JP5445720B1 - High strength steel plate with excellent arrestability - Google Patents

Abstract

This high strength thick steel plate has a carbon equivalent Ceq. Has a component composition of 0.3 to 0.5%, has a microstructure containing 70% or less of ferrite in area ratio, and bainite of 30% or more in area ratio, and is 1/4 of the plate thickness. Part, the crystal grain boundary density, which is the total length per unit area of the crystal grain boundary having a crystal orientation difference of 15 ° or more, is 400 to 1000 mm / mm ² , and 15 ° to the plane perpendicular to the main rolling direction. The area ratio of the {100} plane forming an angle of within 10 to 40%, and at 1/2 part of the plate thickness, the grain boundary density is 300 to 900 mm / mm ² and the main rolling direction The area ratio of the {110} plane forming an angle of 15 ° or less with respect to the plane perpendicular to the plane is 40 to 70%.

Description

The present invention relates to a high-strength thick steel plate excellent in arrestability.
The present application claims priority based on Japanese Patent Application No. 2012-087384 filed in Japan on April 6, 2012, the contents of which are incorporated herein by reference.

  For steel plates used in structures such as shipbuilding, architecture, tanks, offshore structures, line pipes, arrestability (the ability to suppress the propagation of brittle fracture in order to suppress brittle fracture of structures) Brittle fracture propagation stopping performance) is required. In recent years, with the increase in size of structures, there are increasing cases of using high-strength thick steel plates having a yield stress of 390 to 690 MPa and a plate thickness of 60 to 95 mm. However, the above-described arrestability generally tends to conflict with strength and thickness. For this reason, the technique which improves arrestability in a high intensity | strength thick steel plate is desired.

  As a method for improving the arrestability, for example, a method of controlling the crystal grain size, a method of controlling the embrittled second phase, and a method of controlling the texture are known.

As a method for controlling the crystal grain size, there are techniques described in Patent Documents 1 to 3 and 21.
In the technique described in Patent Document 1, the arrestability is improved by using ferrite as a parent phase and making the ferrite finer. In order to obtain such fine-grained ferrite, cooling is performed so that 1/8 or more of the slab thickness becomes Ar3 or less from the front and back layer portions toward the center of the slab thickness, and rolling is performed in a cryogenic region, and then Ac3 is used. It is necessary to reheat to a temperature exceeding that and recrystallize the ferrite.
In the techniques described in Patent Documents 2 and 3, fine ferrite recrystallized grains are formed by rolling in the process where the surface layer part is once cooled to Ar1 or less and then the surface layer part is reheated, using ferrite as a parent phase. obtain.
In the technique described in Patent Document 21, the average crystal grain size in the major axis direction of ferrite is 5 μm or more and the aspect ratio is 2 or more, or the average grain size in the major axis direction of prior austenite grains is 10 μm or more and the aspect ratio is By setting it to 2 or more, the brittle crack propagation stopping property is enhanced.

As a method for controlling the embrittled second phase, there is a technique described in Patent Document 4.
In the technique described in Patent Document 4, a fine embrittled second phase (for example, martensite) is dispersed in a ferrite as a parent phase, whereby a microcrack is formed in the embrittled second phase at the brittle crack tip. To relieve the stress state at the crack tip.

  Moreover, there exists a technique described in patent documents 5-17 as a method of controlling a texture. In the techniques described in Patent Documents 5 to 17, by controlling the X-ray plane intensity ratio as a texture at each plate thickness position, for example, the surface layer portion, 1/4 portion of the plate thickness, and 1/2 portion of the plate thickness. , Change the crack propagation direction, improve arrestability.

Furthermore, as a method for controlling both the crystal grain size and the texture, there are techniques described in Patent Documents 18 to 20.
In the technique described in Patent Document 18, the arrestability is improved by controlling the crystal grain size and the X-ray plane intensity ratio by setting the ferrite fraction at ½ part of the plate thickness to 80% or more.
In the technique described in Patent Document 19, the arrestability is improved by controlling the texture size ratio measured by X-rays and the crystal grain size of ½ part of the surface layer and the plate thickness.
In the technique described in Patent Document 20, the arrestability is improved by controlling the crystal grain size of the surface layer, ½ part of the plate thickness, and the area ratio of the {100} plane perpendicular to the external stress.

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. 2008-156751

  In the techniques described in Patent Documents 1 to 3, since ferrite is mainly used by utilizing recrystallization of ferrite in the front and back layers of the steel sheet, it is difficult to obtain a steel sheet having high strength and a large thickness. . In addition, as in the techniques described in Patent Documents 1 to 3 and 21, it is difficult to improve the arrestability with a steel plate having high strength and a large plate thickness only by controlling the crystal grain size. In addition, it is necessary to go through cooling, rolling, and recuperation steps, and the manufacturing process becomes complicated. Therefore, it is extremely difficult to obtain a stable material. Furthermore, in such a manufacturing process, shape defects are likely to occur due to uneven cooling of the steel sheet surface. When a shape defect occurs, a great deal of cost is required for shape correction.

  In the technique described in Patent Document 4, since martensite is dispersed in ferrite, brittle crack generation characteristics are significantly deteriorated. Furthermore, since the main component is ferrite, it is difficult to obtain a steel plate having high strength and a large plate thickness as described above. Moreover, it is difficult to improve the arrestability with a steel plate having a high strength and a thick plate thickness only by controlling the embrittlement second phase.

In the techniques described in Patent Documents 5 to 17, 19 and 21, the crystal grain size, which is the most effective factor, is not controlled in order to improve the arrestability of a steel plate having high strength and a large thickness. That is, only by controlling the texture, the arrestability cannot be dramatically improved with a steel plate having a high strength and a large thickness. Moreover, since the X-ray plane intensity ratio represents a local texture, the variation is large. These techniques are not techniques for improving arrestability and obtaining high productivity during hot rolling. In the first place, the techniques of Patent Documents 5 to 8, 11 and 21 are techniques for improving the brittle crack propagation stopping characteristics in the plate thickness direction, such as a direction parallel to the steel sheet surface as in the present application, for example, perpendicular or parallel to the rolling direction. It is not a technique related to improving the direction of brittle crack propagation stoppage. With such a technique, it is impossible to improve the brittle crack propagation stopping characteristics in the direction parallel to the steel plate surface.
Even if the arrestability of the steel sheets disclosed in Patent Documents 9 and 10 is the highest, the Kca at −10 ° C. with a plate thickness of 60 mm is about 6500 to 6600 N · mm ^−0.5 . This level is considered that Kca at −20 ° C. is 6000 N · mm ^−0.5 or less, and further improvement in arrestability is necessary.
Patent Documents 12, 13, 16, and 19 disclose techniques for obtaining high arrestability. However, in order to form a peculiar texture in the plate thickness direction, the cumulative reduction rate is 30% or more and the average pass pressure is lowered in the temperature range where the temperature at the center of the plate thickness is Ar3 −10 ° C. or lower and Ar3 −50 ° C. Rolling with a rate of 8% or more is necessary. That is, rolling at a very low temperature is indispensable, productivity at the time of rolling is very low, and mass production is difficult.
The arrestability in the plate thickness direction of the steel sheet disclosed in Patent Documents 14, 15 and 17 is the large-sized hybrid ESSO test (running plate length: 1600 mm, test plate length: 800 mm, specimen width: 2400 mm, load stress: 235 kg. ^{-Mm -0.5} ). The Kca at −20 ° C. is considered to be 6000 N · mm ^−0.5 or less. Moreover, in order to form a texture, rolling at a low temperature is indispensable, and mass production is difficult.

In the technique described in Patent Document 18, since the crystal grain size and texture are controlled by only ½ part of the plate thickness, it is difficult to improve the arrestability when the plate thickness is thick. Moreover, since ferrite is the main component, it is difficult to obtain a steel plate having high strength and a large plate thickness. Further, since the X-ray plane intensity ratio represents a local texture, it has a large variation and is not suitable as a factor for improving arrestability. Further, this technique is a technique for improving the arrestability at an angle of 45 ° with respect to the rolling direction by forming a texture by rolling at a low temperature. Rolling at a low temperature is indispensable, and mass production is difficult.
In the technique described in Patent Document 20, the crystal grain size and texture in the surface layer and ½ part of the plate thickness are controlled. However, in improving the arrestability when the plate thickness is thick, the contribution of the surface layer, which is in the state of plane stress and hardly cleaves by nature, is extremely small, so the control of the surface layer can dramatically improve the arrestability. It is difficult. Further, it is disclosed that Kca at −10 ° C. with a plate thickness of 70 mm is 210 MPa · mm ^−0.5 (that is, equivalent to about 6600 N · mm ^−0.5 ). This level is considered to be 6000 N · mm ^−0.5 or less in Kca at −20 ° C., and further improvement in arrestability is necessary.

  The present invention has been made in consideration of the above-mentioned circumstances, and the purpose thereof is an arrest that has low manufacturing cost, high productivity, high strength, thick plate thickness, and no degradation of HAZ toughness. An object of the present invention is to provide a high-strength thick steel plate having excellent properties.

  The gist of the present invention is as follows.

(1) The high-strength thick steel plate according to one embodiment of the present invention is mass%, C: 0.04 to 0.16%, Si: 0.01 to 0.5%, Mn: 0.75 to 2. 5%, Al: 0.001-0.1%, Nb: 0.003-0.05%, Ti: 0.003-0.05%, N: 0.001-0.008%, P is 0.03% or less, S is 0.02% or less, Cu is 1% or less, Ni is 2% or less, Cr is 1% or less, Mo is 0.5% or less, V is 0.15% or less, B is 0.005% or less, Ca is 0.01% or less, Mg is 0.01% or less, REM is limited to 0.01% or less, and the balance contains iron and unavoidable impurities. Carbon equivalent Ceq. Has a component composition of 0.30 to 0.50%, and has a microstructure containing 70% or less ferrite in area ratio and 30% or more bainite in area ratio, and is ¼ of the plate thickness. Part, the crystal grain boundary density, which is the total length per unit area of the crystal grain boundary having a crystal orientation difference of 15 ° or more, is 400 to 1000 mm / mm ² , and 15 ° to the plane perpendicular to the main rolling direction. The area ratio of the {100} plane forming an angle of within 10 to 40%, and at 1/2 part of the plate thickness, the grain boundary density is 300 to 900 mm / mm ² and the main rolling direction The area ratio of the {110} plane forming an angle of 15 ° or less with respect to the plane perpendicular to the plane is 40 to 70%.
Ceq. = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5
... Formula A (2) The high-strength thick steel plate described in (1) above may have a thickness of 60 to 95 mm.
(3) The high-strength thick steel plate described in (1) or (2) above may have a yield stress of 390 to 690 MPa.
(4) In the high-strength thick steel plate according to any one of (1) to (3), the microstructure may contain pearlite having an area ratio of 10% or less.
(5) In the high-strength thick steel plate according to any one of (1) to (4), the microstructure has a ferrite area ratio of less than 50%, a pearlite area ratio of 5% or less, and a bainite area. The rate may be 50% or more.
(6) In the high-strength thick steel plate according to any one of the above (1) to (5), the grain boundary density of the ¼ part of the plate thickness is 500 to 900 mm / mm ² , and the plate thickness 1 The crystal grain boundary density of / 2 parts may be 400 to 800 mm / mm ² .
(7) In the high-strength thick steel sheet according to any one of (1) to (6), the Cu is 0.5% or less, the Ni is 1% or less, and the Cr is 0.5% or less. The Mo may be further limited to 0.2% or less, and the V may be further limited to 0.07% or less.
(8) In the high-strength thick steel plate according to any one of (1) to (7), B may be further limited to 0.002% or less.
(9) In the high-strength thick steel plate according to any one of (1) to (8), the Ca is 0.003% or less, the Mg is 0.003% or less, and the REM is 0.003%. You may further restrict | limit to the following.

  According to the present invention, a steel sheet that is extremely excellent in arrestability in a direction parallel to the surface of the steel sheet, for example, a direction perpendicular to or parallel to the rolling direction, has high strength even when the plate thickness is thick, and has no degradation of HAZ toughness Therefore, it is possible to reduce the cost and improve the safety of the welded steel structure.

It is a photograph which shows the aspect of the crack propagation which generate | occur | produced and applied the impact from the V notch of the photograph left direction with respect to the steel plate which concerns on one Embodiment of this invention. It is a photograph of the fracture surface of the crack shown to FIG. 1A. It is a photograph which shows the aspect of the crack propagation which generate | occur | produced and applied the impact from the V notch of the photograph left direction with respect to the steel plate which concerns on a comparative example. It is a photograph of the fracture surface of the crack shown to FIG. 2A. It is a photograph which shows the aspect of the crack propagation which generate | occur | produced and applied the impact from the V notch of the photograph left direction with respect to the steel plate which concerns on a comparative example. It is a photograph of the fracture surface of the crack shown to FIG. 3A.

  The present inventors have intensively studied to solve the above problems, and as a result, by controlling the component composition, microstructure, grain boundary density in the plate thickness direction, and texture in the plate thickness direction of the high-strength steel plate. The present inventors have found that a high-strength steel sheet having high productivity during hot rolling and improved arrestability in a direction parallel to the steel sheet surface, for example, a direction perpendicular to or parallel to the rolling direction can be obtained.

Hereinafter, a high-strength thick steel plate (hereinafter sometimes simply referred to as a steel plate) according to an embodiment of the present invention made based on the above-described knowledge will be described.
The steel sheet according to this embodiment controls the component composition, the microstructure, the grain boundary density in the plate thickness direction, and the texture in the plate thickness direction, thereby controlling the direction parallel to the steel plate surface, for example, the rolling direction. Improve the arrestability in the direction perpendicular to or parallel to.

(Micro structure)
The steel sheet according to the present embodiment is a mixed structure of ferrite and bainite, or a mixed structure of ferrite, pearlite, and bainite, and has a microstructure with a ferrite area ratio of 70% or less and a bainite area ratio of 30% or more.
If the ferrite area ratio is more than 70%, it is difficult to obtain a steel plate having a large plate thickness and high strength. If a steel plate having a desired thickness and strength can be obtained, bainite, pearlite and bainite can be used as the second phase. The present invention is intended for thick high-strength steel, and the upper limit of the ferrite area ratio may be limited to less than 50%, less than 30%, less than 20%, or less than 10%.
If the bainite area ratio is less than 30%, it is difficult to obtain a steel sheet having a large thickness and high strength. The upper limit of the bainite area ratio may be 95% in order to secure the ferrite area ratio and increase the crystal grain boundary that hinders brittle crack propagation. The present invention is intended for thick high-strength steel, and the lower limit of the bainite area ratio may be limited to 50% or more, 60% or more, 70% or more, or 80% or more.
Perlite may be contained as long as a steel plate having a desired thickness and strength can be obtained. Therefore, the pearlite area ratio may be limited to 10% or less, 5% or less, or 3% or less. The lower limit of the pearlite area ratio is 0%.
In addition to ferrite, pearlite, and bainite, fine island martensite (MA) may exist, but the MA area ratio is 5% or less. The MA area ratio may be limited to 3% or less, 2% or less, or 1% or less, and 0% is most desirable.

(Grain boundary density in the plate thickness direction)
The dominant factor in improving arrestability is the largest contribution of grain boundaries. This is because the crystal grain boundary becomes an obstacle to brittle crack propagation. That is, since the crystal orientation differs between adjacent crystal grains at the crystal grain boundary, the direction in which the crack propagates changes in this portion. For this reason, an unbroken region is generated, and the stress is dispersed by the unbroken region, resulting in a crack closing stress. Therefore, the driving force for crack propagation is reduced and the arrestability is improved. Further, since the unbroken region finally undergoes ductile fracture, energy required for brittle fracture is absorbed. For this reason, arrestability improves.

  Until now, it was considered necessary to reduce the crystal grain size in order to increase the grain boundary. This is true in the structure mainly composed of ferrite, but the use of bainite is indispensable for steel with a large thickness and high strength. Unlike ferrite, this bainite has a complicated substructure, so it is very difficult to define crystal grains. For this reason, even if the relationship between the crystal grain size and the arrestability is calculated in terms of the equivalent circle diameter, the variation is large, and it is difficult to determine the crystal grain size necessary for improving the arrestability. Therefore, returning to the basic principle that the grain boundary becomes an obstacle to crack propagation, the total length of the grain boundary per unit area (hereinafter referred to as the grain boundary density) is defined and used to relate to the arrest property. We found that the best correlation was obtained by organizing

Therefore, in the steel sheet according to this embodiment,
(A) The grain boundary density is 400 to 1000 mm / mm ² at ¼ part of the plate thickness,
(B) The grain boundary density is set to 300 to 900 mm / mm ² at ½ part of the plate thickness.
Here, “crystal grain boundary density” means “total length per unit area of crystal grain boundaries having a crystal orientation difference of 15 ° or more”. The reason why the crystal orientation difference is set to 15 ° or more is that if it is less than 15 °, the crystal grain boundary is unlikely to be an obstacle to brittle crack propagation, and the effect of improving arrestability is reduced.

That is, the arrest toughness value (Kca) at −20 ° C. is 6000 N · when the grain boundary density satisfies the requirements of 400 and 300 mm / mm ² or more at ¼ part and ½ part of the plate thickness, respectively. High arrestability of mm ^−0.5 or more is shown. In order to further improve the arrestability, the grain boundary density is preferably set to 500 or 400 mm / mm ² or more at 1/4 part or 1/2 part of the plate thickness, or 600 or 400 mm, respectively. / Mm ² or more is more preferable.

The arrestability improves as the grain boundary density increases, but excessively increasing the rolling load increases the productivity, so the upper limit of the grain boundary density is 1/4 of the plate thickness. Part and 1/2 part are 1000 and 900 mm / mm ² , respectively. Each limit may be respectively limited 900,800mm / mm ² or, respectively 800,700mm / mm ^2.

  The reason why the grain boundary density is defined by 1/4 part and 1/2 part of the plate thickness is that it is necessary to increase the grain boundary density of the entire plate thickness in order to improve the arrestability of the extra-thick material. This is because by controlling ¼ part and ½ part of the thickness, it is possible to obtain a representative value of the average grain thickness of the grain boundary. In addition, according to the manufacturing method described later that mainly controls the grain boundary density of 1/2 part of the plate thickness, the other plate thickness positions inevitably have a low temperature and a high cooling rate. Since the field density tends to increase, it is not necessary to limit the special value. However, depending on the heating method, a large temperature gradient may occur in the thickness direction, and the grain boundary density at 1/4 and 1/2 parts of the thickness may be reversed. ing.

For the measurement of the crystal grain boundary, 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. By using the EBSD method, it is possible to identify a crystal grain boundary having a complicated structure such as bainite.
More specifically, the crystal grain boundary density was measured by measuring the area of 500 μm × 500 μm at ¼ part and ½ part of the plate thickness at 1 μm pitch by EBSD method, and the crystal orientation difference between adjacent grains was 15 It can be determined by defining a boundary of more than 0 ° as a grain boundary and dividing the total length of the grain boundary at that time by the measurement area.

(Texture in the thickness direction)
In the case of a high-strength steel plate having a large plate thickness, it is difficult to stably improve the arrestability only by controlling the grain boundary density in the plate thickness direction. Therefore, it is important to control the crack propagation direction using the texture. A brittle crack generated in the steel sheet when the steel sheet is subjected to external stress propagates along a {100} face cleavage plane. Therefore, it has been found that if the {100} plane texture develops in a plane perpendicular to the external stress, the effect of improving the arrestability when the crystal grain size is controlled as described above is reduced. External stress is a stress externally applied to a steel structure. Brittle cracks often occur and propagate in a direction perpendicular to the highest external stress. Therefore, here, the highest stress externally applied 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 surface perpendicular to the external stress can be handled as a surface perpendicular to the main rolling direction of the steel sheet.
The main rolling direction of the steel sheet can be specified by, for example, corroding the steel sheet surface with picric acid and measuring the aspect ratio of prior austenite. That is, the direction in which the aspect ratio of the prior austenite is large can be specified as the main rolling direction of the steel sheet.

  If the texture of the {110} plane forming an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction of the steel sheet is 40 to 70% in terms of the area ratio at 1/2 part of the plate thickness, It has been found that the driving force of crack propagation can be reduced by the fact that the brittle crack in the vicinity of ½ part does not propagate straight but propagates in an inclined manner. However, if a similar texture is developed in plate thickness parts other than 1/2 part of the plate thickness, the crack propagates in an inclined state, and a sufficient arrestability improving effect cannot be exhibited. Therefore, conversely, in order to propagate the crack straight at ¼ part of the plate thickness, an angle of 15 ° or less is formed with respect to a plane perpendicular to the main rolling direction of the steel plate at ¼ part of the plate thickness { By setting the texture of the 100} plane to 10 to 40% in terms of area ratio, it is possible to suppress the propagation of a ½ part inclined crack to a plate thickness part other than ½ part. found.

Based on the above knowledge, in the steel sheet according to the present embodiment,
(C) In the ¼ part of the plate thickness, the area ratio of the {100} plane that forms an angle within 15 ° with respect to the plane perpendicular to the main rolling direction is 10 to 40%,
(D) In the ½ part of the plate thickness, the area ratio of the {110} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction is 40 to 70%.

By satisfying the above (C) and (D), as shown in FIGS. 1A and 1B, the ½ part crack propagates in an inclined manner, and the ¼ part crack becomes straight. Propagation of the cracks is further increased. Thereby, the arrestability improvement effect by the increase in the crystal grain boundary density can be exhibited, and the arrestability can exhibit a sufficient value. FIG. 1A is a photograph showing a mode of crack propagation generated by applying an impact from a V-notch in the left direction of the steel plate according to one embodiment of the present invention, and FIG. 1B is a fracture view of the crack. It is a photograph which shows a cross section.

The reason why the area ratio of the {100} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction at ¼ part of the plate thickness is 10% or more is that the crack is straightened if it is less than 10%. This is because the effect of propagation cannot be obtained.
In addition, the reason why the area ratio of the {100} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction at ¼ part of the plate thickness is 40% or less is shown in FIGS. 2A and 2B. Thus, when it exceeds 40%, the crack propagation of ¼ part is more dominant than ½ part, and the arrest property is deteriorated by the propagation of the crack straight. Note that FIG. 2A is a photograph for a steel sheet with an area ratio of the {100} plane exceeding 40% at an angle of 15 ° or less with respect to a plane perpendicular to the main rolling direction at ¼ part of the plate thickness. FIG. 2B is a photograph showing an aspect of crack propagation generated by applying an impact from the left V-notch, and FIG. 2B is a photograph showing a fracture surface of the crack.

The area ratio of the {100} plane, which forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction, at ¼ part of the plate thickness is preferably 13 to 37%, more preferably 15 to 15%. 35%.

The reason why the area ratio of the {110} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction in the ½ part of the plate thickness is 40% or more is that the crack is inclined if it is less than 40%. This is because the effect of propagating it is not obtained.
The reason why the area ratio of the {110} plane that forms an angle within 15 ° with respect to the plane perpendicular to the main rolling direction at ½ part of the plate thickness is 70% or less is shown in FIGS. 3A and 3B. As shown in the figure, when it exceeds 70%, the arrestability is deteriorated by propagating while being inclined without receiving a resistance of ¼ part. Note that FIG. 3A is a photograph for a steel sheet in which the area ratio of the {110} plane, which forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction, is more than 70% at 1/2 part of the plate thickness. FIG. 3B is a photograph showing an aspect of crack propagation generated by applying an impact from the left V-notch, and FIG. 3B is a photograph showing a fracture surface of the crack.

In the ½ part of the plate thickness, the area ratio of the {110} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction is preferably 45 to 65%, and more preferably 40%. ~ 60%.

The texture 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, according to the EBSD method, at {fraction (1/4)} of the plate thickness, the {100} plane forming an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction of the steel plate, By creating maps of {110} planes and dividing the total area by the measured area, the area ratio can be obtained.

The measures for improving the arrestability as described above can be applied to a steel plate having a yield stress of 390 to 690 MPa, a tensile strength of 500 to 780 MPa, and a steel plate having a thickness of 60 to 95 mm. This is because in a region where the yield stress is less than 390 MPa or the plate thickness is less than 60 mm, it is relatively easy to improve the arrestability without relying on the means of the present invention, the yield stress exceeds 690 MPa, the plate thickness is In the region exceeding 95 mm, even if the grain boundary density and texture defined in the present invention are formed, the mechanical condition becomes severe, so the arrest toughness value (Kca) at −20 ° C. is 6000 N · mm ^−0.5. This is because it is difficult to achieve the above high arrestability. The lower limit of the yield stress may be limited to 440 MPa or 470 MPa, and the upper limit may be limited to 640 MPa or 590 MPa. The lower limit of the tensile strength may be limited to 520 MPa, 540 MPa, or 560 MPa, and the upper limit may be limited to 730 MPa, 680 MPa, or 630 MPa.

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

C: 0.04 to 0.16%
C is contained in an amount of 0.04% or more in order to ensure the strength and toughness of the thick base material. If the C content exceeds 0.16%, it becomes difficult to ensure good HAZ toughness, so the C content is set to 0.16% or less.
Therefore, the lower limit value of C is 0.04%, preferably 0.05%, more preferably 0.06%, and the upper limit value of C is 0.16%, preferably 0.14%, more preferably 0. .12%.

Si: 0.01 to 0.5%
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 0.5%, the HAZ toughness is greatly deteriorated, so the Si content is 0.5% or less.
Accordingly, the lower limit of Si is 0.01%, preferably 0.03%, more preferably 0.05%, and the upper limit of Si is 0.5%, preferably 0.4%, more preferably 0. .35% or 0.3%.

Mn: 0.75 to 2.5%
Mn is contained in an amount of 0.75% or more in order to economically secure the strength and toughness of the thick base material. If the Mn content exceeds 2.5%, the center segregation becomes prominent, and the toughness of the base material and the HAZ where the center segregation has occurred deteriorates, so the Mn content is 2.5% or less. .
Therefore, the lower limit value of Mn is 0.75%, preferably 0.9%, more preferably 1.2%, and the upper limit value of Mn is 2.5%, preferably 2.0%, more preferably 1%. .8% or 1.6%.

P: Limited to 0.03% or less P is one of the impurity elements. In order to stably secure the HAZ toughness, the P content may be limited to 0.03% or less. Preferably, it is 0.02% or less, more preferably 0.015% or less. The lower limit value is 0%, but considering the cost for reducing the P content, 0.0001% may be set as the lower limit value.

S: Limited to 0.02% or less S is one of the impurity elements. In order to stably ensure the characteristics of the base material and the HAZ toughness, the S content may be limited to 0.02% or less. Preferably, it is 0.01% or less, More preferably, it is 0.008% or less. The lower limit is 0%, but 0.0001% may be set as the lower limit in consideration of the cost for reducing the S content.

Al: 0.001 to 0.1%
Al is responsible for deoxidation and reduces O, which is one of the impurity elements. In addition to Al, Mn and Si also contribute to deoxidation. However, even when Mn or Si is added, if the Al content is less than 0.001%, O cannot be stably reduced. However, if the Al content exceeds 0.1%, alumina-based coarse oxides and clusters thereof are generated, and the base material and the HAZ toughness are impaired. Therefore, the Al content is 0.1% or less. .
Accordingly, the lower limit of Al is 0.001%, preferably 0.01%, more preferably 0.015%, and the upper limit of Al is 0.1%, preferably 0.08%, more preferably 0. .05%.

Nb: 0.003 to 0.05%
Nb is an important element in the present invention. In order to form a predetermined grain boundary density and texture, rolling in an unrecrystallized austenite region is required. Nb is an effective element for expanding the non-recrystallization temperature range, and raises the rolling temperature and contributes to productivity improvement. In order to acquire this effect, it is necessary to make it contain 0.003% or more. However, if the Nb content exceeds 0.05%, the HAZ toughness and weldability deteriorate, so the Nb content is set to 0.05% or less.
Therefore, the lower limit value of Nb is 0.003%, preferably 0.005%, more preferably 0.008%, and the upper limit value of Nb is 0.05%, preferably 0.025%, more preferably 0. .018%.

Ti: 0.003 to 0.05%
Ti is an important element in the present invention. TiN is formed by containing Ti, and it suppresses that an austenite particle size becomes large at the time of steel bill heating. When the austenite grain size is increased, the crystal grain size of the transformed structure is also increased, so that it becomes difficult to obtain a predetermined grain boundary density, and toughness and arrestability are deteriorated. In order to obtain a necessary amount of grain boundary density so as not to lower toughness and arrestability, it is necessary to contain 0.003% or more of Ti. However, if the Ti content exceeds 0.05%, TiC is formed and the HAZ toughness decreases, so the Ti content is set to 0.05% or less.
Therefore, the lower limit value of Ti is 0.003%, preferably 0.006%, more preferably 0.008%, and the upper limit value of Ti is 0.05%, preferably 0.02%, more preferably 0. .015%.

N: 0.001 to 0.008%
N is an important element in the present invention. In order to form TiN as described above and suppress the austenite grain size from becoming large when the steel slab is heated, it is necessary to contain 0.001% or more. However, if the N content exceeds 0.008%, the steel material becomes brittle, so the N content is set to 0.008% or less.
Therefore, the lower limit value of N is 0.001%, preferably 0.0015%, more preferably 0.002%, and the upper limit value of N is 0.008%, preferably 0.0065%, more preferably 0. 0.006%.

  In the component composition of the steel sheet according to the present embodiment, the balance of the elements described above may be Fe and inevitable impurities. However, the component composition of the steel plate according to the present embodiment may contain at least one of Cu, Ni, Cr, Mo, V, B, Ca, Mg, and REM as necessary. The lower limit value of the content of these elements is 0%, but a lower limit value may be set in order to stably obtain the effect of addition. Moreover, even if these elements are contained in trace amounts at the impurity level, it is acceptable in the present invention. 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 unavoidable impurities, a steel sheet whose content is within the scope of claims is considered within the scope of the present invention.

Cu: 0 to 1%
By adding Cu, the strength and toughness of the base material can be improved.
However, if the Cu content is too large, the HAZ toughness and weldability deteriorate, so 1% is made the upper limit. The lower limit value of Cu is 0%, but the lower limit value may be 0.1% in order to stably obtain the effect of addition.
Therefore, the lower limit of Cu is 0%. In order to improve the strength and toughness of the base material, the lower limit may be 0.1% or 0.2%. In order to improve HAZ toughness and weldability, the upper limit value of Cu may be limited to 1%, 0.8%, 0.5%, or 0.3% as necessary.

Ni: 0 to 2%
By adding Ni, the strength and toughness of the base material can be improved.
However, if the Ni content is too large, the HAZ toughness and weldability deteriorate, so 2% is made the upper limit. The lower limit value of Ni is 0%, but the lower limit value may be 0.1% in order to stably obtain the effect of addition.
Therefore, the lower limit of Ni is 0%. In order to improve the strength and toughness of the base material, the lower limit may be 0.1% or 0.2%. The upper limit value of Ni may be limited to 2%, 1%, 0.5%, or 0.3% as necessary.

Cr: 0 to 1%
By adding Cr, the strength and toughness of the base material can be improved.
However, if the Cr content is too high, the HAZ toughness and weldability deteriorate, so 1% is made the upper limit. Although the lower limit of Cr is 0%, the lower limit may be set to 0.1% or 0.2% in order to stably obtain the effect of addition. The upper limit value of Cr may be limited to 1%, 0.8%, 0.5%, or 0.3% as necessary.

Mo: 0 to 0.5%
By adding Mo, the strength and toughness of the base material can be improved.
However, if the Mo content is too large, the HAZ toughness and weldability deteriorate, so 0.5% is made the upper limit. Although the lower limit of Mo is 0%, the lower limit may be set to 0.01% or 0.02% in order to stably obtain the effect of addition. The upper limit value of Mo may be limited to 0.5%, 0.3%, 0.2%, or 0.1% as necessary.

V: 0 to 0.15%
By adding V, the strength and toughness of the base material can be improved.
However, if the V content is too large, the HAZ toughness and weldability deteriorate, so 0.15% is made the upper limit. The lower limit value of V is 0%, but the lower limit value may be 0.01% or 0.02% in order to stably obtain the effect of addition. The upper limit value of V may be limited to 0.15%, 0.1%, 0.07%, or 0.05% as necessary.

B: 0 to 0.005%
By adding B, the strength and toughness of the base material can be improved.
However, if the B content is too large, the HAZ toughness and weldability deteriorate, so 0.005% is made the upper limit. The lower limit value of B is 0%, but the lower limit value may be 0.0002% or 0.0003% in order to stably obtain the effect of addition. The upper limit value of B may be limited to 0.005%, 0.003%, 0.002%, or 0.001% as necessary.

Ca: 0 to 0.01%
By adding Ca, the HAZ toughness is improved. However, if the Ca content is too large, the HAZ toughness and weldability deteriorate, so 0.01% is made the upper limit. The lower limit value of Ca is 0%, but the lower limit value may be 0.0002% or 0.0003% in order to stably obtain the addition effect. The upper limit value of Ca may be limited to 0.01%, 0.005%, 0.003%, or 0.001% as necessary.

Mg: 0 to 0.01%
By adding Mg, the HAZ toughness is improved. However, if the Mg content is too large, the HAZ toughness and weldability deteriorate, so 0.01% is made the upper limit. The lower limit value of Mg is 0%, but the lower limit value may be 0.0002% or 0.0003% in order to stably obtain the effect of addition. The upper limit value of Mg may be limited to 0.01%, 0.005%, 0.003%, or 0.001% as necessary.

REM: 0 to 0.01%
By adding REM, HAZ toughness is improved. However, if the REM content is too large, the HAZ toughness and weldability deteriorate, so 0.01% is made the upper limit. The lower limit value of REM is 0%, but the lower limit value may be 0.0003% or 0.0005% in order to stably obtain the addition effect. The upper limit of REM may be limited to 0.01%, 0.005%, 0.003%, or 0.001% as necessary.

  In order to improve the strength and toughness of the base material, the above-mentioned selective elements can be intentionally added. However, it is not necessary to add any of these selective elements in order to reduce alloy costs. Even if these elements are not intentionally added, Cu: 0.1% or less, Ni: 0.1% or less, Cr: 0.1% or less, Mo: 0.01 % Or less, V: 0.01% or less, B: 0.0002% or less, Ca: 0.0003% or less, Mg: 0.0003% or less, REM: 0.0003% or less can be contained in the steel. .

(Carbon equivalent: 0.30 to 0.50%)
The steel plate according to the present embodiment has a carbon equivalent Ceq. Is 0.30 to 0.50%.
Ceq. = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5
... (1) Formula Here, each component is the mass% of each component contained in the steel plate.

When the carbon equivalent is less than 0.30%, the strength required for the high-strength thick steel plate cannot be satisfied. When the carbon equivalent exceeds 0.50%, the arrestability required for the high-strength thick steel plate cannot be satisfied.
Therefore, the lower limit value of the carbon equivalent is 0.30%, preferably 0.32%, more preferably 0.34%, still more preferably 0.36%, and the upper limit value of the carbon equivalent is 0.50%, preferably Is 0.44%, more preferably 0.42%, still more preferably 0.40%.

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

  First, molten steel adjusted to a desired component composition is melted by a known melting method using a converter or the like, and is made into a steel slab by a known casting method such as continuous casting.

After the steel piece is cooled to a plate thickness center temperature of 600 ° C. or lower, it is charged in a heating furnace having an ambient temperature of 1000 to 1250 ° C. for 30 to 600 minutes and extracted at a plate thickness center temperature of 950 to 1150 ° C.
The charging to the heating furnace with the temperature of the cooled slab exceeding 600 ° C. has not yet completed the transformation from austenite to ferrite during cooling, so there is a refining effect due to reverse transformation to austenite during heating. This is because it is difficult to obtain and it is difficult to increase the grain boundary density after rolling with coarse austenite grains. Preferably, it is 500 degrees C or less.
If the atmospheric temperature of heating is less than 1000 ° C., sufficient heating cannot be performed and solution formation becomes insufficient. When the atmospheric temperature exceeds 1250 ° C., austenite grains become coarse, and it becomes difficult to increase the grain boundary density in the subsequent rolling process. A preferable range of the atmospheric temperature is 1050 to 1200 ° C.
This is because, when the charging time into the heating furnace is less than 30 minutes, solution formation is insufficient, and when more than 600 minutes, austenite grains become coarse. A preferred charging time range is 40 to 500 minutes.
When the plate thickness center temperature during heat extraction is less than 950 ° C., the solution is not sufficiently formed, and the hardenability is reduced due to the refinement of the austenite grains. Therefore, the plate should be thick and have high strength. Is difficult.
When the plate thickness center temperature at the time of heat extraction exceeds 1150 ° C., austenite grains become coarse, and it becomes difficult to increase the grain boundary density in the subsequent rolling process, and furthermore, the time to wait for the temperature to drop until the start of rolling. As a result, productivity is lowered. The range of preferable heat extraction temperature is 1000-1100 degreeC.

  Next, at a plate thickness center temperature of 850 to 1150 ° C., the 1-pass reduction rate is 3 to 30% for 4 to 15 passes, less than 3% is within 3 passes (including 0), and the cumulative reduction rate is 15 to 70%. Rough rolling is performed.

When the plate thickness center temperature exceeds 1150 ° C., the recrystallized austenite grains cannot be made fine even in the subsequent finish rolling. When the plate thickness center temperature is less than 850 ° C., the productivity is lowered. A preferable thickness center temperature is 900 to 1000 ° C.
If the one-pass rolling reduction is less than 3%, austenite grains grow abnormally, so it is necessary to avoid them as much as possible. However, if rolling with a 1-pass reduction rate of less than 3% is limited to 3 passes and rolling with a 1-pass reduction rate of 3 to 30% is performed for 4 passes or more, refinement by sufficient recrystallization can be achieved. However, if it exceeds 30%, the load on the rolling mill is large, and if the number of passes exceeds 15 passes, the productivity is lowered. Therefore, the 1-pass rolling reduction is 30% as the upper limit, and the number of passes is 4 to 15 passes. Rolling with a 1-pass reduction of 5 to 25% is preferably 6 to 13 passes.
The reason why the cumulative rolling reduction of rough rolling is 15 to 70% is that when the cumulative rolling reduction is less than 15%, it is difficult to refine by austenite recrystallization, the porosity remains, internal cracks and ductility, and This is because deterioration of toughness may occur, and if it exceeds 70%, the number of passes increases and productivity decreases. A preferred cumulative rolling reduction is 30 to 60%.

Next, finish rolling with a sheet thickness center temperature of 750 to 850 ° C. and 4 to 15 passes, an average value of the shape ratio (m _j ) of the following formula (2) is 0.5 to 1, and a cumulative reduction ratio is 40 to 80%. Apply.
m _j = 2 {R (H _j−1 −H _j )} ^1/2 / (H _j−1 + H _j ) (2) where j is the number of rolling passes and m _j is the jth pass. , R represents a roll radius (mm), and H _j represents a plate thickness (mm) after j passes.

When the plate thickness center temperature exceeds 850 ° C., it does not sufficiently enter the non-recrystallized region, the increase in dislocation is suppressed, and the crystal grain boundary density cannot be increased. When the sheet thickness center temperature is less than 750 ° C., the productivity is lowered and part of the processed ferrite is included, so that it is within 15 ° with respect to the plane perpendicular to the main rolling direction of the steel sheet having a half thickness. It is difficult to make the area ratio of the {110} plane forming the angle of 40% or more. A preferable thickness center temperature is 760 to 840 ° C. In rolling with less than 4 passes, it is difficult to ensure a shape ratio of 1 or less, and when it exceeds 15 passes, productivity decreases. A preferable number of passes is 5 to 13 passes.
The shape ratio of the equation (2) is an index representing what kind of strain component is imparted to the steel sheet by rolling. If the shape ratio is small, a large amount of shear strain component is applied, and if the shape ratio is large, a large amount of compressive strain component is applied. Since the change of the strain component due to the change in the shape ratio has a great influence on the formation of the texture having a quarter of the plate thickness, the range is set as described above.
The reason why the average value of the shape ratio is 0.5 to 1 is that, in 1/4 part of the plate thickness, if it is less than 0.5, the shear strain of rolling becomes dominant, and the {100} texture is developed. It is difficult to make the area ratio of the {100} plane that forms an angle of 15 ° or less with respect to a plane perpendicular to the main rolling direction of the steel sheet less than 40%. This is because the {110} texture is developed, and it is difficult to make the area ratio of the {100} plane more than 10%. The range of the average value of a preferable shape ratio is 0.6 to 0.9. If the cumulative rolling reduction is less than 40%, it is difficult to increase the grain boundary density due to the accumulation of dislocations and develop the specified texture. If the cumulative rolling reduction exceeds 80%, the effect of increasing the grain boundary density due to the accumulation of dislocations is saturated. In addition, the productivity is reduced, so 40 to 80%. A preferable range of the cumulative rolling reduction is 45 to 75%.

  Subsequent to the hot rolling, accelerated cooling is performed from a sheet thickness center temperature of 700 ° C. or more to a temperature of 550 ° C. or less at a sheet thickness center cooling rate of 2 to 10 ° C./s.

  If the plate thickness center temperature at the start of cooling is less than 700 ° C., ferrite transformation proceeds and coarsens, so it is difficult to increase the grain boundary density. When the sheet thickness center cooling rate is less than 2 ° C./s, it is difficult to increase the grain boundary density. The plate thickness center cooling rate exceeding 10 ° C./s is difficult to achieve with a steel plate having a plate thickness of 60 mm or more, so this is the upper limit. When the cooling stop temperature exceeds 550 ° C., it is difficult to increase the grain boundary density. Although there is no particular need to define the lower limit of the cooling stop temperature, it cannot be a temperature lower than the water temperature. Preferred accelerated cooling conditions are a plate thickness center temperature at the start of cooling of 720 ° C. or higher, a cooling rate of 3 to 8 ° C./s, and a cooling stop temperature of 500 ° C. or lower.

  In addition, the steel plate which concerns on this embodiment can be manufactured by controlling manufacture using the plate | board thickness center temperature of a steel plate. By using the plate thickness center temperature, it is possible to properly control the manufacturing conditions even when the plate thickness changes compared to the case where the surface temperature of the steel plate is used. Can be manufactured efficiently.

  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 the strength. A preferable tempering temperature is 400-600 degreeC.

The steel plate according to the present embodiment is applicable as a steel plate having a plate thickness of 60 to 95 mm and a yield stress of 390 to 690 MPa. In particular, the present invention is applicable to the production of steel sheets having a yield stress of 390 MPa class, 460 MPa class or higher for hulls and offshore structures.
As described above, according to the present embodiment, the arrestability can be improved so that Kca at −20 ° C., which exhibits arrestability, is 6000 N · mm ^−0.5 or more. Moreover, it can be set as the high strength thick steel plate which is low in manufacturing cost, has high productivity, has no HAZ toughness deterioration, and has excellent arrestability.

The effects of the present invention will be described below based on examples.
The component composition of the molten steel was adjusted in the steel making process, and then the steel pieces A to Z were manufactured by continuous casting. Steel slabs A to O are invention steels, and steel slabs P to Z are comparative steels.

  In Examples 1 to 20 and Comparative Examples 21 to 55, the steel slabs A to Z were reheated and further subjected to thick plate rolling to obtain a thick steel plate having a thickness of 60 to 95 mm. Subsequently, the thick steel plate was water-cooled. did. However, in Comparative Example 53, air cooling was performed instead of water cooling. Thereafter, heat treatment was performed as necessary.

Tables 1 and 2 show the composition of the steel pieces A to Z. The underline in Table 1 and Table 2 indicates that the numerical value is outside the range of the present invention, and the italic type indicates the analytical value of the amount contained as an unavoidable impurity.
Tables 3 to 6 show the production methods. For rolling, a rolling mill having a roll radius of 600 mm was used. Productivity was evaluated by the time required from the time of extraction from the heating furnace to the start of cooling after completion of rolling, and the production time of less than 1000 s was defined as good. Underlines in Tables 3 to 6 indicate that the conditions are not preferable, or that the productivity deviates from the value defined as good. The temperature and cooling rate in the manufacturing method are values at the center position of the plate thickness, and were obtained from the measured surface temperature by heat conduction analysis using a known differential method.

  Each manufactured steel plate was measured for the microstructure fraction, texture, grain boundary density, and mechanical properties.

  The microstructure phase fraction was obtained by photographing the microstructure at a magnification of 500 times at 1/2 part of the plate thickness with an optical microscope, obtaining the total area of each phase by image analysis, and dividing by the measured area.

  The grain boundary density was measured by measuring the area of ¼ part and ½ part of the thickness of 500 μm × 500 μm at 1 μm pitch by EBSD method, and the boundary where the crystal orientation difference between adjacent grains was 15 ° or more. It was defined as a crystal grain boundary, and was obtained by dividing the total length of the crystal grain boundary at that time by the measurement area.

The texture is {100} plane that forms an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction of the steel sheet at ¼ part of the plate thickness, and {110} plane at ½ part of the plate thickness. Each map was created, and the total area was divided by the measured area to obtain the area ratio.
Among the mechanical properties, the yield stress and Charpy absorbed energy of the base material were tested using test pieces taken from the center of the plate thickness, and the results were used as representative values for each steel plate.

  The tensile test was carried out in accordance with “Metal Material Tensile Test Method” of JIS Z 2241 (1998), and two of them were tested and measured, and the average value was obtained. The tensile test piece was a No. 4 test piece of JIS Z 2201 (1998).

  Charpy absorbed energy was tested in accordance with JIS Z 2242 (2005) “Charpy impact test method for metallic materials” using a 2 mm V notch Charpy impact test piece. The average value was obtained.

  For the arrestability of the base material, the arrest toughness value Kca at −20 ° C. was determined by a temperature gradient type standard ESSO test (original thickness and plate width of 500 mm).

  As for joint toughness, a butt welded joint was produced by a submerged arc welding method with a welding heat input of 10 kJ / mm, and a notch of a 2 mm V-notch Charpy impact test piece was inserted along the fusion line (FL) at 1/4 part of the plate thickness. The average value of the absorbed energy of each three at -20 ° C. The Charpy impact test was based on “Charpy impact test method for metal materials” of JIS Z 2242 (2005).

These measurement results for the thick steel plates of Examples 1 to 20 and Comparative Examples 21 to 55 are shown in Table 7. Here, it is defined that Charpy absorbed energy is 100 J or more and Kca is 6000 N · mm ^−0.5 or more.
The underline in Table 7 indicates that the conditions are out of the scope of the present invention, or that the characteristics of the steel sheet are outside the values defined as good.

  Since Examples 1-20 satisfy all the conditions of the present invention, the strength, toughness, arrestability, joint toughness, and productivity are all good.

  In Comparative Examples 21 to 55, the underlined condition was not within the scope of the present invention, so that good results were not obtained in the following points.

  Comparative Examples 21 to 31 had a problem in at least one of strength, toughness, arrestability, and joint toughness because the component range deviated from the scope of the present invention.

  In Comparative Example 32, the temperature before heating of the steel slab was too high, so the grain boundary density was low, and the toughness and arrestability were low.

  In Comparative Example 33, since the atmosphere temperature of the heating furnace was too high, the crystal grain boundary density at ¼ part of the plate thickness was small, and the arrestability was low.

  In Comparative Example 34, since the heating time was too short, the grain boundary density was small, and the toughness and arrestability were low.

  In Comparative Example 35, since the heating time was too long, the grain boundary density was low, and the toughness and arrestability were low.

  In Comparative Example 36, the heat extraction temperature was too high, so the productivity was low, the crystal grain boundary density was low, and the toughness and arrestability were low.

  In Comparative Example 37, since the heat extraction temperature was too low, the ferrite fraction was high and the strength was low.

  In Comparative Example 38, since the number of passes of less than 3% in the rough rolling was too large, the grain boundary density was low, and the toughness and arrestability were low.

  In Comparative Example 39, since the number of passes of 3 to 30% in the rough rolling was too small, the grain boundary density at ½ part of the plate thickness was small, and the arrestability was low.

  Since the comparative example 40 had too many 3-30% passes of rough rolling, productivity was remarkably low.

  In Comparative Example 41, the heat extraction temperature was high and the rough rolling temperature was too high. Accordingly, the grain boundary density was small, and the toughness, arrestability, and productivity were low.

  In Comparative Example 42, since the cumulative rolling reduction of rough rolling was too small, the grain boundary density was low, and the toughness and arrestability were low.

  In Comparative Example 43, the number of passes of 3 to 30% of the rough rolling was large, and the cumulative rolling reduction ratio of the rough rolling was excessively large. Accordingly, the productivity was remarkably low.

  In Comparative Example 44, since the finish rolling temperature was too high, the grain boundary density was low, and the toughness and arrestability were low.

  In Comparative Example 45, since the finish rolling temperature was too low, the {110} area ratio of 1/2 part of the plate thickness was small, and the arrestability and productivity were low.

  In Comparative Example 46, the number of passes for finish rolling was small, and the shape ratio was too large. Accordingly, the {100} area ratio of ¼ part of the plate thickness was small, and the arrestability was low.

  Comparative Example 47 had low productivity because the number of finish rolling passes was too large.

  In Comparative Example 48, since the average shape ratio of finish rolling was too large, the {100} area ratio of ¼ part of the plate thickness was small, and the arrestability was low.

  In Comparative Example 49, since the average shape ratio of finish rolling was too small, the {100} area ratio of 1/4 part of the plate thickness was large, and the arrestability was low.

  In Comparative Example 50, since the cumulative rolling reduction of finish rolling was too small, the {100} area ratio of ¼ part of the plate thickness, the {110} area ratio of ½ part of the plate thickness, and the grain boundary density The toughness and arrestability were low.

  In Comparative Example 51, since the cumulative rolling reduction of finish rolling was too large, the productivity was low.

  In Comparative Example 52, since the cooling start temperature was too low, the {100} area ratio of ¼ part of the plate thickness, the {110} area ratio of ½ part of the plate thickness, and the grain boundary density were small. Strength, toughness, arrestability, and productivity were low.

  Since Comparative Example 53 is cooling by air cooling, the {100} area ratio of ¼ part of the plate thickness, the {110} area ratio of ½ part of the plate thickness, and the grain boundary density are small, and the strength, toughness, And arrestability was low.

  In Comparative Example 54, since the cooling stop temperature was too high, the grain boundary density was small, and the toughness and arrestability were low.

  Comparative Example 55 had a low strength because the tempering temperature was too high.

  From the above examples, by applying the present invention, a high-strength thick steel plate having low arresting cost, high productivity, high strength, thick plate thickness, and no deterioration in HAZ toughness. It was confirmed that we can provide.

  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, it is possible to provide a high-strength thick steel plate that is low in manufacturing cost, high in productivity, high in strength, thick in plate thickness, and free from deterioration of HAZ toughness and excellent in arrestability.

Claims (8)

% By mass
C: 0.04 to 0.16%,
Si: 0.01 to 0.5%,
Mn: 0.75 to 2.5%,
Al: 0.001 to 0.1%,
Nb: 0.003 to 0.05%,
Ti: 0.003 to 0.05%,
N: 0.001 to 0.008%
Containing
P is 0.03% or less S is 0.02% or less Cu is 1% or less,
Ni is 2% or less,
Cr is 1% or less,
Mo is 0.5% or less,
V is 0.15% or less,
B is 0.005% or less, Ca is 0.01% or less,
Mg is 0.01% or less,
REM is limited to 0.01% or less,
And the balance of iron and inevitable impurities,
The carbon equivalent Ceq. Has a component composition of 0.30 to 0.50%,
Having a microstructure containing ferrite of 70% or less in area ratio and bainite of 30% or more in area ratio;
At 1/4 part of the plate thickness, the crystal grain boundary density, which is the total length per unit area of the crystal grain boundary with a crystal orientation difference of 15 ° or more, is 400 to 1000 mm / mm ² and is perpendicular to the main rolling direction. The area ratio of the {100} plane that forms an angle of 15 ° or less with respect to the plane is 10 to 40%,
In the ½ part of the plate thickness, the area of the {110} plane forming the crystal grain boundary density of 300 to 900 mm / mm ² and forming an angle of 15 ° or less with respect to the plane perpendicular to the main rolling direction. A high-strength thick steel plate characterized by having a rate of 40 to 70% and a yield stress of 390 MPa to 690 MPa .
Ceq. = C + Mn / 6 + (Cu + Ni) / 15 + (Cr + Mo + V) / 5 (1)   The high-strength thick steel plate according to claim 1, wherein the plate thickness is 60 to 95 mm. The high-strength thick steel plate according to claim 1 or 2 , wherein the microstructure contains pearlite having an area ratio of 10% or less. The microstructure, ferrite area ratio is less than 50%, pearlite area ratio is 5% or less, bainite area ratio according to any one of claims 1 to 3, characterized in that 50% or more High strength thick steel plate. The crystal grain boundary density of the plate thickness ¼ part is 500 to 900 mm / mm ² ,
The grain boundary density of the plate thickness ½ part is 400 to 800 mm / mm ²
The high-strength thick steel plate according to any one of claims 1 to 4 , wherein: Cu is 0.5% or less,
Ni is 1% or less,
Cr is 0.5% or less,
Mo is 0.2% or less,
V is 0.07% or less,
The high-strength thick steel plate according to any one of claims 1 to 5 , further limited to: The high-strength thick steel plate according to any one of claims 1 to 6 , wherein the B is further limited to 0.002% or less. 0.003% or less of the Ca,
Mg is 0.003% or less,
The high-strength thick steel plate according to any one of claims 1 to 7 , wherein the REM is further limited to 0.003% or less.