JP5713135B1 - steel sheet - Google Patents

Abstract

The present invention provides a steel sheet that is low in manufacturing cost and high in productivity, and has excellent brittle crack propagation stopping performance of a base material and toughness of a weld heat affected zone during high heat input welding. In the case of having a predetermined chemical component, a carbon equivalent of 0.30 to 0.40 mass%, a SOLB of −0.0015 to +0.0015 mass%, and a plate thickness of 1/2 part, When the weighted average value DAVE of the diameter is 3.0 to 17.0 μm and the thickness is 1/2 part, the area ratio of the {100} plane that forms an angle within 15 ° with respect to the plane perpendicular to the rolling direction is 2.0. ˜20.0%, with a plate thickness of ¼ part, the total number density of TiN particles, MnS particles, and composite particles each having an equivalent circle diameter of 0.5 to 2.0 μm is 20 to 200 particles / mm 2, In the plate thickness 1/4 part, the number density of oxide particles having an equivalent circle diameter of 1 to 10 μm is 20 to 200 particles / mm 2, the plate thickness is 10 to 35 mm, and the yield stress is 300 to 500 MPa. [Selection figure] None

Description

  The present invention relates to a steel sheet that is excellent in the toughness of the heat affected zone and the brittle crack propagation stopping performance of the base metal when welding is performed.

For steel plates used in structures such as shipbuilding, architecture, tanks, offshore structures, and line pipes, the ability to suppress the occurrence of brittle cracks from welded joints in order to suppress brittle fracture of structures ( Hereinafter, joint toughness) is required. In addition, even if brittle fracture occurs at the welded joint location, the structural steel plate as described above is also required to stop the propagation of brittle cracks in the base metal (hereinafter referred to as arrestability). It is done. In particular, in recent years, oil fields have been actively developed in cold regions such as ice seas where reserves of energy resources such as oil and natural gas are large, so in the technical field related to ships and offshore structures, for example, −60 Even in an extremely low temperature environment of ° C., the above-described fracture resistance is required. Such members often use thick steel plates having a yield stress of 300 to 500 MPa and a plate thickness of 10 to 35 mm. In addition, steel sheets used in the technical field related to ships are required to have tensile strength, yield stress and the like in addition to fracture resistance. For example, a steel sheet for a spherical LNG tank exterior used in an LNG ship having a structure in which a hull structure and a spherical LNG tank exterior are integrated in practical use in recent years has a tensile strength of 450 to 620 MPa, preferably 490 to 620 MPa. (TS), yield stress (YP) of 300-500 MPa, preferably 355-500 MPa, and plate thickness of 10-35 mm, preferably 15-30 mm. Furthermore, the spherical LNG tank exterior steel plate must also have a base metal arrestability such that the arrest toughness value Kca at −60 ° C. is 4000 N / mm ^1.5 or more. In addition, in a weld heat affected zone (HAZ) of a welded joint obtained by welding a spherical LNG tank exterior steel plate under very large welding heat input conditions (for example, 50 to 200 kJ / cm), high low temperature toughness (HAZ) For example, it is necessary to obtain a low temperature toughness in which the average value of Charpy absorbed energy at −60 ° C. is 100 J or more.
However, the above-mentioned fracture resistance generally tends to be significantly reduced at extremely low temperatures. Furthermore, in order to improve welding construction efficiency and reduce costs, it is required to apply large heat input welding that can be welded in one pass. In this case, the joint toughness is likely to cause coarsening of the structure, so that descend. For this reason, in the said thick steel plate, the technique which improves the joint toughness and arrestability in -60 degreeC is desired. In addition, the welding heat input in 1-pass large heat input welding changes with plate | board thickness, and is 50-200 kJ / cm in the range whose plate | board thickness is 10-35 mm.

  As a method for improving the joint toughness, for example, a method for controlling the crystal grain size in a welding heat affected zone (hereinafter referred to as HAZ), a method for controlling an embrittlement second phase, and Ni which is an element imparting high toughness to steel. Methods of adding are known.

  As a method for controlling the crystal grain size, there is a method for suppressing the austenite grain coarsening in the heating process of welding by dispersing a large amount of fine pinning particles in the steel (hereinafter referred to as pinning technology). Illustrated. In addition, as another method for controlling the crystal grain size, a method of promoting intragranular transformation in the cooling process of welding by dispersing particles that become the core of ferrite transformation in steel, and subdividing the grain ( Hereinafter, an intragranular transformation technique) is also exemplified.

  As the pinning technique, techniques described in Patent Documents 1 to 10 are exemplified.

  In the technique described in Patent Document 1, TiN is dissolved in 0.004% or more by reheating, and TiN is finely precipitated and dispersed in the subsequent cooling process to refine the HAZ structure and increase joint toughness. ing.

In the technique described in Patent Document 2, the value obtained by dividing the Ti content by the N content is 1.0 to 6.0, so that TiN having a particle diameter of 0.01 to 0.1 μm is 5 × 10 5. ^{5 to} 5 × 10 ⁶ pieces / mm ² are present to thereby refine the HAZ structure and improve the joint CTOD characteristics.

In the techniques described in Patent Documents 3 to 6, in the steelmaking process, the amount of dissolved oxygen in the molten steel is adjusted to a predetermined value using an equilibrium reaction with Si, and then Ti, which is a deoxidizing element, and Al are further added. By adding in order, the Ti-Al composite oxide is uniformly and finely dispersed. Thereby, in the techniques described in Patent Documents 3 to 6, the HAZ structure is refined and the joint toughness is improved. This Ti—Al composite oxide has a Ti composition ratio of 5% or more and an Al composition ratio of 95% or less, a particle diameter of 0.01 to 1.0 μm, and a particle number of 5 × 10 ³ to 1 × 10. Patent Document 6 describes that the number is ⁵ / mm ² .

In the technique described in Patent Document 7, 0.0010% or less of Mg is added after deoxidizing molten steel with Ti and Al. Thereby, a composite oxide mainly composed of Al—Ti—Mg having a particle diameter of 0.01 to 1.0 μm and a number of particles of 1 × 10 ⁴ to 2 × 10 ⁵ particles / mm ² is uniformly fine. Disperse to refine the HAZ structure and increase joint toughness.

In the techniques described in Patent Documents 8 and 9, oxide particles made of Ca, Al, and O elements having an equivalent circle diameter of 0.005 to 2.0 μm are dispersed in steel. This refines the HAZ structure and improves joint toughness. The oxide particles described in Patent Document 8 have a number density of 100 to 5000 / mm ² , and the chemical composition excluding O is Ca: 5 mass% or more and Al: 5 mass% or more. The oxide particles described in Patent Document 9 have a number density of 100 to 3000 / mm ² , and the chemical composition excluding O is Ca: 3 mass% or more and Al: 1 mass% or more.

In the technique described in Patent Document 10, the amount of dissolved oxygen in the molten steel is adjusted to 0.0010 to 0.0050% in the deoxidation step, first Ti and then Al are added to deoxidize, and further Ca, Mg And one or more elements of REM are added. This disperses 100 oxides / mm ² or more of an oxide having an equivalent circle diameter of 0.005 to 0.5 μm, refines the HAZ structure, and improves joint toughness.

  As the intragranular transformation technique, techniques described in Patent Documents 11 to 17 are exemplified.

  In the technique described in Patent Document 11, the HAZ structure is refined and the joint toughness is improved by generating intragranular ferrite with VN as a nucleus in the cooling process after welding.

  In the technique described in Patent Literature 12, TiN and / or BN particles are finely dispersed in the vicinity of the bond, and a metal structure having an area fraction of ferrite having a particle size of 50 μm or less is 60% or more is obtained. Improve toughness.

  In the technique described in Patent Document 13, B is diffused from the weld metal into the HAZ, BN is precipitated in the HAZ, and fine ferrite is generated using the BN as a nucleus to improve joint toughness.

  In the technique described in Patent Document 14, the chemical composition of steel is a chemical composition in which the transformation point in a CCT curve with a cooling rate of 2 ° C./s or less is 670 ° C. or more, and the content ratio of N—Ti—B. To control. Thereby, the intragranular ferrite which made BN the nucleus is produced | generated, HAZ structure | tissue is refined | miniaturized and joint toughness is improved.

  In the technique described in Patent Document 15, the amount of dissolved oxygen at the time of Ca addition is adjusted to 0.0010 to 0.0030% in the melting step, and the addition amounts of Ca, S, and O are appropriately controlled. . Thereby, the composite sulfide in which MnS is precipitated on CaS is generated in the steel, and this composite sulfide functions as a ferrite transformation nucleus, refines the HAZ structure, and improves joint toughness.

  In the technique described in Patent Document 16, a composite inclusion in which BN is precipitated in Ti oxide is generated in steel. By using this composite inclusion as an intragranular transformation nucleus, a fine HAZ structure having a grain boundary ferrite fraction of 5% or less and an acicular ferrite size of 10 μm or less in equivalent circle diameter is formed, and joint toughness is improved.

  In the technique described in Patent Document 17, Al is added after deoxidizing molten steel using Ti. By using the Ti—Al composite oxide thus formed as a nucleus and further depositing TiN, MnS, and B-based precipitates as ferrite transformation nuclei, the HAZ structure is refined and joint toughness is improved.

  Moreover, there exists a technique described in patent documents 18-20 as a method of controlling an embrittlement 2nd phase in order to improve joint toughness.

  In the technique described in Patent Document 18, the chemical component is C: less than 0.03%, Mn: 0.6 to 1.2%, Ni: 1.0 to 2.3%, and Ni ≦ A weld joint is prepared using steel satisfying −2 × Mn + 4.0. Thereby, joint toughness is improved by suppressing the production | generation of island-like martensite.

  In the technique described in Patent Document 19, Mn is positively added to steel, and the content of P which is an impurity element is reduced to 0.008% by mass or less. As a result, the island-like untransformed austenite generated during cooling after high heat input welding is promoted to decompose and become cementite, and the area ratio of island-like martensite is suppressed to 1% or less, thereby reducing the joint. Improve toughness.

In the technique described in Patent Document 20, appropriate amounts of Mn, Ni and Cr are added to steel, and the C content is further reduced. Thereby, joint toughness is improved by suppressing the average area of the island-like martensite formed in the part heated to a two-phase region at the time of multilayer welding to 3 μm ² or less.

  Moreover, as a method of adding Ni which is a high toughness element in order to improve the toughness of the welded joint, techniques described in Patent Documents 21 and 22 are exemplified.

  In the technique described in Patent Document 21, the Ni content is set to 4.0 to 7.5%, and the martensite transformation start temperature (Ms) is set to 370 ° C. or lower. Thereby, the lath width which is a structural unit of a martensitic structure is refined, and joint toughness is improved.

  In the technique described in Patent Document 22, the Ni content is set to 4.0 to 6.0%, and the carbon equivalent represented by the predetermined formula is set to less than 0.40%. Thereby, the hardness of HAZ is reduced and joint toughness is improved.

  As a method for improving the arrestability, for example, a method of controlling the crystal grain size, a method of controlling the embrittlement second phase, a method of controlling the texture, and a method of adding Ni which is a tough element are known. .

  As a method of controlling the crystal grain size in order to improve arrestability, techniques described in Patent Documents 23 to 25 are exemplified.

  In the technique described in Patent Document 23, the arrestability is improved by using ferrite as a parent phase and making the ferrite finer. In order to obtain such fine-grained ferrite, a steel slab such that, during the production of the steel sheet, 1/8 or more of the steel slab thickness (slab thickness or slab thickness) becomes Ar3 point or less along the thickness direction. It is necessary to recrystallize the ferrite by cooling the steel, rolling in a cryogenic region, and then reheating to a temperature exceeding the Ac3 point.

  In the techniques described in Patent Documents 24 and 25, ferrite is used as a parent phase, the surface layer portion of the steel slab is once cooled to Ar1 or less, and then rolled until the surface layer portion is reheated. A steel plate having ferrite recrystallized grains is obtained.

  Moreover, the technique described in patent document 26 is illustrated as a method of controlling an embrittlement 2nd phase in order to improve arrestability.

  In the technique described in Patent Document 26, a fine embrittled second phase (for example, martensite) is dispersed in ferrite as a parent phase. As a result, a microcrack is generated in the brittle second phase at the tip of the brittle crack, and the stress state at the tip of the crack is relaxed.

  Moreover, there exists a technique described in patent documents 27-39 as a method of controlling a texture in order to improve arrestability.

  In the techniques described in Patent Documents 27 to 39, the X-ray plane intensity ratio as a texture is controlled by, for example, the surface thickness part, the plate thickness 1/4 part, and the plate thickness 1/2 part. Change the direction of crack propagation and improve arrestability.

  Furthermore, as a method of controlling both the crystal grain size and the texture in order to improve arrestability, techniques described in Patent Documents 40 to 42 are exemplified.

  In the technique described in Patent Document 40, the ferrite fraction at 1/2 part of the plate thickness is set to 80% or more, and the crystal grain size and the X-ray plane intensity ratio are controlled, whereby 45 ° with respect to the rolling direction. Improve the direction arrestability.

  In the technique described in Patent Document 41, the arrestability is improved by controlling the crystal grain size of the surface layer and ½ part of the plate thickness and the texture strength ratio measured by X-ray.

  In the technique described in Patent Document 42, the arrestability is improved by controlling the crystal grain size of the surface layer and half the plate thickness and the area ratio of the {100} plane perpendicular to the external stress.

  Moreover, as a method of adding Ni which is a high toughness element in order to improve arrestability, there is a technique described in Patent Document 43. In the technique described in Patent Document 43, the average value of the circle-equivalent grain sizes of the crystal grains surrounded by the large-angle grain boundaries of 15 ° or more with Ni content of more than 5.0 to less than 10.0% is the plate thickness. Arrestability is improved by setting it to 5.5 μm or less at the ¼ position.

Japanese Patent Publication No. 55-26164 Japanese Unexamined Patent Publication No. 2001-164333 Japanese Patent Laid-Open No. 6-293936 Japanese Unexamined Patent Publication No. 7-242985 Japanese Unexamined Patent Publication No. 7-242925 Japanese Unexamined Patent Publication No. 9-003597 Japanese Unexamined Patent Publication No. 9-003598 Japanese Unexamined Patent Publication No. 2001-342537 Japanese Laid-Open Patent Publication No. 2003-313629 Japanese Unexamined Patent Publication No. 2005-320624 Japanese Patent Laid-Open No. 5-186848 Japanese Unexamined Patent Publication No. 9-324238 Japanese Unexamined Patent Publication No. 2003-212268 Japanese Unexamined Patent Publication No. 2005-008967 Japanese Unexamined Patent Publication No. 2005-068519 Japanese Unexamined Patent Publication No. 2007-2777681 Japanese Laid-Open Patent Publication No. 5-239528 Japanese Laid-Open Patent Publication No. 2002-060891 Japanese Unexamined Patent Publication No. 2011-032519 Japanese Unexamined Patent Publication No. 2012-172243 Japanese Unexamined Patent Publication No. 6-136383 Japanese Patent Laid-Open No. 6-192729 Japanese Unexamined Patent Publication No. 61-235534 Japanese Unexamined Patent Publication No. 2003-221619 Japanese Patent Laid-Open No. 5-148542 Japanese Patent Publication No. 59-47323 Japanese Unexamined Patent Publication No. 2008-045174 Japanese Unexamined Patent Publication No. 2008-069380 Japanese Unexamined Patent Publication No. 2008-111165 Japanese Unexamined Patent Publication No. 2008-111166 Japanese Unexamined Patent Publication No. 2008-169467 Japanese Unexamined Patent Publication No. 2008-169468 Japanese Laid-Open Patent Publication No. 2008-214652 Japanese Unexamined Patent Publication No. 2008-214653 Japanese Unexamined Patent Publication No. 2008-214654 Japanese Unexamined Patent Publication No. 2009-185343 Japanese Unexamined Patent Publication No. 2009-221585 Japanese Unexamined Patent Publication No. 2009-235458 Japanese Unexamined Patent Publication No. 2010-0478805 Japanese Unexamined Patent Publication No. 2011-068952 Japanese Unexamined Patent Publication No. 2011-214116 Japanese Unexamined Patent Publication No. 2007-302993 Japanese Unexamined Patent Publication No. 2011-211988

  In the techniques described in Patent Documents 1 to 10, it is necessary to disperse a large amount of pinning particles in steel. At an extremely low temperature of −60 ° C., the pinned particles are likely to be the starting point of brittle fracture, so the toughness of a steel welded joint having such characteristics is lowered. Therefore, it is difficult to stably obtain high joint toughness from the techniques described in Patent Documents 1 to 10. In particular, as in the techniques described in Patent Documents 4 to 10, when oxide particles are used as pinning particles, thermally stable fine particles can be obtained, but oxide particles are generated at a high temperature. Therefore, it is inevitable that coarse particles are also produced. The coarse particles serve as a starting point for fracture and reduce joint toughness. As in the techniques described in Patent Documents 8 to 10, this is a case where high-strength steel with a large plate thickness is welded with a large welding heat input exceeding 200 kJ / cm, and only toughness at a relatively high temperature is evaluated. In this case, the techniques described in Patent Documents 1 to 10 are effective techniques, but it is extremely difficult to improve the joint toughness at −60 ° C. targeted by the present invention by these techniques. In addition, the pinning technique can refine the austenite grains when heated in the welding process, but cannot refine the grains. Since the final particle size of the metal structure of the welded part depends on the austenite particle size, there is a limit to refinement by the pinning technique. Therefore, the pinning technique is inferior to the intragranular transformation technique in terms of the HAZ microstructure refinement effect.

  In the technique described in Patent Document 11, VN that precipitates during cooling after the welding process is used as an intragranular transformation nucleus. However, it is difficult to use VN under the welding conditions (plate thickness, welding heat input) targeted in the present invention. Since VN has a small driving force for precipitation, welding is performed under welding conditions (conditions where the cooling rate is extremely low) such that the welding heat input is 1000 kJ / cm as described in the example of Patent Document 11. Can only be deposited. Furthermore, VN alone has only a ferrite nucleation ability that is inferior to other particles. Therefore, the technique described in Patent Document 11 cannot be applied under the conditions required for the present invention.

  In the technique described in Patent Document 12, since the Al content must be controlled to an extremely low level of 0.005% or less, deoxidation becomes insufficient, and a large amount of coarse oxide is easily generated. This coarse oxide serves as a starting point of fracture and significantly reduces toughness. In addition, since the ferrite nucleation ability of the TiN—BN composite particles is small, it is difficult to obtain a sufficient microstructure using the TiN—BN composite particles.

  In the technique described in Patent Document 13, it is necessary to diffuse B from the weld metal to the HAZ. However, since the diffusion is affected by various conditions, the amount of precipitation of BN as a ferrite transformation nucleus is stably increased. It is difficult to keep. The first of the various conditions that affect diffusion is welding heat input. When the welding heat input is low, and thus the high temperature residence time is short or the cooling rate is high, B cannot be sufficiently diffused. Therefore, the technique described in Patent Document 13 cannot be applied to welding performed under welding conditions with low welding heat input. The second of the various conditions that affect diffusion is the Nb and Ti content. Since the diffusion of B occurs through the vacancies, when an element such as Nb or Ti that has a strong binding force to the vacancies is added, the diffusion rate of B decreases and the B is sufficiently diffused. I can't let you. Therefore, the technique described in Patent Document 13 cannot be applied to steel that needs to contain Nb and Ti.

  In the technique described in Patent Document 14, BN is dispersed in a metal structure, intragranular ferrite having BN as a nucleus is generated, the HAZ structure is refined, and joint toughness is improved. However, ferrite cannot be sufficiently generated only by dispersion of BN. Furthermore, since the dispersion state of BN is not limited in Patent Document 14, it is difficult to stably improve joint toughness by HAZ structure control from the contents described in Patent Document 14. In addition, when a welded joint is prepared using an example having a large plate thickness and high strength disclosed in Patent Document 14, in view of its chemical composition, HAZ becomes too hard and is an embrittled phase. MA (Martensite-Austenite constituent) is easily generated, and it is considered impossible to improve joint toughness at -60 ° C. Furthermore, in the Example disclosed in Patent Document 14, the welding heat input condition is set to 340 to 530 kJ / cm. However, when welding is performed under this welding heat input condition, the joint toughness of −60 ° C. is improved. It is very difficult to refine the HAZ structure.

  In the technique described in Patent Document 15, the HAZ structure is refined by using a composite sulfide in which MnS is precipitated on CaS as a ferrite transformation nucleus. However, since the ferrite transformation ability of the composite sulfide is low, the effect of the composite sulfide cannot be sufficiently exhibited when the welding heat input is small. For this reason, the welding heat input is 400 kJ / cm (this welding heat input is considered to be relatively high in this technical field), and Patent Document 15 targets welding conditions with a low cooling rate. Under such welding conditions, ferrite is easily generated. Therefore, the technique of Patent Document 15 cannot be applied to other welding conditions. Similarly to Patent Document 14 described above, when a welded joint is created using an embodiment disclosed in Patent Document 15 with high strength, HAZ becomes too hard and the embrittlement phase is present in the chemical component. Since it becomes easy to produce MA, it is considered extremely difficult to improve the joint toughness at −60 ° C.

  In the technique described in Patent Document 16, the Al content must be controlled to an extremely low level of less than 0.005%. Thereby, in the steel of patent document 16, deoxidation becomes inadequate and it becomes easy to produce | generate a large amount of coarse oxides. Since this coarse oxide becomes a starting point of fracture, toughness is remarkably lowered. In addition, since it is difficult to disperse a large amount of Ti oxide, in Patent Document 16 in which Ti oxide is used for refining the HAZ structure, there are few ferrite nucleation sites, and the HAZ structure cannot be sufficiently refined. . In the case of multi-layer welding with low welding heat input, which is targeted by Patent Document 16, it is considered that the technique disclosed in Patent Document 16 can also exert a miniaturization effect. Cannot apply the technique disclosed in Patent Document 16.

  In the technique described in Patent Document 17, an oxide is used as a ferrite transformation nucleus. However, as described above, in this case, thermally stable fine particles can be obtained, but oxide particles have the property of being generated at high temperatures, so if sufficient oxide particles are to be obtained, coarse particles are generated. Is inevitable. The coarse particles serve as a starting point for fracture, and joint toughness decreases. In addition, since the oxide particles cannot be dispersed at high density, it is difficult to obtain a fine HAZ structure from the technique described in Patent Document 17.

  Since the technique described in Patent Document 18 needs to control the C content to less than 0.03%, which is a very low level, and to add 1.0 to 2.0% Ni, the alloy cost is extremely high. Has the problem of high. Furthermore, since Ni frequently causes slab cracking during casting due to high temperature embrittlement, the technique described in Patent Document 18 has a problem that slab refining costs also increase significantly. Moreover, even if the control of refinement of the HAZ structure is not performed and only the island martensite is controlled, it is difficult to improve the joint toughness at −60 ° C.

  The techniques described in Patent Documents 19 and 20 also suppress the generation of island martensite in the HAZ, similar to Patent Document 18. However, the joint toughness at −60 ° C. cannot be sufficiently improved unless microstructure refinement control is performed. Patent Document 18 targets a welded joint obtained by performing multi-layer welding with a heat input of 300 kJ / cm or more on steel having a yield stress of 460 MPa or more, and Patent Document 19 discloses a steel having a yield stress of 630 MPa or more. It is intended for welded joints obtained by performing multi-layer welding with a welding heat input of 50 kJ / cm. Under these conditions, the metal structure of the present invention cannot be achieved. Therefore, it is extremely difficult to improve joint toughness at −60 ° C. using the techniques described in Patent Documents 19 and 20.

  In the techniques described in Patent Documents 21 and 22, it is necessary to add a large amount of Ni of 4.0% or more. Since the cost is extremely high, the techniques described in Patent Documents 21 and 22 cannot be employed for industrial use. Further, in the techniques described in Patent Documents 21 and 22, since it is necessary to obtain martensite, the welding heat input must be reduced to about 20 kJ / cm. Even if welding is performed under the welding heat input targeted by the present invention, it is difficult to improve joint toughness using the techniques described in Patent Documents 21 and 22, because the above-described metal structure cannot be obtained. is there.

  In the techniques described in Patent Documents 23 to 25, it is necessary to manufacture a steel sheet through cooling, rolling, and reheating processes in order to make the ferrite ultrafine by utilizing recrystallization of ferrite in the front and back layer portions of the steel sheet. is there. This technique complicates the manufacturing process, so that it is extremely difficult to obtain a steel plate having a stable material, and further low temperature rolling is required, so that productivity is low. Further, in such a manufacturing process, a shape defect due to non-uniform cooling of the steel sheet surface is likely to occur in the steel sheet. When a shape defect occurs, a large cost is required to correct the shape.

  In the technique described in Patent Document 26, since martensite is dispersed in ferrite, the characteristics for suppressing the occurrence of brittle cracks are significantly deteriorated. Moreover, it is difficult to improve the arrestability at −60 ° C. only by controlling the embrittlement second phase.

  Further, in the techniques described in Patent Documents 27 to 39, control of the crystal grain size, which is the most effective factor for improving arrestability, is not performed. That is, the arrestability at −60 ° C. cannot be dramatically improved only by controlling the texture. Further, the X-ray plane intensity ratio represents a local texture, and does not indicate the characteristics of the entire steel sheet. In a steel sheet whose texture is controlled only by controlling the X-ray surface intensity ratio, there may be a large variation in texture. Further, these techniques are intended for steel having a large thickness and high strength, but the arrestability of the steel sheet having the thickness and strength targeted by the present invention is improved, and hot rolling of the steel sheet is performed. It is not a technology that can sometimes provide high productivity. In the first place, the techniques of Patent Documents 27 to 30 and 33 are techniques for increasing the arrestability in the plate thickness direction, and are not techniques for improving the arrestability in a direction perpendicular to or parallel to the rolling direction as in the present invention. The present invention is not applicable to steel.

  Patent Documents 40 and 41 disclose techniques for obtaining high arrestability by controlling the crystal grain size and texture. However, the techniques described in Patent Documents 40 and 41 are intended for extra-thick steel plates having a thickness of 50 mm or more. In the production of a steel sheet having a thickness intended by the present invention, the rolling temperature and the cooling rate are different from the conditions described in Patent Documents 40 and 41. Therefore, it is difficult to obtain a steel sheet having the same crystal grain size and texture by the techniques of Patent Documents 40 and 41. Moreover, since the X-ray surface intensity ratio used in Patent Documents 40 and 41 for texture control represents a local texture, there is a large variation and a factor that improves arrestability. Not suitable for. In the techniques described in Patent Documents 40 and 41, rolling at a low temperature is indispensable in order to form a desired texture. This significantly reduces productivity. Further, Patent Document 40 is a technique for improving the arrestability at an angle of 45 ° with respect to the rolling direction, and a technique relating to improvement of arrestability in a direction perpendicular to or parallel to the rolling direction as in the present invention. is not.

  The technique described in Patent Document 42 controls the crystal grain size and texture in the surface layer and ½ part of the plate thickness, and is effective in improving the arrestability of a steel plate having a large plate thickness. However, when the plate thickness targeted by the present invention is 10 to 35 mm, the rolling and cooling conditions during production are different, so the crystal grain size and texture in the surface layer and 1/2 part of the plate thickness are clearly defined. It is difficult to separate and control and obtain a similar tissue.

  The technique described in Patent Document 43 requires a large amount of Ni addition exceeding 5.0% and is extremely expensive, which is not preferable for industrial use. Moreover, since Ni is added in a large amount, the strength of the steel sheet described in Patent Document 43 is high. In 1-pass high heat input welding, since the HAZ becomes hard and the joint toughness decreases, it is extremely difficult to obtain a steel sheet that has both joint toughness and arrestability by the technique described in Patent Document 43. .

The present invention has been made in consideration of the above-mentioned circumstances, and the purpose thereof is low manufacturing cost and high productivity, and the brittle crack propagation stopping performance of the base metal and welding during high heat input welding. It is providing the steel plate excellent in the toughness of a heat affected zone. Specifically, the yield stress of the base material is 300 to 500 MPa, the average Charpy absorbed energy at −80 ° C. of the base material is 100 J or more, and the arrestability (Kca) of the base material at −60 ° C. is 4000 N. / Mm ^1.5 or more, and has good toughness (for example, toughness with an average Charpy absorbed energy at −60 ° C. of 100 J or more) when welding with a heat input of 50 to 200 kJ / cm is performed. It is an object of the present invention to obtain a steel plate from which a weld heat affected zone can be obtained.

  The gist of the present invention for solving the above problems is as follows.

(1) As for the steel plate which concerns on 1 aspect of this invention, a chemical composition is the mass%, C: 0.040-0.090%, Si: 0.01-0.20%, Mn: 1.30-1 80%, P: 0.020% or less, S: 0.001 to 0.010%, Al: 0.005 to 0.100%, Nb: 0.003 to 0.030%, Ti: 0.003 -0.030%, B: 0.0003-0.0040%, N: 0.0020-0.0080%, O: 0.0005-0.0040%, Cu: 0-1.00%, Ni: 0 to 1.00%, Cr: 0 to 1.00%, Mo: 0 to 0.500%, V: 0 to 0.100%, Ca: 0 to 0.0050%, Mg: 0 to 0.0050 %, REM: 0 to 0.0050%, and the balance: iron and impurities, and the carbon equivalent CE defined by the formula A is 0.30 to 0 40 mass%, SOLB defined by the formula B is −0.0015 to +0.0015 mass%, and the metal structure is a mixed structure including ferrite and bainite, or a mixture including the ferrite, pearlite, and the bainite. a tissue area ratio 50-90% of the ferrite area ratio is 10% to 50% of the previous SL bainite, and the area ratio of 0 to 10% of the pearlite area ratio of MA of the metal structure Is 0 to 5%, and in the plate thickness 1/2 part, the boundary between adjacent crystals whose crystal orientation difference is 15 ° or more is defined as a crystal grain boundary, and is surrounded by the crystal grain boundary. Defined as a crystal grain, and when the number of divisions N is set to an integer of 10 or more, the weighted average value D _AVE of the grain size of the crystal grain defined by the E formula is 3.0 to 17.0 μm. And before In the plate thickness 1/2 part, the area ratio of the {100} plane that forms an angle within 15 ° with respect to the plane perpendicular to the rolling direction is 2.0 to 20.0%. Particles containing 1% by mass or more of Ti, less than 1% by mass of O and 1% by mass or more of N are defined as TiN particles, and 1% by mass or more of Mn, 1% by mass or more of S and less than 1% by mass. When the particle containing O is defined as MnS particle, and the particle satisfying both the definition of TiN particle and the definition of MnS particle is defined as a composite particle, an equivalent circle diameter of 0.5 to 2.0 μm is obtained. The total number density of the TiN particles, the MnS particles, and the composite particles that each have is 20 to 200 particles / mm ² , and particles containing 1 mass% or more of O are oxidized at the thickness of 1/4 part. The number of oxide particles having an equivalent circle diameter of 1 to 10 μm when defined as product particles Degree is 20 to 200 pieces / mm ^2, thickness is 10～35Mm, yield stress is 300～500MPa.
CE = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 (A)
SOLB = 0.226 × Ti + B−0.772 × N (B)
p = (D _MAX −D _MIN ) / N (C)
D _k = D _MIN + p × (k−1 / 2) (D)
D _AVE = (Σ [k = 1, N] (D _k × S _k )) / (Σ [k = 1, N] S _k ) (E)
The element symbol described in the formula indicates the content of each component in the steel sheet in unit mass%, D _MAX indicates the particle size of the largest crystal grain in unit μm, and D _MIN is The grain size of the smallest crystal grain is expressed in μm, k is an integer of 1 or more and N or less, and S _k is (D _MIN + p × (k−1)) μm or more (D _MIN + p × k) The total area ratio of the crystal grains having a circle-equivalent diameter of less than μm is shown in unit%.

(2) In the steel plate according to (1), the chemical composition may be% by mass and Si: 0.01 to 0.10%.

(3) As for the steel plate as described in said (1) or (2), the said chemical composition may be mass% and Al: 0.015-0.060%.

(4) In the steel sheet according to any one of (1) to (3), the chemical composition may be mass% and Mo may be 0 to 0.010%.

(5) In the steel sheet according to any one of (1) to (4), the chemical composition may be mass% and Ti: 0.005 to 0.018%.

(6) In the steel plate according to any one of (1) to (5), the chemical composition may be% by mass and B: 0.0005 to 0.0020%.

(7) In the steel sheet according to any one of (1) to (6), the chemical composition may be% by mass and N: 0.0025 to 0.0060%.

(8) As for the steel plate as described in any one of said (1)-(7), the said chemical composition may be mass% and O: 0.0010-0.0030%.

(9) In the steel sheet according to any one of (1) to (8), the SOLB may be -0.0010 to + 0.0005%.

(10) In the steel plate according to any one of (1) to (9), the weighted average value D _{AVE of the} crystal grains may be 3.0 to 13.0 μm.

(11) In the steel sheet according to any one of (1) to (10), the TiN particles, the MnS particles, and the composite particles each having an equivalent circle diameter of 0.5 to 2.0 μm. The total number density may be 50 to 140 pieces / mm ² .

(12) In the steel sheet according to any one of (1) to (11), the number density of the oxide particles having an equivalent circle diameter of 1 to 10 μm is 20 to 150 particles / mm ^2. Also good.

(13) In the steel sheet according to any one of (1) to (12), the chemical composition is mass%, Cu: 0.10 to 1.00%, Ni: 0.10 to 1. 00%, Cr: 0 to 0.10%, Mo: 0 to 0.100%, V: 0 to 0.005%.

  When a welded joint is manufactured using the steel sheet according to the present invention, a welded joint is obtained that is extremely excellent in the toughness of the weld heat affected zone at −60 ° C. and the brittle crack propagation stopping performance of the base material. By using the steel plate according to the present invention, it is possible to reduce the cost and improve the safety of a welded steel structure used in a low temperature environment such as a cold region and an ice sea region.

  The present inventors have intensively studied to solve the above problems, and as a result, by controlling the chemical composition, metal structure, hardness, diameter and number density of precipitated particles, crystal grain size, and texture of the steel sheet. It was found that a steel sheet with low cost and high productivity during hot rolling can be obtained. Furthermore, the present inventors have found that when a welded joint is produced using this steel sheet, the joint toughness (HAZ toughness) and base metal arrestability at −60 ° C. of the welded joint can be improved.

  Hereinafter, a steel sheet according to an embodiment of the present invention made based on the above knowledge will be described. In the following description, the terms “base metal” and “heat-affected zone” refer to the base material and weld heat-affected zone of a welded joint manufactured using the steel plate according to the present embodiment (unless otherwise noted). HAZ) respectively. The configuration and characteristics of the steel sheet can be considered the same as the configuration and characteristics of the base material.

(Metal structure of steel sheet)
The steel sheet according to the present embodiment is a mixed structure containing ferrite and bainite, or a mixed structure containing ferrite, pearlite and bainite, and the area ratio of ferrite is 50 to 90% and the area ratio of bainite is 10 to 50%. Has a metallographic structure.

(Ferrite area ratio of steel sheet: 50 to 90%)
When the ferrite area ratio of the steel sheet is less than 50%, it is difficult to improve the arrestability by setting the average grain size to 17.0 μm or less. When the ferrite area ratio is more than 90%, it is difficult to set the yield stress of the steel sheet to 270 MPa or more. If a steel sheet having a desired plate thickness and strength (tensile strength and yield stress) can be obtained, bainite or pearlite and bainite can be further included as the balance of ferrite. The lower limit value of the ferrite area ratio may be 55%, 60%, 65%, or 69%. The upper limit of the ferrite area ratio may be 86%, 83%, 80%, or 77%.

(Bainite area ratio of steel sheet: 10 to 50%)
Bainite greatly affects the strength, average crystal grain size, and texture of the steel sheet. When the bainite area ratio is less than 10%, it is difficult to set the yield stress of the steel sheet to 270 MPa or more. When the bainite area ratio is more than 50%, the average crystal grain size of the steel sheet is 17.0 μm or less and, as will be described later, an {100} area that forms an angle of 15 ° or less with respect to a plane perpendicular to the rolling direction It is difficult to improve the arrestability of the steel sheet by setting the rate to 20% or less at the plate thickness ½ part of the steel sheet. The lower limit of the bainite area ratio may be 15%, 18%, 20%, or 23%. The upper limit of the bainite area ratio may be 45%, 40%, 35%, 31%, or 28%.

(Perlite area ratio of steel sheet: preferably 10% or less)
The pearlite may be contained in the steel sheet as long as a steel sheet having a desired thickness and strength can be obtained. Therefore, the pearlite area ratio is not specified. However, since pearlite may deteriorate the fracture resistance of the steel sheet, the pearlite area ratio may be limited to 10% or less, 5% or less, or 3% or less. The lower limit value of the pearlite area ratio is 0%.

(MA area ratio of steel sheet: 0 to 5%)
In addition to ferrite, pearlite, and bainite, fine MA (Martensite-Austenite-constituent) may be present in the metal structure of the steel sheet. However, since MA may reduce the fracture resistance of the steel sheet, the MA area ratio is set to 0 to 5%. The upper limit value of the MA area ratio may be 3%, 2%, or 1%. The MA area ratio is most preferably 0%.

  It is not necessary to define the area ratio of the metal structure (the remainder of the metal structure) other than ferrite, bainite, pearlite, and MA. When the steel sheet is manufactured so that the chemical composition of the steel sheet is within the following range and the area ratio of ferrite, bainite and pearlite is within the above specified range, it is considered that no metal structure other than the above metal structure is generated. It is done. Further, even if a metal structure other than the above-described metal structure is generated, the generation amount is small enough not to affect the characteristics of the steel sheet. The area ratio of metal structures other than ferrite, bainite, pearlite, and MA may be limited to 1% or less in total.

(Average crystal grain size of steel sheet: 3.0 to 17.0 μm)
The factor that mainly controls the arrestability of the steel plate and the weld heat affected zone is the amount (area) of the grain boundary. This is because the crystal grain boundary becomes an obstacle to brittle crack propagation. In other words, since adjacent crystal grains sandwiching the crystal grain boundary have different crystal orientations, the direction in which the crack propagates changes at the crystal grain boundary. For this reason, an unbroken region is generated at the crystal grain boundary, and the stress is dispersed by the unbroken region to generate a crack closing stress. Therefore, when the amount of crystal grain boundaries in the metal structure is large, the driving force for crack propagation is reduced and the arrestability is improved. In addition, at the crystal grain boundary, since the unbroken region finally undergoes ductile fracture, energy required for brittle fracture is absorbed. For this reason, increasing the amount of grain boundaries improves the arrestability.

In order to increase the amount of crystal grain boundaries, it is necessary to make the crystal grain size fine. In the steel sheet according to the present embodiment, a boundary between adjacent crystals having a crystal orientation difference of 15 ° or more is defined as a crystal grain boundary, and a region surrounded by the crystal grain boundary is defined as a crystal grain. When the division number N is set to an integer of 10 or more, the weighted average value D _AVE of the grain sizes of the crystal grains defined by the formula 3 is 3.0 to 17.0 μm at the plate thickness of ½ part. .
_{p = (D MAX -D MIN)} / N ... (1 type)
D _k = D _MIN + p × (k−1 / 2) (Expression 2)
D _AVE = (Σ [k = 1, N] (D _k × S _k )) / (Σ [k = 1, N] S _k ) (Expression 3)
D _MAX indicates the maximum grain size in μm, D _MIN indicates the minimum crystal grain size in μm, k is an integer of 1 to N, and S _k is ( D _MIN + p × (k−1)) μm or more (D _MIN + p × k) The total area ratio of crystal grains having an equivalent circle diameter of less than μm is shown in unit%.

  As described above, the “crystal grain boundary” in the present embodiment means “a boundary between adjacent crystals whose crystal orientation difference is 15 ° or more”. A crystal grain boundary having a crystal orientation difference of less than 15 ° is not considered to be a hindrance to brittle crack propagation and does not have a great arrestability improving effect, and thus is not considered in this embodiment. The “plate thickness 1 / X portion” in the present embodiment indicates a region at a depth of about 1 / X of the plate thickness from the steel plate surface.

In the present embodiment, a weighted average value of crystal grain sizes is used to define the crystal grain size. The weighted average value is calculated by the above-described formulas 1, 2, and 3. “P” is the width of the division range, and is obtained by dividing the difference between the maximum value D _MAX and the minimum value D _MIN of the crystal grain size (the width of all crystal grain size distributions) by the division number N (1 See formula). “K” is an arbitrary integer of 1 or more and N or less. “D _k ” is the median value of the k-th division range (see Formula 2). The weighted average value D _AVE is first a value _{obtained by} multiplying the median value D _k of the kth division range by the amount S _k (unit%) of crystal grains having a crystal grain size within the kth division range (D _K × S _K ), and then D ₁ × S _{1 to} D _N × S _N is obtained (see Equation 3).
The weighted average value is excellent as an index for evaluating the presence of a small number of coarse particles (existence of abnormal values). Arrest is influenced by a small number of coarse particles and thus shows a good correlation with the weighted average value. On the other hand, the arithmetic average value of the crystal grain size that is usually most commonly used as a method for evaluating the crystal grain size is obtained by calculating the square root of the value obtained by dividing the area of the measurement visual field by the number of crystal grains in the measurement visual field. can get. When the crystal grain size is evaluated by the arithmetic average, the presence of a small number of coarse particles is not sufficiently reflected in the finally obtained arithmetic average value, so that the arithmetic average value and arrestability do not show a good correlation.
The weighted average value of the crystal grain size may be about twice the arithmetic average value of the crystal grain size. Further, when bainite having a large variation in crystal grain size is included in the metal structure, the weighted average value of the crystal grain size can be 3 to 10 times the arithmetic average value of the crystal grain size. In this embodiment, the term “average crystal grain size” means “weighted average value of crystal grain size” unless otherwise specified.
Here, the arithmetic average of the crystal grain size is an average crystal grain size (unit: mm) defined in Table 1 of JIS G0551: 2013. Specifically, the grain size is 1 / √m, where m is the number of crystal grains per 1 mm ² of the observed test surface.

When the average crystal grain size of the steel sheet is 3.0 to 17.0 μm at 1/2 part of the plate thickness, the arrest toughness value (Kca) of the steel sheet at −60 ° C. is 4000 N / mm ^1.5 or more. In order to further improve the arrestability, the average crystal grain size may be 15.0 μm or less, 13.0 μm or less, 11.0 μm or less, 10.0 μm or 9.0 μm or less. When the average crystal grain size exceeds 17.0 μm, the arrestability of the steel sheet is insufficient.

  In order to improve various properties of the steel sheet, it is not necessary to define a lower limit value of the average crystal grain size of the steel sheet. However, the arrestability improves as the average crystal grain size becomes finer, but excessive refinement increases the rolling load and lowers the productivity. Therefore, the lower limit value of the average crystal grain size is preferably 3.0 μm. This lower limit may be 4.0 μm, 5.0 μm, or 6.0 μm.

  The reason for prescribing the average crystal grain size of the steel sheet by 1/2 part thickness is that the arrestability of the steel sheet having a thickness of 10 to 35 mm is mainly governed by the average crystal grain diameter of 1/2 part thickness. This is because the average crystal grain size of ½ part of the plate thickness can be a representative value of the average crystal grain size of the steel plate. In addition, according to the manufacturing method described later that mainly controls the average crystal grain size of 1/2 part of the plate thickness, the temperature is inevitably lowered at the plate thickness positions other than 1/2 part of the plate thickness, and the cooling rate is It becomes larger and the crystal grains tend to become finer. Accordingly, it is not necessary to particularly define the average crystal grain size at plate thickness positions other than 1/2 plate thickness.

  For the measurement of the crystal grain size of the steel sheet, it is preferable to use an EBSD (Electron Back Scatter Diffraction Pattern) method capable of accurately measuring crystal orientation information with a wide field of view. If the EBSD method is used, the crystal grain size of a complex structure such as bainite can be measured.

  More specifically, the crystal grain size is determined by the following method. First, the measurement of a 500 μm × 500 μm area of ½ part of the plate thickness by the EBSD method is repeated while moving the measurement position by 1 μm. Here, the boundary where the crystal orientation difference between adjacent grains is 15 ° or more is defined as a crystal grain boundary, the equivalent circle diameter of this crystal grain boundary is defined as the crystal grain diameter, and the average crystal grain diameter is calculated from the above measurement results. Ask.

(Area ratio of {100} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the rolling direction in the sheet thickness 1/2 part of the steel sheet: 2.0 to 20.0%)
Only by controlling the average crystal grain size, it is difficult to stably improve the arrestability in an extremely low temperature environment of −60 ° C. Control of the crack propagation direction utilizing the texture is required for the steel sheet according to this embodiment. This indicates that a brittle crack generated in a steel sheet when the steel sheet is subjected to external stress propagates along a {100} cleaved surface that forms an angle of 15 ° or less with respect to a plane perpendicular to the rolling direction. The inventors have found. That is, if a {100} plane texture develops in a plane perpendicular to external stress, the effect of improving arrestability obtained by controlling the average crystal grain size to 3.0 to 17.0 μm is reduced. It has been found. The external stress is a stress applied externally to the steel structure. Brittle cracks often occur and propagate in a direction perpendicular to the direction of the largest external stress. Therefore, here, the largest stress among externally applied stresses to the steel structure is defined as external stress. In general, the external stress is applied substantially parallel to the main rolling direction of the steel sheet. For this reason, a plane perpendicular to the external stress can be regarded as a plane perpendicular to the main rolling direction of the steel sheet.

  The main rolling direction of the steel sheet can be determined by, for example, corroding the steel sheet surface with picric acid and measuring the aspect ratio of prior austenite. The direction in which the aspect ratio of the prior austenite is large can be determined as the main rolling direction of the steel sheet.

  The texture of the {100} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the rolling direction of the steel sheet is adjusted so that the area ratio is 2.0 to 20.0% at 1/2 part of the plate thickness. It has been found that the resistance to the propagation of brittle cracks can be increased and the arrestability can be improved.

  Based on the above knowledge, in the steel sheet according to the present embodiment, the area ratio of the {100} plane, which forms an angle of 15 ° or less with respect to the plane perpendicular to the rolling direction, is 2.0 in the plate thickness 1/2 part. ˜20.0%.

  It is desirable that the area ratio of the {100} plane, which forms an angle of 15 ° or less with respect to the plane perpendicular to the rolling direction, in the sheet thickness ½ part is smaller. However, in order to make the area ratio of the {100} plane that forms an angle within 15 ° with respect to the plane perpendicular to the rolling direction less than 2.0%, the metal structure of the steel sheet has a ferrite area ratio of 90%. It is necessary to make the metal structure exceed. In this case, it becomes difficult to set the yield stress to 270 MPa or more. Therefore, the lower limit of the area ratio of the {100} plane that forms an angle within 15 ° with respect to the plane perpendicular to the rolling direction is set to 2.0%. When the area ratio of the {100} plane forming an angle of 15 ° or less with respect to the plane perpendicular to the rolling direction is more than 20.0%, the arrestability is lowered even if the average crystal grain size is finely controlled. Therefore, the upper limit of the area ratio of the {100} plane that forms an angle within 15 ° with respect to the plane perpendicular to the rolling direction is set to 20.0%. The lower limit value of the area ratio of the {100} plane that forms an angle within 15 ° with respect to the plane perpendicular to the rolling direction may be 3.0%, 4.0%, or 5.0%. Further, the upper limit of the area ratio of the {100} plane may be 17.0%, 15.0%, 12.0%, or 10.0%.

  The texture of the steel plate is preferably measured by the EBSD method. When measuring by the EBSD method, it is possible to measure a texture having a wider field of view with higher accuracy than measurement by X-rays.

  More specifically, a map of {100} planes forming an angle of 15 ° or less with respect to a plane perpendicular to the main rolling direction of the steel sheet at a thickness of 1/2 part is created by the EBSD method, and the total area is By dividing by the measurement area, the area ratio can be obtained.

The measures for improving the arrestability as described above are applied to a steel plate having a yield stress of 300 to 500 MPa and a plate thickness of 10 to 35 mm. A steel plate having a yield stress of less than 300 MPa or a thickness of less than 10 mm is rarely used as a material for a member that requires arrestability. Moreover, even if the average crystal grain size and texture defined in the present embodiment are formed in a steel plate having a yield stress of over 500 MPa and a plate thickness of over 35 mm, the mechanical condition becomes severe, and the arrest toughness value at −60 ° C. This is because it may be difficult to impart a high arrestability of (Kca) of 4000 N / mm ^1.5 or more. The lower limit value of the yield stress may be limited to 320 MPa, 340 MPa, or 360 MPa, and the upper limit value may be limited to 480 MPa, 460 MPa, or 440 MPa. The lower limit value of the plate thickness may be limited to 12 mm, 14 mm, or 16 mm, and the upper limit value may be limited to 30 mm, 26 mm, or 22 mm.
Although it is not necessary to specify the tensile strength, it may be set to 450 to 700 MPa. If necessary, the lower limit value of the tensile strength may be 490 MPa, and the upper limit value of the tensile strength may be 650 MPa or 600 MPa.

(Total number density of TiN particles, MnS particles, and composite particles each having an equivalent circle diameter of 0.5 to 2.0 μm in the steel sheet: 20 to 200 particles / mm ² )
Particles such as precipitates or inclusions contained in the steel plate greatly affect the toughness of the weld heat affected zone. TiN particles, MnS particles, and composite particles, when steel plates are heated by welding, suppress the growth of austenite grains by the pinning effect, and when the steel plates are cooled after welding, the ferrite transforms As a result, the structure can be refined and the toughness of the welded joint can be improved. However, TiN particles, MnS particles, and composite particles may become the starting point of brittle fracture, and the toughness of the steel sheet may be reduced. Therefore, it is necessary to precisely control the number of TiN particles, MnS particles, and composite particles. In order to improve joint toughness by refining the HAZ structure in a welded joint, it is desirable to use TiN particles, MnS particles, and composite particles as ferrite transformation nuclei. This is because when only the above-described pinning effect is used, the metal structure is not sufficiently subdivided by ferrite transformation in the austenite grains, and thus the metal structure is not sufficiently refined. In addition, when only the pinning effect is used to sufficiently refine the metal structure, it is necessary to disperse a large amount of TiN particles, MnS particles, and composite particles, which can be a starting point for brittle fracture. The number of points increases and the toughness of the welded joint decreases.

Based on the above knowledge, the steel sheet according to the present embodiment has the following configuration in order to use TiN particles, MnS particles, and composite particles as ferrite transformation nuclei. A precipitate or inclusion containing 1% by mass or more of Ti, less than 1% by mass of O and 1% by mass or more of N is defined as TiN particles in the present embodiment, and 1% by mass or more of Mn and 1% by mass or more. The precipitates or inclusions containing S and less than 1% by mass of O are defined as MnS particles in this embodiment, and the precipitates or inclusions that simultaneously satisfy the definition of TiN particles and the definition of MnS particles are defined in this embodiment. (Hereinafter, unless otherwise specified, the terms “TiN particles”, “MnS particles”, and “composite particles” refer to particles that satisfy the above definition). In this case, the total number density of TiN particles, MnS particles, and composite particles each having an equivalent circle diameter of 0.5 to 2.0 μm is set to 20 to 200 particles / mm ² .

TiN particles, MnS particles, and composite particles in the present embodiment serve as ferrite transformation nuclei. Further, TiN particles, MnS particles, and composite particles also serve as precipitation nuclei for BN (boron nitride), which is another ferrite transformation nucleus. The reason for selecting TiN particles, MnS particles, and composite particles as ferrite transformation nuclei is as follows. TiN particles have low interfacial energy with ferrite (that is, good lattice matching). The MnS particles generate a Mn-deficient layer in the vicinity of the interface with the matrix to enhance the driving force for ferrite transformation. Further, TiN particles and MnS particles have high performance as the core of ferrite transformation and BN precipitation by increasing the interfacial energy with austenite, and greatly contribute to the refinement of the HAZ structure.
It is preferable that TiN and MnS exist as composite particles (particles satisfying both the definition of TiN particles and the definition of MnS particles) rather than each of them alone. The composite particles further enhance the above effect. However, the influence of the coexistence of TiN and MnS on the refinement of the structure is negligibly small as compared with the influence of the particle diameter and the number density on the refinement of the structure. Therefore, any one or more of TiN particles and MnS particles may be present.

The reason why the equivalent circle diameter of the TiN particles, MnS particles, and composite particles is defined as 0.5 to 2.0 μm and the number density is defined as 20 to 200 particles / mm ² is as follows. When the particle diameter of TiN particles, MnS particles, and composite particles is 0.5 μm and / or the number density is less than 20 particles / mm ² , the interfacial energy between TiN particles, MnS particles, composite particles, and austenite Cannot be made sufficiently large, the ability of these particles as transformation nuclei is reduced, and the amount of nucleation sites is reduced, so that the average crystal grain size of the weld heat affected zone may not be 80 μm or less. is there. When the particle diameter of TiN particles, MnS particles, and composite particles is 2.0 μm and / or the number density exceeds 200 particles / mm ² , these particles are likely to be the starting point of brittle fracture, and the toughness of the steel sheet is reduced. May end up. The lower limit value of the equivalent circle diameter of each of the TiN particles, the MnS particles, and the composite particles may be 0.6 μm, 0.7 μm, or 0.8 μm. Further, the upper limit value of the equivalent circle diameter of each of the TiN particles, the MnS particles, and the composite particles may be 1.8 μm, 1.5 μm, or 1.3 μm. For improving toughness, TiN particles, MnS particles, and 30 the lower limit of the composite particles each number density / ^mm 2, 40 pieces / ^mm 2, may be 50 / mm ² or 60 / mm ^2. In view of the practical range of the number density of each particle, the upper limit of the number density of each of the TiN particles, MnS particles, and composite particles is 180 particles / mm ² , 160 particles / mm ² , 140 particles / mm ^2, or It may be 120 / mm ² . Note that TiN particles, MnS particles, and composite particles (coarse particles) having a particle diameter of more than 2.0 μm may be the starting point of brittle fracture and may reduce joint toughness. Therefore, the number density of each particle diameter is 2.0μm greater is preferably 50 pieces / mm ² or less or 20 / mm ² or less. The number density of TiN particles, MnS particles, and composite particles having a particle diameter of less than 0.5 μm has low ferrite nucleation ability and little influence on the metal structure. Therefore, it is not necessary to particularly define the number density of each particle having a particle diameter of less than 0.5 μm.
It is not necessary to limit the positions that define the number and particle size of TiN particles, MnS particles, and composite particles. This is because the position has a relatively small influence on the number and particle size of the particles. In the present embodiment, the numbers and particle sizes of TiN particles, MnS particles, and composite particles at a thickness of 1/4 part are used as representative values for steel plates.

(Number density of oxide particles having an equivalent circle diameter of 1 to 10 μm: 20 to 200 / mm ² )
Furthermore, in the steel plate according to the present embodiment, it is preferable that the number density of oxide particles having a circle-equivalent diameter of 1 to 10 μm at a thickness of ¼ part is 20 to 200 particles / mm ² .

  The oxide particles are not used as pinning particles and / or ferrite transformation nuclei. Furthermore, the oxide particles serve as a starting point for brittle fracture and reduce joint toughness. Therefore, it is preferable to reduce the number density of the oxide particles as much as possible. However, if the number density of the oxide particles is excessively reduced, problems such as an increase in steel sheet manufacturing cost occur, which is not preferable for industrial use.

When the deoxidation step is performed to the maximum extent acceptable for industrial use, the equivalent circle diameter of the oxide particles present as impurities is about 1 μm, and the number density is about 20 particles / mm ² . Therefore, the lower limit of the equivalent circle diameter of the oxide particles is 1 μm, and the lower limit of the number density of the oxide particles is 20 particles / mm ² . The lower limit of the number density of the oxide particles may be 0 / mm ² , 10 / mm ² , 20 / mm ² , 40 / mm ² , or 60 / mm ² . Presence of oxide particles having an equivalent circle diameter of less than 1 μm is acceptable. This is because oxide particles having an equivalent circle diameter of less than 1 μm hardly affect the toughness. Therefore, it is not necessary to define the number density of oxide particles having an equivalent circle diameter of less than 1 μm.

When the equivalent circle diameter of the oxide particles is more than 10 μm and / or the number density is more than 200 pieces / mm ² , the oxide particles become a starting point of brittle fracture and may reduce the toughness of the steel sheet. Therefore, the upper limit of the equivalent circle diameter of the oxide particles was 10 μm, and the upper limit of the number density of the oxide particles was 200 particles / mm ² . The upper limit of the number density of preferable oxide particles is 150 / mm ² , and more preferably 100 / mm ² . In particular, the upper limit of the oxide particles may be set to 8 μm or 6 μm in order to reduce oxide particles that are likely to be the starting point of destruction. Oxide particles (coarse particles) having an equivalent circle diameter of more than 10 μm are preferably not contained in the steel plate. However, when the number density of the oxide particles having an equivalent circle diameter of 1 to 10 μm is reduced as described above, it is considered that the oxide particles having an equivalent circle diameter of more than 10 μm are sufficiently removed. It is not necessary to define the number density of oxide particles having an equivalent circle diameter of more than 10 μm.

  There is no need to limit the position for defining the number and particle size of the oxide particles. This is because the position has a relatively small influence on the number and particle size of the oxide particles. In the present embodiment, the number of oxide particles and the particle size at a ¼ part thickness are used as representative values of the steel plate.

  In order to measure the number density and particle size of the oxide particles, it is preferable to measure with a wide field of view using a scanning electron microscope (SEM: Scanning Electron Microscope).

  More specifically, in a 1 mm × 1 mm region, oxide particles are identified by a backscattered electron image, and the composition is analyzed by EDS (Energy Dispersive X-ray Spectrometer) included in the SEM to identify the oxide particles. Further, an oxide particle image is taken, an equivalent circle diameter is obtained by image analysis, and the number density of the oxide particles is obtained by counting the number. You may measure this oxide particle using the software attached to SEM which can be performed automatically and continuously. The plate thickness position at the time of measurement may be arbitrary, but the thickness is measured with 1/4 part of the plate thickness as a representative.

(Chemical composition)
Hereinafter, the chemical composition of the steel sheet according to the present embodiment will be described. “%” For a component means mass%.

(C: 0.040-0.090%)
C is contained by 0.040% or more in order to ensure the strength of the base material. If the C content exceeds 0.090%, the embrittled phase such as cementite and MA increases, and it becomes difficult to ensure joint toughness. Therefore, the C content is set to 0.090% or less. . Therefore, the lower limit of the C content is 0.040%, preferably 0.050%, and more preferably 0.060%. The upper limit of the C content is 0.090%, preferably 0.080%, more preferably 0.070%.

(Si: 0.01-0.20%)
Since Si is effective as a deoxidizing element and a strengthening element, it is contained in an amount of 0.01% or more. If the Si content exceeds 0.20%, MA increases and the joint toughness deteriorates greatly, so the Si content is 0.20% or less. The lower limit of the Si content is 0.01%, preferably 0.03%, more preferably 0.05%. The upper limit of the Si content is 0.20%, preferably 0.15%, more preferably 0.10%.

(Mn: 1.30 to 1.80%)
Mn is contained in an amount of 1.30% or more in order to ensure the strength and toughness of the thick base material. If the Mn content exceeds 1.80%, the center segregation of Mn becomes prominent, and the toughness of the base metal and the HAZ in the portion where the center segregation has occurred deteriorates, so the Mn content is 1.80% or less. And The lower limit of the Mn content is 1.30%, preferably 1.35%, more preferably 1.40%. The upper limit of the Mn content is 1.80%, preferably 1.70%, more preferably 1.60%.

(P: limited to 0.020% or less)
P is one of impurity elements. In order to ensure joint toughness stably, the P content is limited to 0.020% or less. Preferably, the P content is 0.015% or less, more preferably 0.010% or less. The smaller the P content, the better. Therefore, the lower limit of the P content is 0%. However, considering the cost for reducing the P content, 0.0005% may be set as the lower limit of the P content.

(S: 0.001 to 0.010%)
S is one of the impurity elements, but is an element necessary for generating MnS particles. In order to generate MnS particles and ensure joint toughness, 0.001% or more of S is contained. However, if the S content exceeds 0.010%, coarse MnS particles increase and the joint toughness decreases, so the S content is set to 0.010% or less. The lower limit of the S content is 0.001%, preferably 0.0015%, more preferably 0.002%. The upper limit of the S content is 0.010%, preferably 0.008%, more preferably 0.006% or 0.005%.

(Al: 0.005-0.100%)
Al has a role of deoxidizing steel and reduces O which is one of impurity elements. In addition to Al, Si and Mn also contribute to deoxidation. However, even when Si and Mn are added, if the Al content is less than 0.005%, O cannot be stably reduced. However, when the Al content exceeds 0.1%, alumina-based coarse oxides and clusters of coarse oxides are generated in the metal structure, and the toughness of the base material and the HAZ is impaired. Therefore, the Al content is 0.100% or less. The lower limit of the Al content is 0.005%, preferably 0.010%, more preferably 0.020%. The upper limit of the Al content is 0.100%, preferably 0.080%, more preferably 0.060% or 0.050%.

(Nb: 0.003-0.030%)
Nb is an important element in the present embodiment. In order to form a predetermined average crystal grain size and texture, a steel sheet is manufactured by performing rolling with the rolling temperature set to an unrecrystallized austenite region (a temperature region where unrecrystallized austenite is maintained in the metal structure). It is necessary to do. Nb is an element effective for expanding the non-recrystallized austenite region. When the non-recrystallized austenite region is expanded, the rolling temperature can be raised, which contributes to productivity improvement. In order to acquire this effect, it is necessary to contain Nb 0.003% or more. However, if the Nb content exceeds 0.030%, the hardness of the HAZ excessively increases and the joint toughness decreases, so the Nb content is set to 0.030% or less. The lower limit of the Nb content is 0.003%, preferably 0.004%, and more preferably 0.005%. The upper limit of Nb content is 0.020%, preferably 0.015%, more preferably 0.010%.

(Ti: 0.003-0.030%)
Ti is an important element in the present invention. By containing Ti, TiN particles are formed in the steel sheet. The TiN particles suppress the austenite particle size from increasing when the steel plate is heated by the pinning effect. When the austenite grain size is increased, the average crystal grain size of the transformed structure is also increased, so that it becomes difficult to obtain a predetermined average crystal grain size, and the toughness and arrestability of the steel sheet are lowered. In addition, the fine TiN particles used for the pinning effect in the base material are mostly dissolved during the heating of welding, but the coarse TiN particles remain undissolved in the HAZ, and the nuclei of precipitation of BN particles and the transformation nuclei of ferrite And contributes to the refinement of the HAZ structure. When TiN particles are not included, it is difficult to obtain a predetermined average crystal grain size, and joint toughness is reduced. In order to obtain a necessary average crystal grain size and improve toughness, arrestability, and joint toughness, it is necessary to contain 0.003% or more of Ti. However, if the Ti content exceeds 0.030%, coarse TiN particles and TiC particles that can become the starting point of brittle fracture are formed, and the toughness and joint toughness of the steel sheet are lowered. Therefore, the Ti content is 0.030% or less. The lower limit of the Ti content is 0.003%, preferably 0.005%, more preferably 0.007%. The upper limit of the Ti content is 0.030%, preferably 0.020%, more preferably 0.015%.

(B: 0.0003 to 0.0040%)
B is an important element in the present invention. In HAZ, B forms BN particles. Since the BN particles have a function as a ferrite transformation nucleus, the B content contributes to the refinement of the HAZ structure. In order to obtain this effect, it is necessary to contain 0.0003% or more of B. However, if the B content is too large, joint toughness and weldability deteriorate, so 0.0040% is made the upper limit of the B content. If the content of B exceeds 0.0040%, coarse BN particles that may become the starting point of brittle fracture are formed, or the hardness of the HAZ becomes excessive due to excessive increase in hardenability by increasing the solid solution B. The joint toughness decreases due to the increase. Therefore, the B content is 0.0040% or less. The lower limit of the B content is 0.0003%, preferably 0.0005%, and more preferably 0.0007%. The upper limit of the B content is 0.0040%, preferably 0.0030%, more preferably 0.0020% or 0.0016%.

(N: 0.0020 to 0.0080%)
N is an important element in the present invention. In order to form TiN particles and BN particles as described above and improve toughness, arrestability, and joint toughness, it is necessary to contain 0.0020% or more of N. However, if the N content exceeds 0.0080%, the steel sheet becomes brittle, so the N content is set to 0.0080% or less. The lower limit of the N content is 0.0020%, preferably 0.0025%, more preferably 0.0030%. The upper limit of the N content is 0.0080%, preferably 0.0070%, more preferably 0.0060% or 0.0050%.

(O: 0.0005-0.0040%)
O is one of impurity elements. In order to suppress the generation of coarse oxide particles and to ensure joint toughness stably, the O content is limited to 0.0040% or less. Preferably, it is 0.0030% or less, More preferably, it is 0.0020% or less. The lower limit value is 0.0005%, preferably 0.0010%, more preferably 0.0015% in consideration of the cost for reducing the O content.

  In the chemical composition of the steel sheet according to the present embodiment, the balance of the elements described above may be Fe and impurities. However, the chemical composition of the steel sheet according to the present embodiment may contain at least one of Cu, Ni, Cr, Mo, V, Ca, Mg, and REM as necessary. Moreover, even if these elements are contained in trace amounts at the impurity level, it is acceptable in this embodiment. Hereinafter, the addition effect and content of each element will be described. Even if these elements are intentionally added, even if they are mixed as impurities, the steel sheet whose content of these elements is within the specified range is regarded as the steel sheet according to the present embodiment.

(Cu: 0 to 1.00%)
By containing Cu, the strength and toughness of the base material can be improved. In order to obtain the effect, 0.10% or more of Cu may be contained. However, when there is too much content of Cu, joint toughness and weldability will deteriorate, so 1.00% is made the upper limit of Cu content. The lower limit of the Cu content is 0%, and may be 0.10%, 0.15%, or 0.20% in order to improve the strength and toughness of the base material. In order to improve joint toughness and weldability, the upper limit of the Cu content is 1.00%, and if necessary, even if limited to 0.80%, 0.50%, or 0.30% Good.

(Ni: 0 to 1.00%)
By containing Ni, the strength and toughness of the base material can be improved. In order to obtain the effect, 0.10% or more of N may be included. However, when there is too much content of Ni, manufacturing cost will become high and also joint toughness and weldability will deteriorate. Therefore, 1% is taken as the upper limit of the Ni content. The lower limit of the Ni content is 0%, and may be 0.15% or 0.20% in order to improve the strength and toughness of the base material. In order to reduce costs and / or improve joint toughness and weldability, the upper limit of the Ni content is 1.00%, depending on necessity, 0.70%, 0.50%, or 0.30. % May be limited.

(Cr: 0 to 1.00%)
By containing Cr, the strength and toughness of the base material can be improved. In order to obtain the effect, 0.10% or more of Cr may be included. However, if the Cr content is too high, joint toughness and weldability deteriorate, so 1.00% is made the upper limit of the Cr content. Therefore, the lower limit of the Cr content is 0.05% or 0.10%. The upper limit of Cr may be 1.00%, 0.80%, 0.50%, 0.30%, 0.20%, or 0.12% as necessary.

(Mo: 0 to 0.50%)
By containing Mo, the strength and toughness of the base material can be improved. In order to obtain the effect, 0.01% or more of Mo may be contained. However, if the Mo content is too large, joint toughness and weldability deteriorate, so 0.50% is made the upper limit of the Mo content. The lower limit of the Mo content is 0%, and may be 0.01%, 0.02%, or 0.03%. The upper limit of the Mo content is 0.50%, and may be 0.30%, 0.20%, 0.10%, or 0.06% as necessary.

(V: 0 to 0.1%)
By containing V, the strength and toughness of the base material can be improved. In order to obtain the effect, 0.01% or more of V may be contained. However, if the V content is too high, joint toughness and weldability deteriorate, so 0.10% is made the upper limit of the V content. The lower limit of the V content is 0%, and may be 0.01%, 0.015%, or 0.02%. The upper limit of the V content is 0.10%, and may be 0.08%, 0.06%, or 0.04% as necessary.

(Ca: 0 to 0.0050%)
By containing Ca, joint toughness is improved. In order to obtain the effect, 0.0003% or more of Ca may be contained. However, when there is too much content of Ca, joint toughness and weldability deteriorate, so 0.0050% is made the upper limit of Ca content. The lower limit of the Ca content is 0%, and may be 0.0003%, 0.0005%, or 0.0007%. The upper limit of the Ca content is 0.0050%, and may be 0.0040%, 0.0030%, or 0.0020% as necessary.

(Mg: 0 to 0.0050%)
By containing Mg, joint toughness is improved. In order to obtain the effect, 0.0003% or more of Mg may be contained. However, if the Mg content is too high, joint toughness and weldability deteriorate, so 0.0050% is made the upper limit of the Mg content. The lower limit of the Mg content is 0%, and may be 0.0003%, 0.0005%, or 0.0007%. The upper limit of the Mg content is 0.0050%, and may be 0.0050%, 0.0040%, 0.0030%, or 0.0020% as necessary.

(REM: 0-0.0050%)
The term “REM” refers to a total of 17 elements composed of Sc, Y and lanthanoid, and “REM content” means the total content of these 17 elements. By including REM, joint toughness is improved. In order to obtain the effect, 0.0003% or more of REM may be contained. However, if the content of REM is too large, joint toughness and weldability deteriorate, so 0.0050% is made the upper limit of the REM content. The lower limit of the REM content is 0%, and may be 0.0003%, 0.0005%, or 0.0007%. The upper limit of the REM content is 0.0050%, and may be 0.0040%, 0.0030%, or 0.0020% as necessary.

In order to improve the strength and toughness of the steel plate and the welded joint, the above-mentioned selective elements can be intentionally included. However, there is no problem even if these selective elements are not included in order to reduce alloy costs. Therefore, the lower limit values for the contents of the above-mentioned selective elements are all 0%. However, even when not intentionally added, as impurities, Cu: less than 0.10%, Ni: less than 0.10%, Cr: less than 0.10%, Mo: less than 0.010%, V: Less than 0.010%, Ca: less than 0.0003%, Mg: less than 0.0003%, REM: less than 0.0003% may be contained in the steel. Even if these elements are contained as impurities in the steel, they do not affect the properties of the steel sheet according to this embodiment.
The steel sheet according to the present embodiment contains the essential components, the content of the selective components is limited, and the balance contains Fe and impurities. However, the welded steel material according to the present embodiment contains the following alloy elements in addition to the above components, for the purpose of further improving the strength, toughness, etc. of the steel material itself, or as impurities from auxiliary materials such as scrap. Also good.
Sb impairs joint toughness. Accordingly, the Sb content [Sb] is preferably 0.010% or less, more preferably 0.005% or less, and most preferably 0.003% or less.
Sn impairs joint toughness. Accordingly, the Sn content [Sn] is preferably 0.010% or less, more preferably 0.005% or less, and most preferably 0.003% or less.
As deteriorates joint toughness. Accordingly, the As content [As] is preferably 0.010% or less, more preferably 0.005% or less, and most preferably 0.003% or less.
Moreover, in order to fully exhibit the above-described effects of the essential component and the selection component, the contents of Zr, Co, Zn, and W may be limited to 0.01% or less or 0.005% or less. preferable.
There is no need to limit the lower limit of Sb, Sn, As, Zr, Co, Zn, and W, and the lower limit of the content of each element is 0%. In addition, alloy elements (for example, Cu, Ni, Cr, Mo, V, Ca, Mg, REM, Sb, Sn, As, Zr, Co, Zn, and W) that do not have a lower limit are intentionally added. Even if it is mixed as an impurity, if the content is within the specified range, it is determined that the steel material is a steel material according to the present embodiment.

(Carbon equivalent CE: 0.30 to 0.40%)
In the steel plate according to the present embodiment, the carbon equivalent CE obtained by the following four formulas is set to 0.30 to 0.40%.
CE = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 (4 formulas)
Here, the element symbol described in the formula indicates the content of each component contained in the steel sheet in mass%. When there is a component having a content of 0%, the calculation is performed by substituting 0% as the content of the component in the above four formulas.

  When the carbon equivalent of the steel sheet is less than 0.30%, a steel sheet that satisfies the requirements regarding yield strength of 300 MPa or more, arrestability at −60 ° C., and joint toughness cannot be obtained. When the carbon equivalent exceeds 0.40%, a steel sheet that satisfies the requirements regarding the arrestability at −60 ° C. and the joint toughness cannot be obtained. Therefore, the lower limit value of the carbon equivalent of the steel sheet is 0.30%, preferably 0.31%, more preferably 0.32%, still more preferably 0.33%, and the upper limit of the carbon equivalent is 0.40%. , Preferably 0.39%, more preferably 0.38%, still more preferably 0.37% or 0.36%.

(SOLB: -0.0015 to + 0.0015%)
In the steel plate according to the present embodiment, the SOLB (Solution B) obtained by the following formula 5 is set to -0.0015 to + 0.0015%.
SOLB = 0.226Ti + B−0.772N (5 formulas)
Here, the element symbol described in the formula indicates the content of each component contained in the steel sheet in mass%.
SOLB is an index for controlling the average crystal grain size and hardness of HAZ within a predetermined range and ensuring joint toughness. If the SOLB of the steel sheet is less than -0.0015%, the average crystal grain size of HAZ becomes coarse and the joint toughness at -60 ° C cannot be satisfied. When SOLB exceeds + 0.0015%, the presence of B excessively increases the hardenability of the steel sheet, makes the average particle size of HAZ 80 μm or less, and makes HAV Vickers hardness 190 or less. The joint toughness at −60 ° C. cannot be satisfied. Therefore, the lower limit value of SOLB is -0.0015%, preferably -0.0010%, more preferably -0.0005%, still more preferably 0%, and the upper limit value of SOLB is + 0.0015%, preferably + 0.0013%, more preferably + 0.0010%, still more preferably + 0.0008%.

  Next, the preferable manufacturing method of the steel plate which concerns on this embodiment is demonstrated.

(Melting process)
First, molten steel adjusted to a desired chemical composition is melted by a known melting method using a converter or the like. However, deoxidation and adjustment of the chemical composition must be performed as shown below during the melting process.
In the secondary refining by the RH method, first, the amount of dissolved oxygen in the steel is suppressed to 40 ppm or less by using Al that provides a strong deoxidizing power at low cost. Next, Ar is blown and refluxed for 5 minutes or more, thereby suppressing the amount of dissolved oxygen in the steel to 20 ppm or less. By performing deoxidation to the above-mentioned level, the number density of oxide particles having an equivalent circle diameter of 1 to 10 μm, which is the starting point of brittle fracture, can be reduced to 200 particles / mm ² or less. Decline can be prevented. The reason why Ar is blown in secondary refining by the RH method is that there is an effect of floating and removing coarse oxides such as alumina.
In order to adjust the chemical composition of the molten steel, it is necessary to input an alloy element in the melting process. Ti and B are preferably added after Al deoxidation. Ca, Mg, and REM are also preferably added after Al deoxidation if they are contained in the steel sheet. Since Ti, B, Ca, Mg, and REM form oxides, if added before Al deoxidation, the deoxidation reaction may be hindered and the yield may be reduced. Other elements may be added before or after deoxidation, but are usually added after deoxidation. Since Si and Mn have a deoxidizing effect, deoxidation can be performed more efficiently when introduced before Al deoxidation. However, even if Si and Mn are added after deoxidation, the properties of the steel sheet are not deteriorated.

Further, the molten steel is made into a steel piece (slab, slab) by a known casting method such as continuous casting. However, the cooling rate in the range of 900-1300 degreeC of the steel plate 1/4 thickness at the time of casting shall be 0.05-5.00 degreeC / s. This is for precipitating TiN particles, MnS particles, and composite particles. TiN particles, MnS particles, and composite particles used as ferrite transformation nuclei in HAZ will melt or become coarse during reheating in subsequent steps unless the precipitation control in the steel slab manufacturing stage is performed appropriately. . When dissolution or coarsening occurs, it becomes impossible to obtain a steel sheet containing TiN particles, MnS particles, and composite particles having a predetermined particle size and number density. A preferable cooling condition is that the cooling rate is 0.1 to 4 ° C./s in the range of 950 to 1250 ° C., and a more preferable cooling condition is 0.2 to 3 in the range of 1000 to 1200 ° C. .0 ° C./s.
It is desirable that the steel slab is cooled until the half thickness of the steel slab becomes 700 ° C. or less. That is, it is desirable that the cooling stop temperature of the slab thickness 1/2 part is 700 ° C. or less. Since the transformation from austenite to ferrite that occurs during cooling is not completed when the cooling stop temperature of the slab thickness of 1/2 part exceeds 700 ° C, the reverse transformation from ferrite to austenite during heating in the subsequent process It is difficult to obtain a fine graining effect using. When the effect of refinement cannot be obtained, austenite grains become coarse, and it becomes difficult to refine the crystal grain size after rolling. The cooling stop temperature is preferably 650 ° C. or lower, and more preferably 600 ° C. or lower. However, when the heating temperature in the subsequent process is 1100 ° C. or less and the rolling reduction of rough rolling can be 30% or more, the austenite grains can be refined by recrystallization. It is permissible to set the cooling stop temperature of half the thickness of the piece to more than 700 ° C.

(Reheating and soaking process)
When the slab is reheated for the production of a steel plate, the steel plate is soaked for 20 to 90 minutes in a state where the plate thickness center temperature is 950 to 1100 ° C in a heating furnace having a soaking atmosphere at 1000 to 1150 ° C. Keep heat.

  If the atmospheric temperature in the soaking zone is less than 1000 ° C., the steel slab cannot be heated sufficiently and the solution is insufficient. When the atmospheric temperature exceeds 1150 ° C., the austenite grains of the steel slab become coarse, and it becomes difficult to refine the crystal grains in the subsequent rolling process.

  If the sheet thickness center temperature during soaking is less than 950 ° C., solutionization becomes insufficient, and the austenite grains become finer and hardenability decreases, so that a steel sheet satisfying a predetermined strength can be obtained. Have difficulty. When the plate thickness center temperature during soaking is over 1100 ° C., the austenite grains become coarse, and it becomes difficult to refine the crystal grains in the subsequent rolling process. Further, the slab is heated to a temperature for starting rolling. Since there is a time to wait for the temperature to drop, productivity is lowered.

  If the holding time during soaking is less than 20 minutes, the solution is insufficient, and if it exceeds 90 minutes, the austenite grains become coarse and the TiN particles, MnS particles and composite particles are dissolved, and a predetermined particle diameter is obtained. And number density is not obtained.

(Rough rolling process)
Next, rough rolling with a cumulative rolling reduction of 30 to 90% is performed. If the cumulative reduction ratio is less than 30%, it becomes difficult to refine by austenite recrystallization. If the cumulative reduction ratio exceeds 90%, the austenite refinement effect is saturated, and the number of passes increases to increase productivity. descend.

(Finish rolling process)
Next, finish rolling is performed in which the sheet thickness center temperature at the start of rolling is 760 ° C. to 880 ° C., the average one-pass rolling reduction is 10 to 25%, and the cumulative rolling reduction is 60 to 85%. The temperature at the end of finish rolling is set to 760 ° C to 840 ° C.

  When the plate thickness center temperature during finish rolling exceeds 880 ° C., the steel slab temperature does not sufficiently reach the non-recrystallized region, the increase in dislocation is suppressed, and the crystal grains cannot be refined. When the sheet thickness center temperature is less than 760 ° C., the productivity decreases, and further processed ferrite is generated in the metal structure. Therefore, the area of the {100} plane forming an angle of 15 ° or less with respect to the plane perpendicular to the rolling direction. It becomes difficult to make the rate 20% or less.

  The reason why the average one-pass reduction ratio during finish rolling is 10 to 25% is as follows. When the average 1-pass rolling reduction is less than 10%, {100} texture develops at the thickness 1/2 part, and forms an angle within 15 ° with respect to a plane perpendicular to the rolling direction of the steel plate {100} It becomes difficult to make the area ratio of the surface 20% or less. Setting the average one-pass rolling reduction to more than 25% is difficult to realize because the burden on the rolling mill becomes extremely large.

  The reason why the cumulative rolling reduction during finish rolling is 60 to 85% is as follows. When the cumulative rolling reduction during finish rolling is less than 60%, it becomes difficult to refine crystal grains and develop a predetermined texture due to the accumulation of dislocations. If the cumulative rolling reduction during finish rolling exceeds 85%, the effect of grain refinement due to the accumulation of dislocations is saturated, and the productivity is further reduced. In order to refine crystal grains, the temperature at the end of finish rolling is set to 840 ° C. or lower.

(Accelerated cooling process)
Subsequent to the above hot rolling, accelerated cooling is performed at a sheet thickness center cooling rate of 15 to 60 ° C./s until the sheet thickness center temperature becomes 740 ° C. or more to 550 ° C. or less.

  If the plate thickness center temperature at the start of accelerated cooling (accelerated cooling start temperature) is less than 740 ° C., the ferrite transformation proceeds and the metal structure becomes coarser, so it is difficult to refine the crystal grains. When the plate thickness center cooling rate of accelerated cooling is less than 15 ° C./s, it becomes difficult to make crystal grains fine. When the plate thickness center cooling rate at the start of accelerated cooling exceeds 60 ° C./s, the ferrite area ratio becomes less than 50%, and it becomes difficult to obtain a predetermined average crystal grain size and texture. If the plate thickness center temperature at the start of accelerated cooling (accelerated cooling stop temperature) exceeds 550 ° C., it becomes difficult to make crystal grains fine. Although it is not necessary to specify the lower limit value of the accelerated cooling stop temperature, it cannot be a temperature lower than the water temperature, so the water temperature or room temperature (for example, about 25 ° C.) is set as the lower limit value. Although it is not necessary to specify the upper limit value of the plate thickness center cooling rate for accelerated cooling, considering the cooling capacity of the apparatus, 70 ° C./s is the upper limit value of the plate thickness center cooling rate for accelerated cooling.

  The steel plate according to the present embodiment can be manufactured by setting the center thickness of the steel slab and the steel plate to be controlled. It is preferable to control the plate thickness center temperature of the steel sheet as the control object rather than the surface temperature of the steel sheet as the control object. By controlling the center temperature of the steel sheet, the manufacturing conditions can be controlled appropriately even when the sheet thickness changes, etc., and high-quality steel sheets with small variations in materials are efficiently manufactured. Can do.

  In the rolling process, usually the temperature distribution inside the steel sheet is calculated while measuring the surface temperature of the steel sheet from heating to rolling, and the rolling control is performed while predicting the rolling reaction force from the calculation result of the temperature distribution. It is carried out. Thus, the steel plate center temperature can be easily obtained during rolling. In the case of performing accelerated cooling, the accelerated cooling is controlled while predicting the temperature distribution inside the plate thickness.

  After performing accelerated cooling, you may temper at 300-650 degreeC as needed.

  When tempering at less than 300 ° C., the effect of tempering is difficult to obtain. When the tempering temperature exceeds 650 ° C., the amount of softening increases, and it becomes difficult to ensure strength and toughness. The tempering temperature is preferably 400 to 600 ° C, more preferably 450 to 550 ° C.

  The manufacturing method of the steel plate which concerns on this embodiment is not limited to the above-mentioned manufacturing method. Even if the manufacturing method of a steel plate is a manufacturing method other than the above, if the content and structure of the steel material are within the specified range, the steel plate is regarded as the steel plate according to the present embodiment.

  The effects of the present invention will be described below based on examples.

  In the steel making process, the chemical composition of the molten steel was adjusted, and then the steel pieces A to AI were produced by continuous casting. Steel slabs A to O are invention steels, and steel slabs P to AI are comparative steels.

  In Examples 1 to 20 and Comparative Examples 21 to 53, the steel pieces A to AI are first reheated, and further the steel pieces A to AI are rolled to form steel plates having a thickness of 10 to 38 mm. The steel plate was water cooled. Thereafter, the steel sheet was heat-treated as necessary.

  Tables 1A and 1B show chemical compositions of the steel pieces A to AI. The underline in Table 1A and Table 1B indicates that the underlined numerical value is outside the specified range of the present invention, and the italic type indicates the analytical value of the amount contained as an impurity.

  Tables 2A to 2D show steel sheet manufacturing methods. The underline in Table 2A to Table 2D indicates that the numerical value with the underline is an unfavorable manufacturing condition. The temperature and the cooling rate in the manufacturing method are values at the center position of the plate thickness. This value was obtained from the measured surface temperature by heat conduction analysis using a known difference method.

  For each steel plate produced, the amount of dissolved oxygen after Al addition, the amount of dissolved oxygen after reflux, the area fraction of each metallographic phase, texture, average crystal grain size, number density of each particle, and mechanical properties were measured. did.

  The amount of dissolved oxygen in the molten steel after the addition of Al and the amount of dissolved oxygen in the molten steel after the reflux are, for example, “rapid analysis of molten steel oxygen in the converter furnace by electromotive force measurement” (Imuta et al., Iron and Steel, Japan Steel) It was calculated | required by the well-known electromotive force measuring method as described in an association, 1972 No. 10, P125-132).

  The area fraction of each metallographic phase (ferrite area ratio, pearlite area ratio, and bainite area ratio, MA area ratio) was first photographed with a magnification of 500 times the metal structure of 1/2 plate thickness with an optical microscope, Next, the total area of each phase in the photographed image was obtained by image analysis, and the total area of each phase was further divided by the area of the whole photographed image.

  The average crystal grain size was determined by the following method. First, by measuring the area of 500 μm × 500 μm with a thickness of 1/2 part by the EBSD method, the measurement position was repeatedly moved by 1 μm. Here, a boundary where the crystal orientation difference between adjacent grains is 15 ° or more is defined as a crystal grain boundary, a circle equivalent diameter of this crystal grain boundary is defined as a crystal grain diameter, and a crystal grain diameter is obtained from the above measurement result. It was. Next, the number distribution of crystal grain sizes was determined, and from this number distribution, the weighted average value of crystal grains was determined, and this was used as the average crystal grain size. The number of divisions N when determining the weighted average value was 20.

  The {100} area ratio indicating the texture is prepared by first creating a map of {100} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction in the 1/2 thickness of the steel sheet. It was determined by dividing the total area of the {100} plane by the measured area.

  Measurement of the number density of each particle (total number density of TiN particles having a particle diameter of 0.5 to 2.0 μm, MnS particles and composite particles, and number density of oxide particles having a particle diameter of 1 to 10 μm) is as follows. The procedure was performed. First, using a scanning electron microscope (SEM: Scanning Electron Microscope), a reflected electron image was photographed in a 1 mm × 1 mm region of ½ part thickness, and particles (particle size) in the reflected electron image were analyzed by image analysis. Various particles including both 0.5-2.0 μm TiN particles, MnS particles and composite particles, and oxide particles with a particle size of 1-10 μm). Next, the composition of each particle was analyzed by EDS, and particles containing 1% by mass or more of Ti, less than 1% by mass of O and 1% by mass or more of N were determined as TiN particles, and 1% by mass or more of Mn And 1% by mass or more of S and less than 1% by mass of O are determined as MnS particles, and a particle that simultaneously satisfies the definition of TiN particles and the definition of MnS particles is determined as composite particles. The particles containing O were determined as oxide particles. Further, by taking a particle image and obtaining the equivalent circle diameter and number of each particle by image analysis, the total number density of TiN particles, MnS particles and composite particles having a particle diameter of 0.5 to 2.0 μm, and The number density of oxide particles having a diameter of 1 to 10 μm was determined.

  Among the mechanical properties, the yield stress and Charpy absorbed energy of the base material were obtained by conducting tests on test pieces taken from the center of the plate thickness and setting the results as representative values of each steel plate.

  The tensile test for determining the yield stress was performed using a No. 1B test piece in accordance with “Metal material tensile test method” of JIS Z 2241 (2011). Two test pieces were collected from one sample for test measurement, and an average value of the two measurement values was obtained. The average value was shown in Tables 3C and 3D as yield stress.

  The Charpy impact test for obtaining Charpy absorbed energy was performed in accordance with “Charpy impact test method for metal material” of JIS Z 2242 (2005) using a 2 mm V notch Charpy impact test piece. Three test pieces were collected from one sample, the test was conducted at a test temperature of −80 ° C., the average value of the three measured absorption energy values was determined, and this average value was used as Charpy absorbed energy. Shown in 3C and Table 3D.

  The test to determine the arrestability of the base material is the NK Classification Society Steel Ship Rules Inspection Procedure K, Appendix K3.12.2-1. (2012) “Inspection Procedure for Brittle Crack Propagation Stop Toughness Value Kca Test Method” was performed using an ESSO test piece having an original thickness and a sheet width of 500 mm. By the test, the arrest toughness value Kca at −60 ° C. was determined.

  Joint toughness was determined as follows. First, welding with a submerged arc with a welding heat input of 58 to 216 kJ / cm, or electrogas arc welding, was tested. 1-53 butt welded joints were made on steel plates 1 to 53. Next, 30 Charpy impact test pieces were collected from each weld joint. At this time, a notch of a 2 mmV notch Charpy impact test piece was formed along the melting line (FL) at a thickness of 1/4 part. A Charpy impact test was performed on these test pieces, and an average value of absorbed energy at −60 ° C. was obtained. Furthermore, statistical analysis was performed assuming that the distribution of absorbed energy values follows a normal distribution. The probability that the absorbed energy of steel plates 1 to 53 was 50 J or less was determined. The Charpy impact test was based on “Charpy impact test method for metal materials” of JIS Z 2242 (2005).

These measurement results for the steel plates of Examples 1 to 20 and Comparative Examples 21 to 53 are shown in Tables 3A to 3D. When the average Charpy absorbed energy at −80 ° C. of the base material was 100 J or more, it was judged that the base material toughness was good. When Kca of the base material was 4000 N / mm ^1.5 or more, it was judged that the base material arrestability was good. When the average Charpy absorbed energy at −60 ° C. of the heat-affected zone of the welded joint was 100 J or more and the probability that the Charpy absorbed energy was 50 J or less was 10% or less, the joint toughness was judged to be good. Underlined in Tables 3A to 3D indicates that the numerical value with the underline is outside the specified range of the present invention, or is rejected in view of the above pass / fail criteria.

  Since Examples 1 to 20 satisfied all the conditions of the present invention, the base material strength, base material toughness, base material arrestability, and joint toughness were all good.

  In Comparative Examples 21 to 53, the numerical value underlined was out of the specified range of the present invention, so that good results were not obtained. This will be specifically described below.

  In Comparative Examples 21 to 40, since the chemical components deviated from the specified range of the present invention, there was a problem in at least one of the base material strength, base material toughness, base material arrestability, and joint toughness.

  In Comparative Example 41, since the atmospheric temperature of the heating furnace was too high, the average crystal grain size of the base material was large, and the base material toughness and the base material arrestability were low. In Comparative Example 41, the joint number toughness was low because the particle number density of the TiN particles, MnS particles, and their composite particles was small.

  In Comparative Example 42, since the soaking temperature in the heating furnace was too high, the average crystal grain size of the base material was large, and the base material toughness and base material arrestability were low.

  In Comparative Example 43, since both the atmosphere temperature of the heating furnace and the soaking temperature were too high, the ferrite fraction of the base material was low, the bainite fraction was high, the average crystal grain size of the base material and the base material The {100} area ratio increased. Thereby, the base material toughness and the base material arrestability were low. In Comparative Example 43, the joint number toughness was low because the particle number density of the TiN particles, MnS particles, and their composite particles was small.

  In Comparative Example 44, since the soaking time in the heating furnace was too long, the average crystal grain size of the base material was large, and the base material toughness and base material arrestability were low. Further, since the particle number density of TiN particles, MnS particles, and composite particles thereof was small, joint toughness was low.

  In Comparative Example 45, since the cumulative rolling reduction of rough rolling was too small, the ferrite fraction of the base material was low, the bainite fraction of the base material was high, the average crystal grain size of the base material, and the {100} area of the base material The rate was great. Thereby, the base material toughness and the base material arrestability were low.

  In Comparative Example 46, since the finish rolling temperature was too high, the average crystal grain size of the base material was large, and the base material toughness and the base material arrestability were low.

  In Comparative Example 47, since the finish rolling temperature was too low, the {100} area ratio of the base material was large, and thus the base material toughness and the base material arrestability were low.

  In Comparative Example 48, since the average rolling reduction ratio for finish rolling was too small, the {100} area ratio of the base material was large, and thereby the base material arrestability was low.

  In Comparative Example 49, since the cumulative rolling reduction of finish rolling was too small, the average crystal grain size of the base material was large, and thereby the base material toughness and the base material arrestability were low.

  In Comparative Example 50, since the plate thickness was too large, the average crystal grain size of the base material was large, and thus the base material toughness and the base material arrestability were low. Further, since the plate thickness was large, the welding heat input was too large, so the joint toughness was low.

  In Comparative Example 51, since the cooling start temperature was too low, the ferrite fraction of the base material was large, the bainite fraction of the base material was low, and the average crystal grain size of the base material was large, thereby increasing the base material strength and the base material. Toughness and base metal arrestability were low.

  In Comparative Example 52, since the cooling rate was too low, the ferrite fraction of the base material was large, the bainite fraction of the base material was low, and the average crystal grain size of the base material was large, thereby the base material strength and base material toughness. And, the base metal arrestability was low.

  In Comparative Example 53, since the cooling stop temperature was too high, the ferrite fraction of the base material was large, the bainite fraction of the base material was low, and the average crystal grain size of the base material was large, thereby increasing the base material strength and the base material. Toughness and base metal arrestability were low.

  From the above examples, by applying the present invention, a high-strength steel sheet excellent in arrestability with low manufacturing cost, high productivity, high strength, thick plate thickness, and no deterioration in joint toughness. It was confirmed that it could be provided.

  In addition, this invention is not limited to embodiment mentioned above. Various modifications can be made without departing from the spirit of the present invention.

  According to the present invention, a steel plate having a low manufacturing cost and high productivity, a toughness of a weld heat-affected zone and a brittle crack propagation stopping performance of a base material, and a welded joint having an excellent toughness of a weld heat-affected zone. Can be provided.

Claims (13)

Chemical composition is mass%,
C: 0.040-0.090%,
Si: 0.01-0.20%,
Mn: 1.30 to 1.80%,
P: 0.020% or less,
S: 0.001 to 0.010%,
Al: 0.005 to 0.100%,
Nb: 0.003-0.030%,
Ti: 0.003-0.030%,
B: 0.0003 to 0.0040%,
N: 0.0020 to 0.0080%,
O: 0.0005 to 0.0040%,
Cu: 0 to 1.00%,
Ni: 0 to 1.00%,
Cr: 0 to 1.00%,
Mo: 0 to 0.500%,
V: 0 to 0.100%,
Ca: 0 to 0.0050%,
Mg: 0 to 0.0050%,
REM: 0-0.0050%, and the balance: iron and impurities,
The carbon equivalent CE defined by the formula A is 0.30 to 0.40% by mass,
SOLB defined by the formula B is -0.0015 to +0.0015 mass%,
The metal structure is a mixed structure containing ferrite and bainite, or a mixed structure containing the ferrite, pearlite and the bainite,
Area ratio 50-90% of the ferrite, the area ratio is 10-50% of the previous SL bainite, and the area ratio of the pearlite 0-10%,
The area ratio of MA in the metal structure is 0 to 5%,
In the plate thickness 1/2 part, the boundary between adjacent crystals having a crystal orientation difference of 15 ° or more is defined as a crystal grain boundary, and a region surrounded by the crystal grain boundary is defined as a crystal grain. When the division number N is set to an integer of 10 or more, the weighted average value D _AVE of the grain sizes of the crystal grains defined by the E formula is 3.0 to 17.0 μm,
In the plate thickness ½ part, the area ratio of the {100} plane forming an angle within 15 ° with respect to the plane perpendicular to the rolling direction is 2.0 to 20.0%.
In a thickness of 1/4 part, particles containing 1% by mass or more of Ti, less than 1% by mass of O and 1% by mass or more of N are defined as TiN particles, and 1% by mass or more of Mn and 1% by mass or more. When particles containing S and less than 1% by mass of O are defined as MnS particles, and particles satisfying both the definition of TiN particles and the definition of MnS particles are defined as composite particles, 0.5-2 The total number density of the TiN particles, the MnS particles, and the composite particles each having an equivalent circle diameter of 0.0 μm is 20 to 200 particles / mm ² ;
In the case where the plate thickness is 1/4 part, the number density of the oxide particles having a circle-equivalent diameter of 1 to 10 μm is 20 to 200 / particles when particles containing 1 mass% or more of O are defined as oxide particles. mm ²
The plate thickness is 10 to 35 mm,
A steel sheet characterized by a yield stress of 300 to 500 MPa.
CE = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 (A)
SOLB = 0.226 × Ti + B−0.772 × N (B)
p = (D _MAX −D _MIN ) / N (C)
D _k = D _MIN + p × (k−1 / 2) (D)
D _AVE = (Σ [k = 1, N] (D _k × S _k )) / (Σ [k = 1, N] S _k ) (E)
The element symbol described in the formula indicates the content of each component in the steel sheet in unit mass%, D _MAX indicates the particle size of the largest crystal grain in unit μm, and D _MIN is The grain size of the smallest crystal grain is expressed in μm, k is an integer of 1 or more and N or less, and S _k is (D _MIN + p × (k−1)) μm or more (D _MIN + p × k) The total area ratio of the crystal grains having a circle-equivalent diameter of less than μm is shown in unit%.   The steel sheet according to claim 1, wherein the chemical composition is Si: 0.01 to 0.10% in terms of mass%.   The said chemical composition is the mass% and is Al: 0.015-0.060%, The steel plate of Claim 1 or 2 characterized by the above-mentioned.   The said chemical composition is the mass% and is Mo: 0-0.010%, The steel plate of any one of Claims 1-3 characterized by the above-mentioned.   The said chemical composition is the mass% and is Ti: 0.005-0.018%, The steel plate of any one of Claims 1-4 characterized by the above-mentioned.   The said chemical composition is the mass% and is B: 0.0005-0.0020%, The steel plate of any one of Claims 1-5 characterized by the above-mentioned.   The said chemical composition is the mass% and is N: 0.0025-0.0060%, The steel plate of any one of Claims 1-6 characterized by the above-mentioned.   The said chemical composition is the mass% and is O: 0.0010 to 0.0030%, The steel plate of any one of Claims 1-7 characterized by the above-mentioned.   The steel plate according to any one of claims 1 to 8, wherein the SOLB is -0.0010 to +0.0005 mass%. The steel plate according to any one of claims 1 to 9, wherein the weighted average value D _{AVE of the} crystal grains is 3.0 to 13.0 µm. The total number of the number densities of the TiN particles, the MnS particles, and the composite particles each having an equivalent circle diameter of 0.5 to 2.0 μm is 50 to 140 particles / mm ^2. The steel plate according to any one of 1 to 10. The steel sheet according to any one of claims 1 to 11, wherein the number density of the oxide particles having an equivalent circle diameter of 1 to 10 µm is 20 to 150 particles / mm ² . The chemical composition is mass%,
Cu: 0.10 to 1.00%,
Ni: 0.10 to 1.00%,
Cr: 0 to 0.10%,
Mo: 0 to 0.100%,
V: 0 to 0.005%
The steel sheet according to any one of claims 1 to 12, wherein the steel sheet is a steel sheet.