straight line portion in the low strain region of the stress-strain curve obtained by the tensile test shown in Fig. 1, but as the plate thickness is decreased, higher stress is applied to the material. It is not sufficient to evaluate the stiffness with E, and it is necessary to evaluate with the slope of the stress-strain curve near the stress actually applied to the material. In particular, high-strength cold-rolled steel sheets with a low-temperature transformation phase are thought to have many movable dislocations, but the stress-strain curve slope decreases from the low-strain range, ensuring sufficient rigidity as described above. Inability to do so causes problems such as bending and twisting.
そこで、本発明者らが、TSが590MPa以上の高強度冷延鋼板をコラム状の構造部材に適用した場合に、どの程度の応力が外部から負荷されるかを調査したところ200MPa程度の応力が負荷されることが明らかになった。したがって、TSの40%の応力(約240MPa)が負荷された状態、すなわち0.40×TS近辺の応力が負荷されたときの応力-歪曲線の傾きを230GPa以上にすれば、板厚を10%以上減少してもコラム状の構造部材の剛性を確保できることになる。 Therefore, the present inventors investigated how much stress is applied from the outside when a high-strength cold-rolled steel sheet having a TS of 590 MPa or more is applied to a columnar structural member. It became clear that it was loaded. Therefore, if 40% stress of TS (approx. 240 MPa) is applied, that is, if the stress-strain curve slope is set to 230 GPa or more when stress around 0.40 × TS is applied, the plate thickness will be 10% or more. Even if it decreases, the rigidity of the columnar structural member can be secured.
なお、応力-歪曲線の傾きを高精度で求めることは難しいので、図1に示すように、0.40×TSと0.45×TSの2点の傾きを上記の式(1)を用いて計算したXで評価した。すなわち、Xが230GPa以上であれば、コラム状の構造部材の剛性を確保できることになる。なお、Xが230GPa以上であれば、必然的にEは230GPa以上となる。また、0.60×TS近辺の傾きが230GPa以上であれば、より高い応力が負荷された場合でも剛性を確保できることになる。 Since it is difficult to determine the slope of the stress-strain curve with high accuracy, as shown in FIG. 1, the slopes of two points of 0.40 × TS and 0.45 × TS were calculated using the above equation (1). It was evaluated with. That is, if X is 230 GPa or more, the rigidity of the columnar structural member can be secured. If X is 230 GPa or more, E is inevitably 230 GPa or more. Further, if the slope near 0.60 × TS is 230 GPa or more, rigidity can be ensured even when higher stress is applied.
本発明の高強度冷延鋼板の表面には、電気めっき法あるいは溶融めっき法などにより、純亜鉛、亜鉛系合金、純Al、Al系合金などのめっき層を設けることができる。 On the surface of the high-strength cold-rolled steel sheet of the present invention, a plated layer of pure zinc, zinc-based alloy, pure Al, Al-based alloy or the like can be provided by electroplating or hot dipping.
4)製造方法
本発明の一例である製造方法における基本的な考え方は、Xを230GPa以上とするために、まずEを高めて230GPa以上を確保し、次に0.4×TS近辺の応力が負荷されたときの応力-歪曲線の傾き、すなわちXを向上させることにある。
4) Manufacturing method The basic idea in the manufacturing method which is an example of the present invention is that in order to set X to 230 GPa or more, first, E is increased to ensure 230 GPa or more, and then stress around 0.4 × TS is applied. It is to improve the slope of the stress-strain curve, that is, X.
4-1) 熱間圧延条件
転炉や電炉などで溶製後、鋳造されたスラブを、950℃以下における総圧下量を30%以上とし、Ar3変態点〜900℃の仕上温度で熱間圧延すると、熱間圧延直後に{112}<111>方位からなる未再結晶のオーステナイトが発達し、その後の冷却過程においてこの未再結晶オーステナイトが{113}<110>方位のフェライトに変態し、冷間圧延後の焼鈍時に{112}<110>方位のフェライトが生成してヤング率を向上させることができる。
4-1) Hot rolling conditions After the slab has been melted in a converter or electric furnace, the total reduction at 950 ° C or lower is 30% or higher, and the hot rolling is performed at a finishing temperature of Ar 3 transformation point to 900 ° C. When rolled, unrecrystallized austenite consisting of {112} <111> orientation develops immediately after hot rolling, and in the subsequent cooling process, this unrecrystallized austenite transforms into ferrite of {113} <110> orientation, During annealing after cold rolling, ferrite of {112} <110> orientation is generated, and Young's modulus can be improved.
4-2)巻取温度
熱間圧延後の鋼板を巻取るに当り、巻取温度が650℃を超えると、Nbの炭窒化物が粗大化し、焼鈍時の昇温過程においてフェライトの再結晶を抑制する効果が小さくなり、未再結晶フェライトからオーステナイトに変態させることが困難となる。その結果、その後の冷却過程で変態するフェライトの方位を制御することができず、ヤング率が大きく低下してしまう。したがって、巻取温度は650℃以下とする。なお、巻取温度は低くなり過ぎると形状不良となりやすいため、350℃以上とすることが好ましい。
4-2) Winding temperature When winding the steel sheet after hot rolling, if the winding temperature exceeds 650 ° C, the Nb carbonitrides will become coarse and ferrite recrystallization will occur during the heating process during annealing. The effect of suppressing becomes small, and it becomes difficult to transform from non-recrystallized ferrite to austenite. As a result, the orientation of ferrite that transforms in the subsequent cooling process cannot be controlled, and the Young's modulus is greatly reduced. Therefore, the coiling temperature is 650 ° C. or less. Note that if the coiling temperature is too low, the shape is likely to be defective, and therefore, the temperature is preferably 350 ° C. or higher.
4-3)圧下率
熱間圧延後の鋼板を冷間圧延することにより、熱延鋼板で発達した{113}<110>方位を{112}<110>方位に回転させることができ、その後の焼鈍時に{112}<110>方位の発達したフェライトを生成でき、ヤング率を高くすることができる。このような効果を得るには、冷間圧延時の圧下率を40%以上とする必要がある。
4-3) Reduction ratio By cold rolling the hot-rolled steel sheet, the {113} <110> orientation developed in the hot-rolled steel sheet can be rotated to the {112} <110> orientation. During annealing, ferrite with {112} <110> orientation can be generated, and Young's modulus can be increased. In order to obtain such an effect, the rolling reduction during cold rolling needs to be 40% or more.
4-4)焼鈍時の加熱条件
冷間圧延後の鋼板は焼鈍されるが、焼鈍の加熱時に{112}<110>方位をもつフェライトをオーステナイトへ変態させ、その後の冷却過程でオーステナイトから{112}<110>方位をもつフェライトに逆変態させてヤング率を向上させるために、加熱時に十分な量のフェライトをオーステナイトに変態させる必要がある。それには、加熱温度を(875-400×[C]-30×[Mn])℃以上とする必要がある。一方、加熱温度が高過ぎると、オーステナイトが粗大になり、逆変態したフェライトの{112}<110>方位への集積が低下する。そのため、加熱温度は(915-400×[C]-30×[Mn])℃以下とする必要がある。なお、ここで、(875-400×[C]-30×[Mn])〜(915-400×[C]-30×[Mn])℃の温度はAc3変態点近傍の温度域を表す。すなわち、本発明鋼では、Ac3変態点は(895-400×[C]-30×[Mn])(発明者らの求めた実験式)で表され、上記温度範囲は概ねAc3変態点±20℃の範囲を表す。加熱時間は、フェライトからオーステナイトへの変態を促進するために、1s以上必要である。一方、長過ぎるとオーステナイトが粗大になるため、加熱時間は600s以下とする必要がある。さらに、オーステナイトの粗大化を抑制してヤング率の向上を図ったり、組織を微細化して高強度化を図るには、加熱時間を300s以下とすることが好ましい。
4-4) Heating conditions during annealing Although the steel sheet after cold rolling is annealed, the ferrite with {112} <110> orientation is transformed to austenite during heating during annealing, and then the austenite is transformed into {112 } In order to reversely transform to ferrite with <110> orientation and improve Young's modulus, it is necessary to transform a sufficient amount of ferrite to austenite during heating. For this purpose, the heating temperature must be (875-400 × [C] -30 × [Mn]) ° C. or higher. On the other hand, if the heating temperature is too high, austenite becomes coarse, and the accumulation of reverse transformed ferrite in the {112} <110> orientation decreases. Therefore, the heating temperature needs to be (915-400 × [C] -30 × [Mn]) ° C. or less. Here, the temperature from (875-400 × [C] -30 × [Mn]) to (915-400 × [C] -30 × [Mn]) ° C. represents the temperature range near the Ac 3 transformation point. . That is, in the steel of the present invention, the Ac 3 transformation point is represented by (895-400 × [C] -30 × [Mn]) (empirical formula obtained by the inventors), and the above temperature range is approximately the Ac 3 transformation point. Represents a range of ± 20 ° C. The heating time is 1 s or more in order to promote the transformation from ferrite to austenite. On the other hand, if it is too long, austenite becomes coarse, so the heating time needs to be 600 s or less. Furthermore, in order to suppress the coarsening of austenite and improve the Young's modulus, or to refine the structure and increase the strength, the heating time is preferably set to 300 s or less.
4-5)焼鈍時の冷却条件
加熱後の鋼板は、冷却過程において、{112}<110>方位をもつフェライトを十分に生成させるために、700〜500℃の温度域に20s以上滞留させる必要がある。一方、長時間滞留させると炭化物が生成して低温変態相の生成が困難になり、強度が低下してしまう傾向があるため、滞留時間は200s以下とすることが好ましい。なお、滞留時間は、上記温度域で冷却速度を変えたり、恒温保持して制御できる。
4-5) Cooling conditions during annealing Heated steel sheets must be retained for at least 20 s in the temperature range of 700-500 ° C in order to sufficiently generate ferrite with the {112} <110> orientation during the cooling process. There is. On the other hand, if the residence time is long, carbides are generated and it is difficult to produce a low temperature transformation phase, and the strength tends to decrease. Therefore, the residence time is preferably 200 s or less. The residence time can be controlled by changing the cooling rate in the above temperature range or keeping the temperature constant.
4-6)焼鈍時の焼戻し条件
上記温度域に滞留させた鋼板は、さらに200〜500℃の温度域に60s以上滞留させて焼戻すか、あるいは上記温度域で滞留後200℃以下に冷却した後、再加熱して200〜500℃の温度域に60s以上滞留させて焼戻して、低温変態相の周辺に多量に存在する可動転位を減少させたり、あるいは可動転位に固溶C、Nのような侵入型元素を析出させて転位の移動を妨げると、上述したXを230GPa以上にすることができる。このとき、焼戻し温度が200℃未満では、その効果が得られず、一方、温度が高くなり過ぎ500℃を超えると、低温変態相が著しく軟化するので、200〜500℃の温度域で60s以上滞留させて焼戻す必要がある。好ましくは200〜450℃の温度域で60s以上滞留させる必要がある。なお、滞留時間が60s未満だと、上記効果を十分に得ることが困難になる。また、長時間滞留させると炭化物が粗大化し、低温変態相が著しく軟化して、強度が低下してしまうため滞留時間は1hr以下とすることが好ましい。なお、60s以上滞留は、加熱速度や冷却速度を調整したり、あるいは恒温保持で行える。
4-6) Tempering conditions during annealing The steel sheet retained in the above temperature range was further tempered by retaining for 60 s or more in the temperature range of 200 to 500 ° C, or cooled to 200 ° C or less after retention in the above temperature range. After that, reheat and stay in the temperature range of 200-500 ° C for 60s or more and temper it to reduce the mobile dislocations that exist in large amounts around the low-temperature transformation phase, or as a solid solution C, N in the mobile dislocations If the interstitial elements are precipitated to prevent dislocation movement, the above-mentioned X can be increased to 230 GPa or more. At this time, if the tempering temperature is less than 200 ° C., the effect cannot be obtained. On the other hand, if the temperature becomes too high and exceeds 500 ° C., the low temperature transformation phase is remarkably softened, so that the temperature range of 200 to 500 ° C. is over 60 s. It must be tempered and retained. Preferably, it is necessary to retain for 60 seconds or more in a temperature range of 200 to 450 ° C. If the residence time is less than 60 seconds, it is difficult to sufficiently obtain the above effect. Further, if the residence time is long, the carbide is coarsened, the low temperature transformation phase is remarkably softened, and the strength is lowered. Therefore, the residence time is preferably 1 hour or less. In addition, residence for 60 seconds or more can be performed by adjusting the heating rate and the cooling rate, or by keeping the temperature constant.
表1に示す成分の鋼Aを、実験室にて、真空溶解炉で溶製し、熱間圧延、冷間圧延、焼鈍を行って冷延鋼板を作製した。このとき、熱間圧延に先立つ加熱条件:1250℃で1時間、熱間圧延における950℃以下の総圧下率:40%、仕上温度:860℃、熱間圧延後の板厚:3.0mm、巻取条件:600℃で1時間保持後炉冷する巻取相当処理(巻取温度600℃)、冷間圧延の圧下率:50%、冷間圧延後の板厚:1.5mm、焼鈍時の加熱条件:820℃で180s、700〜500℃の平均冷却速度:8℃/s(滞留時間:25s)、500〜300℃の平均冷却速度:15℃/s、焼戻し条件:300℃で180s保持し、次いで200℃まで平均冷却速度10℃/sで冷却後空冷を基本条件とし、冷間圧延の圧下率、焼鈍の加熱温度、700〜500℃の滞留時間、焼戻しの温度と時間を変化させた。すなわち、変化させた条件以外は上記の条件である。また、700〜500℃の温度域を平均冷却速度8℃/sで冷却後、室温まで空冷し、再加熱して焼戻しの温度と時間を変化させた条件も加えた。 Steel A having the components shown in Table 1 was melted in a laboratory in a vacuum melting furnace and subjected to hot rolling, cold rolling, and annealing to produce a cold rolled steel sheet. At this time, heating conditions prior to hot rolling: 1 hour at 1250 ° C., total rolling reduction of 950 ° C. or lower in hot rolling: 40%, finishing temperature: 860 ° C., plate thickness after hot rolling: 3.0 mm, winding Rolling conditions: Holding at 600 ° C for 1 hour and then furnace cooling (winding temperature 600 ° C), cold rolling reduction: 50%, plate thickness after cold rolling: 1.5mm, heating during annealing Conditions: 180 s at 820 ° C, average cooling rate from 700 to 500 ° C: 8 ° C / s (residence time: 25 s), average cooling rate from 500 to 300 ° C: 15 ° C / s, tempering conditions: 180 s at 300 ° C Then, after cooling to 200 ° C at an average cooling rate of 10 ° C / s, air cooling was the basic condition, and the cold rolling reduction, annealing heating temperature, 700-500 ° C residence time, and tempering temperature and time were changed. . That is, the above conditions are the same except for the changed conditions. In addition, the temperature range of 700 to 500 ° C. was cooled at an average cooling rate of 8 ° C./s, then air-cooled to room temperature, and reheated to change the tempering temperature and time.
そして、作製した冷延鋼板から、圧延方向に対し直角な方向より10×60mmの試験片を切り出し、共振法によりヤング率Eを測定した。また、焼鈍後0.5%の調質圧延を施した冷延鋼板から、圧延方向に対し直角な方向よりJIS 5号引張試験片を切り出し、引張特性(TSと伸びEl)を測定した。さらに、調質圧延を施さない冷延鋼板から、圧延方向に対し直角な方向よりJIS 5号引張試験片を切り出し、引張試験により応力を負荷し、歪みゲージを用いて、0.40×TSと0.45×TSおよび0.60×TSと0.65×TSの応力負荷時の歪み量を測定し、上記式(1)からX(0.4TS)とX(0.6TS)を求めた。さらにまた、上述した方法で低温変態相の体積率を求めた。なお、低温変態相以外の相はフェライト相であった。 Then, a 10 × 60 mm test piece was cut out from the produced cold-rolled steel sheet in a direction perpendicular to the rolling direction, and Young's modulus E was measured by a resonance method. In addition, JIS No. 5 tensile specimens were cut from a direction perpendicular to the rolling direction from cold rolled steel sheets subjected to temper rolling at 0.5% after annealing, and tensile properties (TS and elongation El) were measured. Furthermore, from a cold rolled steel sheet not subjected to temper rolling, a JIS No. 5 tensile test piece was cut out from a direction perpendicular to the rolling direction, stress was applied by a tensile test, and 0.40 × TS and 0.45 × using a strain gauge. The amount of strain at the time of stress loading of TS and 0.60 × TS and 0.65 × TS was measured, and X (0.4TS) and X (0.6TS) were obtained from the above formula (1). Furthermore, the volume fraction of the low temperature transformation phase was determined by the method described above. The phases other than the low temperature transformation phase were ferrite phases.
基本条件で作製した冷延鋼板は、TS:790MPa、El:20%、E:245GPa、X:243GPaを有し、優れた強度-延性バランスと高い剛性を示す。 The cold-rolled steel sheet produced under the basic conditions has TS: 790 MPa, El: 20%, E: 245 GPa, X: 243 GPa, and exhibits an excellent strength-ductility balance and high rigidity.
表2に、冷間圧延の圧下率を変えて得られた結果を示す。本発明である40%以上の圧下率では、優れた強度-延性バランスを示すとともに、Xが230GPa以上となり高い剛性が得られる。 Table 2 shows the results obtained by changing the cold rolling reduction ratio. At a rolling reduction of 40% or more according to the present invention, an excellent balance between strength and ductility is exhibited, and X is 230 GPa or more and high rigidity is obtained.
表3に、焼鈍時の加熱温度を変えて得られた結果を示す。本発明である(875-400×[C]-30×[Mn])〜(915-400×[C]-30×[Mn])℃の加熱温度では、優れた強度-延性バランスを示すとともに、X(0.4TS)およびX(0.6TS)がともに230GPa以上となり高い剛性が得られる。 Table 3 shows the results obtained by changing the heating temperature during annealing. The heating temperature of (875-400 × [C] -30 × [Mn]) to (915-400 × [C] -30 × [Mn]) ° C. according to the present invention shows an excellent strength-ductility balance. , X (0.4TS) and X (0.6TS) are both 230 GPa or more, and high rigidity is obtained.
表4に、700〜500℃の滞留時間を変えて得られた結果を示す。なお、滞留時間は冷却速度を変化させることにより変えた。本発明である20s以上の滞留時間では、優れた強度-延性バランスを示すとともに、X(0.4TS)およびX(0.6TS)がともに230GPa以上となり高い剛性が得られる。 Table 4 shows the results obtained by changing the residence time at 700 to 500 ° C. The residence time was changed by changing the cooling rate. In the residence time of 20 s or more according to the present invention, an excellent balance between strength and ductility is exhibited, and both X (0.4 TS) and X (0.6 TS) are 230 GPa or more, and high rigidity is obtained.
表5に、再加熱せずにそのまま焼戻したときに焼戻し温度と時間を変えて得られた結果を示す。なお、500℃から焼戻し温度までは平均冷却速度15℃/sで冷却し、表に記載の焼戻し温度および時間で焼戻し後、200℃まで平均冷却速度10℃/sで冷却して空冷した。ここで、100℃、200℃で焼戻した場合は、焼戻し後空冷とした。本発明である焼戻し温度を200〜500℃とし、この温度域での滞留時間を60s以上とした場合では、優れた強度-延性バランスを示すとともに、X(0.4TS)およびX(0.6TS)がともに230GPa以上となり高い剛性が得られる。 Table 5 shows the results obtained by changing the tempering temperature and time when tempering without reheating. In addition, from 500 degreeC to the tempering temperature, it cooled with the average cooling rate of 15 degree-C / s, and after tempering with the tempering temperature and time as described in a table | surface, it cooled to 200 degreeC with the average cooling rate of 10 degree-C / s, and air-cooled. Here, when tempering at 100 ° C. and 200 ° C., air cooling was performed after tempering. When the tempering temperature of the present invention is 200 to 500 ° C. and the residence time in this temperature range is 60 s or more, an excellent balance between strength and ductility is exhibited, and X (0.4TS) and X (0.6 TS) are Both are over 230 GPa and high rigidity is obtained.
表6に、再加熱して焼戻したときに焼戻し温度と時間を変えて得られた結果を示す。なお、焼戻し温度までは平均10℃/sで加熱し、表に記載の焼戻し温度および時間で焼戻し後、200℃まで平均冷却速度10℃/sで冷却して空冷した。また、ここで、100℃、200℃で焼戻した場合は、焼戻し後空冷とした。本発明である焼戻し温度を200〜500℃とし、この温度域での滞留時間を60s以上とした場合では、優れた強度-延性バランスを示すとともに、X(0.4TS)およびX(0.6TS)がともに230GPa以上となり高い剛性が得られる。 Table 6 shows the results obtained by changing the tempering temperature and time when reheating and tempering. It was heated at an average temperature of 10 ° C./s until the tempering temperature, tempered at the tempering temperature and time shown in the table, then cooled to 200 ° C. at an average cooling rate of 10 ° C./s and air-cooled. Here, when tempering at 100 ° C. and 200 ° C., air cooling was performed after tempering. When the tempering temperature of the present invention is 200 to 500 ° C. and the residence time in this temperature range is 60 s or more, an excellent balance between strength and ductility is exhibited, and X (0.4TS) and X (0.6 TS) are Both are over 230 GPa and high rigidity is obtained.
表7に示す成分の鋼B〜Yを、実験室にて、真空溶解炉で溶製し、実施例1の基本条件で熱間圧延、冷間圧延、焼鈍を行って冷延鋼板1〜24を作製した。このとき、焼鈍時の加熱温度は、C、Mn量にしたがって変え、本発明範囲内とした。そして、実施例1と同様な調査を行った。なお、実施例1と同様、低温変態相以外の相はフェライト相であった。 Steels B to Y having the components shown in Table 7 were melted in a laboratory in a vacuum melting furnace and subjected to hot rolling, cold rolling, and annealing under the basic conditions of Example 1 to cold rolled steel plates 1 to 24. Was made. At this time, the heating temperature at the time of annealing was changed according to the amounts of C and Mn, and was within the scope of the present invention. Then, the same investigation as in Example 1 was performed. As in Example 1, the phases other than the low temperature transformation phase were ferrite phases.
結果を表8に示す。本発明である成分の条件を満たす鋼板1〜10、12〜15、18、21〜24では、優れた強度-延性バランスを示すとともに、X(0.4TS)が230GPa以上となり高い剛性が得られる。一方、[Nb]-(93/14)×[N]が本発明範囲外の鋼板11、Mn量が本発明範囲外の鋼板16、17、およびC量が本発明範囲外の鋼板19、20では、Eが低く、X(0.4TS)も230GPa未満で、剛性が低い。また、C量が著しく低い鋼板20では、TSも450MPaで高強度が得られない。 The results are shown in Table 8. In the steel plates 1 to 10, 12 to 15, 18, and 21 to 24 that satisfy the component conditions of the present invention, an excellent strength-ductility balance is exhibited, and X (0.4TS) is 230 GPa or more and high rigidity is obtained. On the other hand, [Nb] − (93/14) × [N] is a steel plate 11 outside the scope of the present invention, steel plates 16 and 17 whose Mn amount is outside the scope of the present invention, and steel plates 19 and 20 whose C amount is outside the scope of the present invention. Then, E is low, X (0.4TS) is less than 230GPa, and rigidity is low. Further, with the steel plate 20 having a remarkably low amount of C, TS is 450 MPa and high strength cannot be obtained.
Claims (6)
X=0.005×TS/(ε2-ε1)[GPa] ・・・(1)
ここで、[M]は元素Mの含有量[質量%]を表し、TSは鋼板の引張強度[MPa]、ε1は公称応力が0.40×TSのときの公称歪[%]、ε2は公称応力が0.45×TSのときの公称歪[%]を表す。 In mass%, C: 0.02 to 0.20%, Si: 1.5% or less, Mn: 1.0 to 3.5%, P: 0.05% or less, S: 0.01% or less, Al: 1.5% or less, N: 0.01% or less, Nb: Made of 0.01 to 0.40%, balance Fe and inevitable impurities, satisfying [Nb]-(93/14) × [N] ≧ 0.005, obtained by tempering by staying in the temperature range of 200-500 ° C for 60s or more It has a microstructure composed of a ferrite phase and a low temperature transformation phase, the volume ratio of the low temperature transformation phase is 1 to 60%, the tensile strength is 590 MPa or more, and X in the following formula (1) is 230 GPa or more. High strength cold-rolled steel sheet with high rigidity;
X = 0.005 × TS / (ε2-ε1) [GPa] (1)
Here, [M] represents the content of element M [% by mass], TS is the tensile strength [MPa] of the steel sheet, ε1 is the nominal strain [%] when the nominal stress is 0.40 × TS, and ε2 is the nominal stress Represents the nominal strain [%] when is 0.45 × TS.
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