JP2006193810A - Method for producing low yield ratio high tensile strength steel sheet having excellent gas cutting crack resistance and high heat input welded joint toughness and reduced acoustic anisotropy - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide an advantageous method for producing a high tensile strength steel sheet having excellent gas cutting crack resistance and high heat input welded joint toughness, also having reduced acoustic anisotropy, further having high plastic deformability, and having tensile strength in the class of 590 MPa. <P>SOLUTION: A steel stock in which a chemical componential composition is suitably controlled is heated to the temperature range of 950 to 1,300°C, is then rolled at a cumulative draft in the temperature range of an austenite unrecrystallization temperature t(°C) shown by prescribed formula or below so as to be finished at ≥(austenite unrecrystallization temperature t-80°C) to ≤1,100°C, is thereafter directly quenched from a temperature of ≥780°C to ≤300°C at the cooling rate of ≥3°C/s, is successively subjected to reheating quenching in the temperature range of 760 to 840°C, and is subsequently tempered in the temperature range of 450 to 550°C. <P>COPYRIGHT: (C)2006,JPO&NCIPI

Description

  The present invention is excellent in gas cut cracking resistance, joint toughness in super high heat input welding up to about 100 KJ / mm, small acoustic anisotropy, large plastic deformability, and tensile strength of 590 MPa. The present invention relates to a useful method for producing the above-described high-strength steel sheet for welded structure.

In recent years, with the progress of super-high-rise buildings and fewer main girder bridges, in addition to reducing preheating during welding, application of high heat input welding, application of ultrasonic flaw detection tests to welded joints, and ultimate earthquake Application of proof stress design is being promoted on the design and construction side. Accordingly, the following items (1) to (4) have been studied as the steel material side used for them.
(1) In response to the requirement for preheating reduction, reduction of weld cold cracking sensitive composition (P _CM )
(2) For the application of high heat input welding, toughness of all parts of the welded joint that can withstand the increase in heat input,
(3) For application to the ultrasonic flaw detection test (UT) to the welded joint, the sound velocity ratio between the steel material of the transverse wave used for the oblique angle UT and the STB test piece is, for example, more than 25 mm thick, When flaw detection is performed at a nominal bending angle of 70 °, the acoustic anisotropy is small by being within the range of V / V _STB 0.995 to 1.015 (superstructure of welded steel structures in the Architectural Institute of Japan). According to the definition of sonographic standards),
(4) For the application of ultimate strength design for earthquakes, it is required that the yield ratio (yield point / tensile strength x 100%) is small (ie, the plastic deformability is high) (for construction) 80% or less), and the thickness and tensile strength of the steel materials used are also increased in thickness and strength (maximum thickness is 100mm, for architectural use, 590MPa or more, for bridges,
570 MPa or more).

For example, a technique as shown in Non-Patent Document 1 has been proposed as one of the methods for producing a 590 MPa class steel that requires a low yield ratio and small acoustic anisotropy required for building steel. In this technique, a chemical composition of 0.13C-0.25Si-1.44Mn-0.21Cu-0.20Ni-0.08Cr-0.17Mo-0.04V and a carbon equivalent Ceq of 0.45% is 80 mm. Reheating quenching (RQ) consisting of a temperature of Ac ₃ point or higher, quenching from a two-phase region (Ac ₁ point or higher and less than Ac ₃ point) (Q '), and tempering at a temperature lower than Ac ₁ point. By using the QQ'-T method for applying (T), a yield ratio of 80% or less and a tensile strength of 590 MPa or more are obtained as a two-phase structure of soft and excellent ductile ferrite and harder bainite.

  The medium carbon 590 MPa grade steel sheet obtained by this technology exhibits low yield ratio and low acoustic anisotropy, and is used for construction purposes. In order to ensure the strength of the steel, it is necessary to make the carbon equivalent Ceq higher than that of the conventional 590 MPa grade steel sheet (for example, the carbon equivalent Ceq is more than 0.42%), and the preheating temperature for preventing low temperature cracking is inevitably as high as 100 ° C. Therefore, the toughness of the high heat input welded joint is also low.

  As a method for utilizing the ultra-low carbon bainite structure to improve acoustic anisotropy and producing a steel material having a tensile strength of 570 MPa class or 590 MPa class, for example, each technique as described in Patent Documents 1 to 3 is also available. Proposed.

  Among them, in the technique of Patent Document 1, C: 0.001% or more and less than 0.030%, Si: 0.60% or less, Mn: 0.20 to 3.00%, Ni: 2.0% or less, Cu : After the steel material which becomes a composition which respectively contains 0.7-2.0% and Al: 0.10% or less is heated and cooled to the temperature of 860 degreeC or more, it is in the temperature range of 500 degreeC or more and less than 800 degreeC. A method for producing a high strength steel material of 570 MPa class with little material variation and small acoustic anisotropy by reheating and cooling is disclosed.

  In the technique of Patent Document 2, C: 0.005 to 0.025%, Si: 0.60% or less, Mn: 0.4 to 1.6%, P: 0.025% or less, S: 0.00. 010% or less, Al: 0.1% or less, Cu: 0.6-2%, Ni: 0.25-2.0%, Ti: 0.001-0.050% and B: 0.0002-0 A slab comprising a chemical composition composed of .0030% and a weight ratio of Mn / Cu: 2.0 or less and 117 Mn (%) + 163 Cu (%): 250 to 350%, with the balance being Fe and inevitable impurities , After reheating to 1050 to 1250 ° C., by performing hot rolling with a cumulative rolling reduction in a temperature range of 950 ° C. or less of 50% or less and a finish rolling temperature of 800 ° C. or more, toughness in the thickness direction of the steel sheet as it is rolled 590MPa welding electrode with excellent acoustic anisotropy It discloses a method of manufacturing the steel sheet.

  Furthermore, Patent Document 3 discloses that C: 0.025 to 0.045%, Nb: 0.005 to 0.1%, Mo-free acoustic anisotropy is small, and tensile strength 590 MPa excellent in weldability. A method for producing the above non-tempered low yield ratio high strength steel sheet is shown.

  However, the following problems are pointed out also in the high-tensile steel sheets having an ultra-low carbon bainite structure proposed so far. First, in the technique of Patent Document 1, although the acoustic anisotropy of the obtained steel material is small, the yield ratio exceeds 80% and cannot be used for architectural purposes.

In the technique of Patent Document 2, although the obtained steel sheet has small acoustic anisotropy and the yield ratio satisfies the required value (80% or less) for the steel sheet for construction, the effect of increasing the strength due to precipitation of Cu is reduced by cooling after rolling. Depending on the process, when the cooling rate is fast, the above effect is not always obtained stably, and such an effect depends on the plate thickness. The large heat input HAZ toughness is also about 52 to 71 J in the heat cycle Charpy (absorption energy vE ₀ at 0 ° C.) with a heat input of 50 KJ / mm (Table 2 in Examples), and the average of 70 J targeted in the present invention. The above cannot be achieved stably. Further, in this technique, Cu and Ni are included as essential components, but no consideration is given to their appropriate blending ratio (Ni / Cu is 0.47 to 0.95 in Table 1 of the Examples). The gas cut surface is susceptible to cracking due to Cu liquefaction in a direction parallel to the surface.

  The technique of Patent Document 3 satisfies the required characteristics as a steel sheet for construction by realizing a small acoustic anisotropy and a low yield ratio while being air-cooled, but the heat input amount of the large heat input welded joint is 20 KJ. / Mm, which is relatively small (for example, Table 3 of Examples), there is no guarantee that high HAZ toughness can be secured at the welding heat input up to about 100 KJ / mm which is the target of the present invention.

  On the other hand, technologies such as Patent Documents 4 to 6 have been proposed as steel materials having a tensile strength of 590 MPa and excellent in toughness of high heat input welded joints at a heat input of 20 KJ / mm or more.

Among these, Patent Document 4 includes C: 0.001 to 0.03%, Mn: 0.8 to 3.0%, B: 0.0003 to 0.0050%, and Ti / Al is 5. 0 or more, and Ti oxide: 20 to 90%, Al ₂ O ₃ : 70% or less, one or two of Ca oxide and REM oxide: 5 to 50%, MnO: A method for producing a tensile strength 570 MPa grade steel material in which oxide inclusions of 15% or less are dispersed is disclosed.

  In Patent Document 5, C: 0.02% or less, Mn: 0.5 to 2.0%, Nb: 0.010 to 0.10%, B: 0.0003 to 0.0040%, and Non-tempered steel for low temperature use with a tensile strength of 552 to 605 MPa having a composition of B / N of 0.3 to 1.0 is disclosed.

  Further, in Patent Document 6, C: 0.01 to 0.06%, Mn: 1.25 to 2.5%, Cr: 0.1 to 2.0%, Mo: 1.5% or less (0 %), Ti: 0.005 to 0.03%, B: 0.0006 to 0.005%, O: 0.0025 to 0.015%, ([Mn] + 1.5 × [ A high-strength steel sheet having a parameter KP defined by Cr] + 2 × [Mo]) of 2.4% by mass or more is disclosed. Moreover, as a process for producing such a high-tensile steel plate, rolling is completed at 850 to 950 ° C., then cooled, then re-heated to 750 to 800 ° C. and then water-quenched, and finally tempered at 550 to 600 ° C. (Eg, Tables 3 and 4 in the Examples).

  However, in the technique of the above-mentioned Patent Document 4, Nb is contained in an amount of about 0.005 to 0.10% (0.04 to 0.05% in the examples), and the rolling end temperature is 800 ° C. or more (Examples). 820 to 850 ° C.), and cannot stably satisfy the acoustic anisotropy required for a steel material for welded structures for architectural purposes (this point will be described later).

  In the technique of Patent Document 5, Nb: 0.010 to 0.10% is contained as in Patent Document 4, and the rolling of these steel materials is finished at 750 ° C. or higher, and air cooling or accelerated cooling is performed. As a result, the tensile strength of the 590 MPa class is satisfied, but the yield ratio is 82% or more and cannot be applied to architectural purposes.

In the technique of Patent Document 6, although low carbon steel (based on a C content of 0.03%) is also shown in which Cu, Nb and Mo are not added (No. 1 in Table 1 of Examples). The amount of Mn in this steel material exceeds the prescribed amount (1.6% or less) of the amount of Mn in SM570 of JIS G 3106 and high performance steel SA440 certified by the Minister of Construction. Moreover, although what made Mn amount 1.60% or less is also shown (No. 12 of Table 1), in this steel material, the strength reduction by the addition of Cu and Nb is compensated by addition of Mo, Such a steel material has a toughness as low as 71 J in vE _-40 (absorbed energy at -40 ° C), and is inferior in gas cutting property. That is, in such a steel material, it is difficult for the surface roughness of the gas cut surface to satisfy WES2801, and a notch is generated and becomes a starting point of fracture, so that it is not suitable as a structural member.

By the way, in the steel material containing Cu, as a technique for suppressing surface cracks caused by Cu and having excellent toughness of the weld heat affected zone, for example, a technique as shown in Patent Document 7 is also known. Yes. This technology is a Cu-containing steel for welded structures containing Si: 0.05 to 0.5%, Cr: 0.1 to 0.6%, B: 0.0005% or less, and cracks during rolling It is said that the toughness vE ₀ of the weld heat affected zone in CO ₂ welding at 12 KJ / mm is good. Moreover, when adding Ni, it has shown that it is desirable to prescribe | regulate to less than 1/3 of the amount of Cu in steel.

However, even if cracking during rolling can be prevented with this technique, when the Cu or Cu alloy in the steel remains in a molten state up to a low temperature in the process of slagging the steel component during gas cutting, Compared to the fact that the amount of heat input is so large that the steel is once melted in comparison, it is easy to easily enter the grain boundary of the gas cut surface of the molten Cu, and may not prevent gas cut cracks. In this state, the starting point of fracture is inherent, and it cannot be applied to a site where axial stress acts in the direction in which the crack opens.
“Japan Steel Pipe Technical Report” 122 (1988), pp. 5-10 Japanese Patent Laid-Open No. 9-256042 Claims etc. Japanese Patent Application Laid-Open No. 11-193445 Patent Claims, Examples Table 1, 2 etc. JP, 2002-53912, A Claims, Table 3 of an example, etc. JP, 2000-345239, A Claims etc. JP, 2001-20034, A Claims etc. JP, 2001-335883, A Claims, Tables 3 and 4 of an example, etc. JP, 2002-371337, A Claims etc.

  The present invention has been made in order to solve such problems in the prior art, and its purpose is excellent in gas cut cracking resistance and high heat input weld joint toughness, and has low acoustic anisotropy and plasticity. An object of the present invention is to provide an effective method for producing a high-tensile steel plate having a large deformability and a tensile strength of 590 MPa.

The production method of the high-strength steel sheet of the present invention that can achieve the above-mentioned object is: C: 0.015-0.045% (meaning mass%, the same shall apply hereinafter), Si: 0.4% or less (0% Mn: 0.8 to 1.6%, Cr: 0.5 to 1.3%, sol. Al: 0.08% or less (including 0%), B: 0.0004 to 0.003%, Cu: 0.5 to 0.95%, Ni: 0.7 to 5.0% (however, Ni The ratio [Ni] / [Cu] ≧ 1) of the content [Ni] and the Cu content [Cu], Ti: 0.005 to 0.03% and N satisfying the following formula (1) A steel material having a chemical composition that does not substantially contain Nb and Mo and has a CE _N value in the range of 0.27 to 0.33% represented by the following formula (2) is 950 to 1300 ° C. Then, the cumulative rolling reduction in the temperature range below the austenite non-recrystallization temperature t (° C.) represented by the following formula (3) is set to 60% or less (the austenite non-recrystallization temperature t− 80 ° C.) or more and 1100 ° C. or less after rolling, and then at a cooling rate of 3 ° C./second or more from a temperature of 780 ° C. or more to 300 ° C. It has a gist in that it is directly quenched until it reaches the following, followed by reheating quenching in the temperature range of 760 to 840 ° C. and then tempering in the temperature range of 450 to 550 ° C.
[Ti] × 14.0 / 47.9−0.001 ≦ [N] ≦ [Ti] × 14.0 / 47.9 + [B] × 14.0 / 10.8 (1)
However, [Ti], [N], and [B] indicate the contents (mass%) of Ti, Ni, and B, respectively.
CE _N = [C] + A (c) ・ [[Si] / 24 + [Mn] / 6 + [Cu] / 15 + [Ni] / 20 + ([Cr] + [V]) / 5 + 5 [B]] (2)
However, A (c) = 0.75 + 0.25 · tanh [20 ([C] -0.12)], [C], [Si], [Mn], [Cu], [Ni], [Cr], [ V] and [B] indicate the contents (mass%) of C, Si, Mn, Cu, Ni, Cr, V, and B, respectively.
t (℃) = 887 + 464 [C] + (732 × [V] -230 × √ [V]) + 890 × [Ti] +363 [sol.Al] -357 × [Si]
(3)
In the production method of the present invention, it is also useful to perform online leveler correction after the rolling is completed.

  The steel materials used in the present invention include (a) V: 0.005 to 0.10%, (b) Ca: 0.0005 to 0.01%, and (c) La: 0.002 as necessary. It is also effective to contain one or more selected from the group consisting of 0.02%, Ce: 0.0003 to 0.0050% and Mg: 0.0005 to 0.0030%. The characteristics of the high-tensile steel sheet can be further improved depending on the components contained.

  According to the manufacturing method of the present invention, high HAZ toughness of 70 J or more on average can be secured even if high heat input welding is performed up to about 100 KJ / mm without requiring a low yield ratio, and an earthquake with a large strain rate. In addition, it can prevent brittle fracture of welded joints, has no internal defects of gas cutting cracks, and has low acoustic anisotropy, making it an extremely reliable tensile material as the main welded structural member of high-rise buildings. A high strength steel plate with a strength of 590 MPa was obtained.

  In addition to the conventional plastic deformability with a yield ratio of 80% or less, the welding heat input of 50 to 100 KJ is applicable to the high-strength steel sheet with a tensile strength of 590 MPa for high-rise buildings. In order to realize an extremely large heat input of / mm and to resist external forces having a large strain rate like a large earthquake, the average value of 15 J or more shown in the welding guidelines for diaphragm welds in all parts of the welded joint A much higher average toughness of 70 J or more has been demanded.

  On the other hand, in the case of bridges, in addition to high heat input welding, the welded joint is in a tendency to switch from the conventional radiation transmission test to the ultrasonic oblique flaw detection test, so as in the case of high-rise buildings, The acoustic anisotropy with respect to the transverse wave is required to be small.

  In addition to a low yield ratio and small acoustic anisotropy in a 590 MPa class high-tensile steel sheet for construction, the present inventors have applied a heat-affected zone to a strength member subjected to electroslag welding up to about 100 KJ / mm. We have intensively studied to provide an average toughness of 70 J or more.

  As a result, after heating to 1350-1400 ° C., in addition to the effect of suppressing grain coarsening by TiN for the heat affected zone with a very low cooling rate, (i) crystals in the heat affected zone in the vicinity of the melting line It has been found that it is effective to suppress the precipitation of ferrite at the grain boundaries and the formation of ferrite side plates from the grain boundaries, and (ii) to make the inside of the crystal grains a fine bainite structure.

For (i) above, it has been found that it is effective to segregate the solid solution B at the grain boundaries even when the cooling rate is extremely low. For that purpose, in order to suppress the formation of BN, free N is fixed with Ti, Mo which raises the Ar ₃ transformation point is not added, and an appropriate amount of Cr which lowers the Ar ₃ transformation point is positively added. Was found to be effective.

In addition to the above (ii), the carbon equivalent CE _N defined by the above formula (2) is controlled to 0.27 to 0.33%, and the constituent C is made extremely low, and Nb is not added. It has been proved that it is effective to be an aggregate of fine bainite blocks in which the microstructure in the prior austenite crystal grains in the heat-affected zone is increased in inclination by the addition. It was also found that when the CE _N is in the range of 0.27 to 0.33%, preheating free at the time of small heat input welding can be realized. In addition, in order to compensate for the strength reduction of the base material due to these measures, it is necessary to actively utilize precipitation strengthening due to the ε-Cu phase.

  By the way, when Cu that exhibits precipitation strengthening is excessively added, a low melting point Cu or Cu alloy melted at the time of gas cutting selectively penetrates into the crystal grain boundary, so that there are many hair cracks extending to a depth of about 0.2 mm. Therefore, the Cu content is limited to 0.5 to 0.95%, and the ratio [Ni] / [Cu] of the Ni content [Ni] and the Cu content [Cu] is set to 1 or more to melt. Sometimes it is necessary to reduce the susceptibility to cracking by eutecticizing and increasing the solidification temperature.

  Also, as a measure for reducing the yield ratio of the base metal (80% or less), unlike the conventional two-phase structure of ferrite bainite, bainite transformation is promoted by addition of Mn and Cr, and the C content is reduced to 0. .015 to 0.045%, and by performing heat treatment such as quenching in a two-phase region of 760 to 840 ° C. and tempering at 450 to 550 ° C., a fine island martensite phase is formed between bainite phases. It is necessary to disperse 0.5 to 3.5%.

  Furthermore, with respect to the acoustic anisotropy of the base material, in order to satisfy the range in which it is determined that there is no sound velocity ratio with STB in Appendix Table 1 of “Ultrasonic Inspection Standards for Steel Structure Building Welds” of the Architectural Institute of Japan. The rolling reduction is finished at a temperature range of (austenite non-recrystallization temperature−80 ° C.) or more and 1100 ° C. or less as a cumulative reduction ratio of 60% or less at the austenite non-recrystallization temperature or less. This can be realized by controlling the average aspect ratio of the austenite grain size (average grain size in the rolling direction / average grain size in the sheet thickness direction) in the range of 1 to 1.2.

By applying the above measures in a comprehensive manner, even if a large heat input welding of 50 to 100 KJ / mm is applied, it has an average _value of 70 J or more at vE ₀ and is not susceptible to gas cutting cracking. The present inventors have found that a steel plate suitable for steel (SA440) can be obtained, and have completed the present invention.

  The reason for limiting the chemical composition and the microstructure for obtaining the characteristics of the present invention will be described below along the background.

The inventors of the present invention have five steel types (0.035C-1.45Mn-0.95Cu-1Ni-0.0) adjusted with the Mn content variable so that the CE _N value is constant at about 0.30%. 7Cr-0.015Ti-0.0012B-0.0040N-based basic steel, 0.4Mo-added steel, 0.01Nb-added steel, 0.4Mo-0.01Nb-added steel, and 0.05C-based steel) skin plate material And a SN490B-TMC (0.14C-0.3Si-1.25Mn-0.008Nb-0.012Ti system, 60 mm thick) diaphragm material, with a gap of 25 mm and an electroslag with a heat input of 100 KJ / mm Welding and comparing the toughness of the heat affected zone (HAZ) on the skin plate side of the welded joint (notch position: bond showing the lowest toughness in high heat input welding HAZ + 0.5 mm) And 査 (see Experimental No.2~7 the following Table 7).

From this result, even if the chemical component composition shows the same level as the CE _N value of about 0.30%, the large heat input HAZ toughness is greatly increased by adding no Mo, no Nb, adding Cr, and reducing C. It was found that, by combining these requirements and by combining these requirements, it is possible to guarantee vE ₀ 70 J or more for the first time even in the vicinity of the bond +0.5 mm position where the HAZ toughness is lowest. It was also found that the strength reduction due to the absence of addition of Mo and Nb can be compensated by precipitation strengthening of Cu and solid solution strengthening by Mn and Cr.

  In addition, steel grade with high heat input welding HAZ toughness (steel grade B in Table 1 below) is used as a basic component, and the content of Cu and Ni is changed to LP gas (LPG) cutting from the gas cut surface to the steel plate surface. A comparative investigation was made on the maximum depth of cracks extending in the parallel direction (see Experiment Nos. 2 and 8 to 17 in Table 7 below).

  From this result, it has been found that by controlling the ratio [Ni] / [Cu] to 1 or more when the Cu content is 0.95% or less, the sensitivity to gas cutting cracks is reduced.

  Furthermore, the inventors of the present invention have a basic component of 0.035C-1.45Mn-0.95Cu-1Ni-0.7Cr-0.015Ti-0.0012B-0.0040N system (steel type B in Table 1 below). Using steel grades with varying C content, the continuous cast slab was heated to 1050-1100 ° C, then hot-rolled to 900mm at 900 ° C and then directly quenched, followed by quenching at 840 ° C (Q 'treatment) A tempering treatment (T treatment) at 500 ° C. was carried out to investigate the influence of the C content on the strength, yield ratio, microstructure and HAZ toughness at 100 KJ / mm (Experiment No. 2 in Table 7 below). 6, 18-21).

From this result, it was found that when C: 0.015% or more, the yield ratio is 80% or less and the strength is compatible. Further, from the investigation results of the high heat input HAZ toughness in these steel types, it was found that the upper limit of the C content needs to be 0.045% in order to ensure an average of 70 J or more in vE ₀ .

  Next, the reason for limiting the chemical component composition in the steel material that is the subject of the present invention will be described. First, in the present invention, as described above, C: 0.015 to 0.045%, Si: 0.01 to 0.4%, Mn: 0.8 to 1.6%, Cr: 0.5 to 1.%. 3%, sol. Al: 0.08% or less (including 0%), B: 0.0004 to 0.003%, Cu: 0.5 to 0.95%, Ni: 0.7 to 5.0% (however, Ni The ratio [Ni] / [Cu] ≧ 1) of the content [Ni] and the Cu content [Cu], Ti: 0.005 to 0.03% and N satisfying the following formula (1) Although it is necessary to substantially not contain Nb and Mo, the reasons for limiting the ranges of these elements are as follows.

C: 0.015-0.045%
C is an element effective for forming low-temperature bainite and ensuring the strength and low yield ratio of high-tensile steel, and for that purpose, it is necessary to contain 0.015% or more. However, if C is contained excessively, a martensite phase is formed on the high cooling rate side and the cold cracking resistance of welding is deteriorated, and the island-like martensite phase is increased by high heat input HAZ to deteriorate toughness. I will let you. For these reasons, the upper limit needs to be 0.045%. In addition, from a viewpoint of coexistence of base material strength and high heat input welding HAZ toughness, a preferable lower limit is 0.02%, and a preferable upper limit is 0.04%.

Si: 0.4% or less (including 0%)
Si is an effective element as a deoxidizer and strengthening element. However, if excessively contained, it increases the island-like martensite phase in the high heat input welding HAZ and deteriorates toughness. For these reasons, the upper limit is set to 0.4%, and the lower content is preferably as low as possible, so the lower limit is set to 0%. When Si is not included, deoxidation can be arbitrarily replaced with Mn, Al, Ti, or the like. In addition, the preferable upper limit of Si content is 0.3%.

Mn: 0.8 to 1.6%
Mn is an element that is effective in causing the ferrite transformation to shift to a low temperature for a long time and to form and strengthen the upper bainite phase at a low temperature. For that purpose, it is necessary to contain 0.8% or more of Mn. However, when Mn is contained excessively, the toughness of the base metal and the high heat input weld HAZ and the deterioration of the weld cold crack resistance are caused, so the upper limit is made 1.6%. The minimum with preferable Mn content is 1.0%, and a preferable upper limit is 1.5%.

Cr: 0.5 to 1.3%
Cr is an element that is effective for forming the upper bainite phase at a low temperature and reducing the yield ratio by comparing the same strength level by improving hardenability. In addition, since the ferrite transformation of the high heat input weld HAZ is shifted to a low temperature and long time side, it is easy to form a bainite phase at a low temperature and effective in improving toughness. In order to exhibit such an effect, Cr needs to be contained 0.5% or more. However, if Cr is excessively contained, the island-like martensite phase is increased in the high heat input welding HAZ, so the upper limit needs to be 1.3%. In addition, the minimum with preferable Cr content is 0.6%, and a preferable upper limit is 1.1%.

sol. Al: 0.08% (including 0%)
sol. Al (soluble Al) is an element effective for deoxidation, but in order to improve toughness by forming intragranular bainite with Ti oxide as the core in high heat input welding HAZ, the content should be as small as possible The lower limit is 0%. In this case, deoxidation can be arbitrarily replaced with Mn, Si, Ti, or the like. Also, sol. Al effectively works to secure the hardenability of the base material by supplementing N fixation with Ti, but if it is excessively contained, it increases nonmetallic inclusions and causes toughness deterioration. Therefore, sol. When Al is contained, the upper limit needs to be 0.08%. In addition, sol. A preferable upper limit of Al is about 0.06%.

B: 0.0004 to 0.003%
B improves the hardenability, suppresses ferrite transformation even at a low cooling rate, and promotes the formation of bainite in a low temperature region, and is therefore effective in improving the base material strength. In the high heat input welding HAZ, the solid solution B segregates at the prior austenite grain boundaries, and the remaining BN precipitates at the grain boundaries and acts as intragranular bainite transformation nuclei. Formation is suppressed and effectively acts to improve toughness. For that purpose, B needs to be contained in an amount of 0.0004% or more. However, if B is contained excessively, the hardenability of the steel becomes too high, the island martensite phase is increased, the toughness of the base metal and the high heat input weld HAZ is deteriorated, and the resistance to welding cold cracking is also improved. Deteriorate. For these reasons, the upper limit of the B content needs to be 0.003%. In addition, the minimum with preferable B content is 0.006%, and a preferable upper limit is 0.002%.

Cu: 0.5 to 0.95%
Cu is an element effective for improving the strength of the base material by solid solution strengthening and precipitation of clusters of ε-Cu phase. In order to exhibit these effects, it is necessary to contain Cu 0.5% or more. However, if Cu is excessively contained, a Cu-concentrated phase is formed on the gas cut surface and enters the prior austenite grain boundary during thermal expansion to induce cracking, so the upper limit is made 0.95%. The minimum with preferable Cu content is 0.7%, and a preferable upper limit is 0.9%.

Ni: 0.7-5.0%
Ni is an element that improves the hardenability and improves the toughness of the base material and the base of the high heat input welding HAZ, and in order to exert these effects, it is necessary to contain 0.7% or more. However, if Ni is contained excessively, the hardenability becomes too high and the island-like martensite phase increases, leading to deterioration of toughness and also being uneconomical, so the upper limit is made 5%. In addition, the minimum with preferable Ni content is 0.9%, and a preferable upper limit is 3%.

However, [Ni] / [Cu] ≧ 1
In the high-tensile steel of the present invention, the ratio [Ni] / [Cu] of the Ni content [Ni] and the Cu content [Cu] needs to be 1 or more. By satisfying these requirements, the melting point of the Cu—Ni alloy phase concentrated on the gas cut surface can be increased, and hot cracking can be prevented.

Ti: 0.005 to 0.03%
Ti is an element effective for fixing solid solution N as TiN, increasing the amount of solid solution B, and improving the hardenability of the base material. Further, when Ti oxide is generated by Ti deoxidation, it acts as a nucleus for formation of an intragranular bainite phase in high heat input welding HAZ and improves toughness. In order to exert such effects, the Ti content needs to be 0.005% or more. However, when Ti is excessively contained, the toughness of the base metal and the HAZ is deteriorated by precipitation of TiC, so the upper limit is made 0.03%. In addition, the minimum with preferable Ti content is 0.008%, and a preferable upper limit is 0.02%.

N: In order to ensure high toughness in the high heat input welding HAZ that satisfies the above formula (1) , fine precipitation of TiN in the prior austenite grains and composite precipitation of BN, It is effective to form the nuclei of intragranular bainite. From such a viewpoint, the lower limit of the N content was (Ti stoichiometric equivalent) −0.001%, and the upper limit was the total amount of Ti and B stoichiometric equivalents. Exceeding this will cause toughness degradation of the HAZ due to the solute N and a decrease in the hardenability of the base metal.

Nb: Nb not substantially contained improves the hardenability by solid solution, but forms a plate-like coarse upper bainite phase in the prior austenite grains in the high heat input welding HAZ, and the crystal orientation is aligned. Therefore, it does not become a barrier for the fracture path and greatly deteriorates the toughness. From such a viewpoint, it is necessary that the high-tensile steel of the present invention is not substantially contained. Note that “substantially free” means that it is allowed to be mixed as an impurity (for example, 0.005% or less).

Mo: Mo which is substantially not contained is an element effective for improving the hardenability and improving the strength. However, the Ar ₃ transformation point is raised to increase the high-temperature bainite phase or the island martensite phase in the high heat input welding HAZ. Promotes formation and degrades toughness. From such a viewpoint, it is necessary that the high-tensile steel of the present invention is not substantially contained. Note that “substantially does not contain” means that it is allowed to be mixed as an impurity (for example, 0.05% or less).

In the steel material that is the subject of the present invention, it is necessary to control the CE _N value defined by the above equation (2) within an appropriate range. The reason for limiting the range is as follows.

CE _N : 0.27 to 0.33%
CE _N defined by the above formula (2) is a carbon equivalent expressing the curability of the welded HAZ. If the CE _N value is less than 0.27%, the thick material cannot satisfy the tensile strength of 590 MPa class. Also the value of CE _N exceeds 0.33%, the resistance to weld cold cracking resistance deteriorates, preheating not only necessary, toughness high heat-input welding HAZ island martensite phase is increased It becomes low, and it becomes difficult to stably secure the target average vE ₀ : 70 J at a heat input of 100 KJ / mm. Therefore, in the high-tensile steel of the present invention, the CE _N value defined by the above equation (2) needs to be in the range of 0.27 to 0.33%. The preferable lower limit of the CE _N value is 0.28%, and the preferable upper limit is 0.32%.

In addition, when the steel material contains Nb and Mo, the carbon equivalent CE _N can be expressed as the following formula (4) including these items. However, in the steel material targeted by the present invention, Since these elements are not substantially contained, the value defined by the above formula (2) is adopted as the value of the carbon equivalent CE _N. Therefore, when evaluating the carbon equivalent of the steel material containing Nb or Mo, a value obtained by the following equation (4) is adopted (the steel types A, C, D, E, etc. described later).
CE _N = [C] + A (c) ・ [[Si] / 24 + [Mn] / 6 + [Cu] / 15 + [Ni] / 20 + ([Cr] + [Mo] + [Nb] + [V]) / 5 + 5 [B]]
(4)
However, A (c) = 0.75 + 0.25 · tanh [20 ([C] -0.12)], [C], [Si], [Mn], [Cu], [Ni], [Cr], [ Mo], [Nb], [V] and [B] indicate the contents (mass%) of C, Si, Mn, Cu, Ni, Cr, Mo, Nb, V and B, respectively.

  The steel materials to be used in the present invention include (a) V: 0.005 to 0.10%, (b) Ca: 0.0005 to 0.01%, and (c) La: 0.002 as necessary. It is also effective to contain one or more selected from the group consisting of ˜0.02%, Ce: 0.0003 to 0.0050% and Mg: 0.0005 to 0.0030%. The reasons for limiting the range when these components are contained are as follows.

V: 0.005-0.10%
V is an element effective for improving the base material strength. In order to exert such an effect, it is preferable to contain V in an amount of 0.005% or more. However, if it is contained in excess of 0.10%, the high heat input welding HAZ toughness is lowered. In addition, the more preferable minimum of V content is 0.03, and a more preferable upper limit is 0.06%.

Ca: 0.0005 to 0.01%
When Ca fixes S as CaS and controls the form as granular non-metallic inclusions, the Z-direction tensile stress generated during column angle joint welding acts on the S segregated portion in the center of the plate thickness Is effective in improving drawing and toughness and preventing breakage from the segregated portion. Moreover, it combines with [O] to form CaO, and the core of bainite transformation after high heat input welding is dispersed in the prior austenite grains, and the bainite block size is refined to improve the high heat input HAZ toughness. Demonstrate. In order to exert these effects, Ca is preferably contained in an amount of 0.0005% or more. However, if the Ca content exceeds 0.01%, the effect is not only saturated, but the toughness of the base material is increased. On the contrary, it deteriorates. In addition, the more preferable minimum of Ca content is 0.01%, and a more preferable upper limit is 0.05%.

One or more selected from the group consisting of La: 0.002-0.02%, Ce: 0.0003-0.0050%, Mg: 0.0005-0.0030% La and Ce are rare earth elements It is a kind of (REM) and works effectively to fix S as a sulfide and improve the drawing and toughness of the segregation part. In addition, Ce and Mg precipitate CeO ₂ or MgO low melting point oxides in the prior austenite grains during cooling after high heat input welding, and transform the bainite block into fine bainite blocks. Thus, the high heat input welding HAZ toughness is improved by complicating the fracture path. When La, Ce, and Mg are less than the above lower limits, these effects are not exhibited. When the La, Ce, and Mg are more than the upper limits, the presence of excessive nonmetallic inclusions causes deterioration of the base metal toughness. More preferable lower limits are La: 0.002%, Ce: 0.005%, Mg: 0.001%, respectively, and more preferable upper limits are La: 0.01%, Ce: 0.002%, Mg: 0. 0020%.

  In the steel material which is the subject of the present invention, in addition to the above components, it consists of Fe and inevitable impurities, but it can also contain trace components (allowable components) to the extent that they do not impede its properties. It is also included in the scope of the present invention.

  In the production method of the present invention, a steel material having the chemical composition as described above is used, and by making the production conditions appropriate, a high-tensile steel sheet exhibiting desired properties can be obtained. The manufacturing method is basically a process of direct quenching after rolling (DQ) → two-phase region quenching (Q ′) → tempering (T) using slag produced by a continuous casting method or an ingot casting method. It is characterized by doing. Next, each manufacturing condition prescribed | regulated by this invention is demonstrated. In the following manufacturing conditions, the position of temperature control is the surface portion of the steel material, and soaking is performed in heating and heat treatment.

Slab heating temperature: 950-1300 ° C
When the slab heating temperature exceeds 1300 ° C., coarsening of prior austenite grains is caused and toughness is deteriorated. In contrast, when the slab heating temperature is less than 950 ° C., the BN precipitates contained in the slab cannot be completely dissociated, and the amount of segregation of B austenite grain boundaries during direct quenching after rolling is reduced. The hardenability is significantly reduced. Therefore, the slab heating temperature needs to be in the temperature range of 950 to 1300 ° C.

Rolling temperature range: austenite non-recrystallization temperature t (° C.) or less represented by the above formula (3), and rolling end temperature: (austenite non-recrystallization temperature−80 ° C.) or more and 1100 ° C. or less. When the temperature falls below the austenite non-recrystallization temperature t (° C.), a processed structure parallel to the (100) plane extending in the rolling direction is formed and remains even after cooling and tempering. While lowering, acoustic anisotropy appears in the rolling direction and the direction perpendicular to the rolling direction. Therefore, the rolling reduction (cumulative rolling ratio) at the austenite non-recrystallization temperature t or less needs to be as small as possible. For this purpose, the cumulative rolling reduction needs to be 60% or less. Incidentally, the lower limit of the cumulative rolling reduction at this time includes 0% (that is, rolling is not performed in the above temperature range) from the above-mentioned purpose.

  On the other hand, the rolling end temperature (rolling finishing temperature) needs to be not less than (austenite non-recrystallization temperature−80 ° C.) from the viewpoint of suppressing the generation of a rolling texture that develops acoustic anisotropy. From the viewpoint of securing the base material toughness material, it is necessary to set the temperature to 1100 ° C. or lower. The preferable lower limit of the rolling end temperature is 850 ° C., and the preferable upper limit is about 1050 ° C. By satisfying these conditions, the average aspect ratio that affects the volume anisotropy is in an appropriate range (for example, 1 To 1.2).

  That is, in order to satisfy the range in which the acoustic anisotropy is determined to have no difference in sound velocity between the STB and the test piece in Appendix Table 1 in the “Ultrasonic Inspection Standard for Steel Structure Building Welds of the Architectural Institute of Japan”. As will be described later, it is preferable to control the average aspect ratio (average particle size ratio in the main rolling direction / sheet thickness direction) within an appropriate range. In order to satisfy these conditions, It is necessary to perform rolling while appropriately setting the rolling temperature range and the rolling end temperature.

The austenite non-recrystallization temperature t (° C.) can be expressed by the following formula (5) including the item when the steel material contains Nb. Since the material does not contain this element, the austenite non-recrystallization temperature t (° C.) is the one defined by the above equation (3). Therefore, when evaluating the austenite non-recrystallization temperature t (° C.) of the steel material containing Nb, the value obtained by the following equation (5) is adopted (the steel types A, D, E, etc. described later).
t (℃) = 887 + 464 [C] + (6445 [Nb] -644√ [Nb]) + (732 × [V] -230 × √ [V]) + 890 ×
[Ti] +363 [sol.Al] -357 × [Si] (5)
However, [C], [Nb], [V], [Ti], [sol. Al] and [Si] are C, Nb, V, Ti, sol. The content (mass%) of Al and Si is shown.

Cooling conditions after completion of rolling: For cooling conditions after completion of rolling in which direct quenching is performed from a temperature of 780 ° C. or higher to a temperature of 300 ° C. or lower at a cooling rate of 3 ° C./second or more , the cooling start temperature is less than 780 ° C. Alternatively, when the cooling rate is less than 3 ° C./second, a microstructure including a pro-eutectoid ferrite phase and a coarse bainite phase is formed after cooling by passing through a ferrite nose in the continuous cooling transformation diagram. The strength and toughness of the material will decrease. Further, when the cooling stop temperature is higher than 300 ° C., the bainite phase cannot be refined, so that the strength and toughness are low. For these reasons, it is necessary to perform direct quenching (DQ) from 780 ° C. or higher to 300 ° C. or lower at a cooling rate of 3 ° C./second or higher as rolling conditions after completion of rolling. The “direct quenching” means a process of quenching the hot pieces after rolling on-line.

Reheating and quenching after cooling: Reheating and quenching in which the reheating degree quenching is performed in the temperature range of 760 to 840 ° C. As a result of excessive austenitization in the reheating temperature range exceeding 840 ° C., the quenching structure is also increased. The fraction of bainite phase becomes excessive, the number of island martensite phases increases, and the yield ratio exceeds the target 80%. On the other hand, if the reheating temperature is less than 760 ° C., the austenite formation by reheating is not sufficient, so that the bainite phase cannot be refined and the island-like martensite phase is too small, and the strength is low. Become. Moreover, even if the reheating temperature is in the temperature range of 760 to 840 ° C., only a low strength can be obtained by normalization, so that it is necessary to quench.

Tempering temperature range after reheating and quenching: 450-550 ° C
With regard to the tempering temperature after reheating and quenching, when the temperature exceeds 550 ° C., the strength decreases due to a decrease in the dislocation density, and when the temperature becomes less than 450 ° C., Cu precipitation becomes insufficient and the yield strength becomes low. Therefore, the tempering temperature after reheating and quenching needs to be in the range of 450 to 550 ° C.

  In the present invention, as a measure for reducing the yield ratio of the base material (80% or less), unlike the conventional two-phase structure of ferrite bainite, (a) the addition of Mn and Cr, and the austenite region after rolling By controlling the ferrite transformation and lowering the start of the bainite transformation by a combination of direct quenching to 300 ° C. at a cooling rate of 3 ° C./second or more and quenching after two-phase region reheating, the microstructure In addition to the fine bainite base, (b) a combination of 0.015 to 0.045% C and a two-phase region heat treatment at 760 to 840 ° C. The island-like martensite phase is finely dispersed.

  In addition, regarding the acoustic anisotropy of the base material, a range in which it is determined that there is no difference in sound velocity between the STB and the test piece in Appendix Table 1 in the “Ultrasonic Inspection Standard for Steel Structure Building Welds of the Architectural Institute of Japan”. In order to satisfy the requirements, it is necessary to appropriately regulate the cumulative rolling reduction in the rolling temperature region as described above, but in addition, by adding no Nb, the austenite recrystallization temperature is lowered. It can be made.

  Furthermore, even with the chemical component composition as described above, by performing tempering at 450 to 550 ° C. after DQ-two-phase quenching, the dislocation of the movable dislocation formed during quenching is promoted and a predetermined amount. By precipitation of the ε-Cu phase, the yield strength can be ensured to be equal to or greater than the lower limit of the standard value of SA440 (440 MPa).

  In the present invention, by satisfying the above conditions (chemical component composition and production conditions), a high-tensile steel sheet exhibiting desired properties can be obtained. It is preferable that the average aspect ratio (average grain size ratio in the main rolling direction / sheet thickness direction) and the dispersion ratio of the island-like martensite phase and ε-Cu phase clusters are controlled within an appropriate range. These preferable ranges and the reasons thereof are as follows.

Average aspect ratio of prior austenite grain size in the cross section in the thickness direction (main rolling direction / thickness direction)
Average particle size ratio): 1.0 to 1.2
When the aspect ratio of the prior austenite grain size in the cross section in the thickness direction (average grain size in the main rolling direction / average grain size in the thickness direction) exceeds 1.2, the so-called texture in which the crystal orientation is oriented in a specific direction Therefore, a range in which the acoustic anisotropy is determined to have no sound velocity ratio with STB in Appendix Table 1 in “Ultrasonic Inspection Standard of Steel Structure Building Welding by the Architectural Institute of Japan” (for example, When flaw detection is performed with a probe having a nominal refraction angle of 70 ° exceeding a plate thickness of 25 mm, it will exceed 0.995 ≦ V / V _STB ≦ 1.015), and the defect position cannot be correctly displayed in the ultrasonic flaw detection test. It becomes a construction problem. Therefore, it is preferable to control the average aspect ratio of the prior austenite (average particle size in the main rolling direction / average particle size in the plate thickness direction) in the range of 1.0 to 1.2. More preferably, the average aspect is in the range of 1.0 to 1.1.

Clusters of 0.5 to 3.5% by volume of island martensite phase and 4 × 10 ^{20 to} 26 × 10 ²⁰ pieces / m ³ of ε-Cu phase are dispersed in the bainite base. in tensile steel, in order to include a yield ratio (yield strength / tensile strength × 100%) is below the yield ratio of 80% required for the main members of the building in terms of seismic design of low-CE _N is the base of the bainite phase It is preferable to finely disperse the harder island martensite phase. When the dispersion ratio of the island-like martensite phase is less than 0.5% by volume, the yield ratio exceeds 80%. On the other hand, when the dispersion ratio of the island-like martensite phase exceeds 3.5% by volume, the toughness of the base material is low although the yield ratio is 80% or less. Therefore, the island-like martensite phase is preferably 0.5 to 3.5% by volume. More preferably, the content is 1 to 3% by volume. Further, in order to secure a tensile strength of 590 MPa or more with low CE _N , it is necessary to add precipitation strengthening of the ε-Cu phase by aging using the bainite phase as a base. When ε-Cu phase clusters are dispersed at less than 4 × 10 ²⁰ / m ³ , tensile strength is insufficient even in bainite. On the other hand, if the number of ε-Cu phase clusters exceeds 26 × 10 ²⁰ / m ³ , the toughness tends to deteriorate, and a large amount of low melting point Cu alloy is formed during gas cutting. Causing intergranular cracking on the gas cut surface. Therefore, it is preferable to disperse 4 × 10 ^{20 to} 26 × 10 ²⁰ clusters / m ³ of ε-Cu phase clusters in bainite. More preferably 7 × 10 ²⁰ ~22 × 10 ²⁰ atoms / m ^3.

  In carrying out the method of the present invention, it is also preferable to perform online leveler correction before direct quenching after the end of rolling. By adding such correction, it is possible to realize a high-tensile steel sheet that satisfies the target material characteristics and does not require cutting off due to poor flatness.

  Hereinafter, the present invention will be described in more detail with reference to examples. However, the following examples are not intended to limit the present invention, and any design modifications may be made in accordance with the gist of the present invention. It is included in the range.

Steel sheets having chemical composition compositions shown in Tables 1 to 3 below were used, and steel sheets were manufactured under the manufacturing conditions shown in Tables 4 to 7 below. Table 1 below shows the range of the formula (1) defined by the present invention, the value of CE _N defined by the formula (2), and the austenite non-recrystallization temperature t (γ unrecrystallized) defined by the formula (3). Temperature) and [Ni] / [Cu] values. In Tables 4 to 7, Q is the reheating quenching temperature consisting of a temperature of Ac ₃ point or higher, Q ′ is the quenching temperature from the two-phase region (Ac ₁ point or higher and less than Ac ₃ point), and T is Ac ₁ point. The tempering temperatures at temperatures below are shown respectively. In addition, the value of CE _N and the austenite non-recrystallization temperature t are those obtained according to the above formula (4) or (5) for those containing Nb and Mo (steel types A, C, D, E). Indicated.

  About each obtained steel plate, the aspect ratio of the prior austenite ((gamma)) particle size, the volume ratio of an island-like martensite phase, the number of (epsilon) -Cu phase clusters, etc. were measured with the following method.

[Aspect ratio of former austenite (γ) particle size]
The prior austenite grain boundaries in the plate thickness section in the main rolling direction are revealed using a corrosive solution consisting of (5 ml hydrochloric acid + 1 g picric acid + 100 ml ethanol), and the average grain size in the main rolling direction and the average grain size in the plate thickness direction are determined. Measurements were taken to determine their ratio as the average aspect ratio.

[Volume ratio of island martensite phase]
The island-like martensite phase was photographed by exposing the main rolling direction and the plate thickness section perpendicular to the main rolling direction using a repeller reagent, and the fraction (volume ratio) was calculated by an image analyzer.

[Number of clusters of ε-Cu phase]
For ε-Cu phase clusters, a thin film is taken from the cross section of the plate thickness, the ε-Cu phase is fixed and the distribution of this cluster is photographed using an analytical electron microscope, and the number per unit area is measured by an image analyzer. Was calculated.

  In addition, for each steel plate obtained, gas cutting crack sensitivity, acoustic anisotropy, tensile properties of the base material [yield strength (0.2% YS), tensile strength (TS), yield ratio (YR), base material toughness, The welding cold crack resistance and high heat input weld HAZ toughness were evaluated by the following methods, respectively.

[Gas cutting crack sensitivity]
Oxygen pressure: 0.6 MPa (6 kgf / cm ² ), LP gas pressure under appropriate cutting conditions (for example, # 5 crater in the case of 100 mm thickness) according to the plate thickness of LP gas widely used for steel plate cutting in the steel frame manufacturing process : 0.06 MPa (0.6 kgf / cm ² ), cutting speed: 210 mm / min. After cutting, the cross section perpendicular to the cut surface was observed with an optical microscope, and the maximum crack depth from the cut surface was measured. . No crack was accepted.

[Acoustic anisotropy]
STB sound velocity ratio (V / V _STB ) defined in the ultrasonic inspection standard for steel structure building welds of the Architectural Institute of Japan was measured in the main rolling direction (L direction) and the direction perpendicular to the main rolling direction (C direction). The presence or absence of a difference in sound velocity from STB was determined according to Appendix Table 1. The range of V / V _{STB in} Appendix Table 1 was considered acceptable. For example, when flaw detection is performed with a flaw having a plate thickness of more than 25 mm and a nominal refraction angle of 70 °, 0.995 ≦ V / V _STB ≦ 1.015 is determined to have no acoustic anisotropy.

[Tensile properties of base material]
A JIS Z 22014 test piece was taken in the C direction from t / 4 (t is the plate thickness) of the steel sheet and subjected to a tensile test in accordance with JIS Z 2241. Yield strength (0.2% proof stress: σ _0.2 ), tensile Strength (TS) and yield ratio (yield strength / tensile strength × 100%: YR) were measured. Yield strength: 440-540 MPa, tensile strength: 590-740 MPa, and yield ratio: 80% or less were accepted.

[Base material toughness]
JIS Z 2204 No. 4 test specimens were collected in the L direction from t / 4 of the steel sheet, subjected to an impact test in accordance with JIS Z 2242, and the fracture surface transition temperature (vTrs) was measured. The target was vTrs of −20 ° C. or less.

[Weld cold crack resistance]
According to the JIS Z 3158 y-type weld crack test method, covered arc welding was performed at a heat input of 1.7 KJ / mm, and the root crack prevention preheating temperature was measured. 25 degrees C or less was set as the pass.

[Large heat input welding HAZ toughness]
A column / diaphragm welded joint is manufactured by electroslag welding with a heat input of 100 KJ / mm. From the column (skin plate) side, the absorbed energy is often the lowest, Z- A Charpy impact test piece (JIS Z 2204 No. 4) with a notch in the T direction was taken at n = 3, and an average impact absorption energy vE ₀ at 0 ° C. was determined. An average of 70 J or more was accepted. Experiment No. For only 74 and 98 electroslag welded joints, 40 mm thick SN490B (0.16% C-0.34% Si-1.34% Mn-0.034% V steel) was used as the diaphragm material.

These results are shown in Tables 8 to 11 below, and can be considered as follows from these results. First, Experiment No. 1 is a conventional C-based SA440 steel in, for CE _N is high, weld cold cracking prevention preheating temperature as high as 100 ° C., HAZ toughness of electroslag welding in heat input 100 kJ / mm is also low.

  Experiment No. 2 is a basic component system of the present invention, the toughness of the base material satisfies the target, the weldability is good as it does not require preheating, and the HAZ toughness of electroslag welding at a heat input of 100 KJ / mm is also a target average of 70 J or more. Is sufficiently satisfied.

Experiment No. 3 to 6 are those of Experiment No. 2 and CE _N (steel types C, D, E, and F in Table 1), which are 0.4Mo series, 0.01Nb series, 0.4Mo-0.01Nb series, and 0.05C series, respectively. The heat input HAZ toughness is low.

  Experiment No. 7 to 12 were obtained by changing the Ni / Cu ratio by changing the amount of Cu while keeping Ni constant at 1% (steel types G to L in Table 1). In addition, Experiment No. 13 to 17 are Cu: 0.95%, and the amount of Ni is changed (steel types M to Q in Table 1). It can be seen that Cu: exceeding 0.95% (Experiment No. 8, 9) has gas cut cracking susceptibility even when the Ni / Cu ratio is 1 or less. In the case where Cu is less than 0.5% (Experiment No. 12), the yield strength and the tensile strength are lower than the target values.

  Experiment No. Nos. 18 to 21 are constant in the composition of the present invention except for the C amount, and only the C amount is changed (steel types R to U in Table 1). It can be seen that when the C content is less than 0.015% (Experiment No. 18), the yield strength (0.2% YS) and the tensile strength (TS) are both below the target values.

  Experiment No. No. 22 has no Q ′ treatment (two-phase quenching treatment), and the volume ratio of the island-like martensite phase constituting the microstructure of the base material falls below the range specified in the present invention, and the yield ratio (YR ) Exceeds 80% of the target value.

  Experiment No. Nos. 23, 24, and 80 are obtained by changing the cumulative reduction ratio at the γ non-recrystallization temperature t or lower and the rolling end temperature (rolling finishing temperature), and the cumulative reduction ratio at the γ non-recrystallization temperature t or lower. Is 60% or less, and the rolling finish temperature is (austenite non-recrystallization temperature t-80 ° C.) or more (Experiment No. 23, 80), the aspect ratio of the prior austenite grains of the base material (main rolling direction) (Particle size / particle size in the plate thickness direction) can be controlled to 1.2 or less, and the target value of acoustic anisotropy is satisfied.

  Experiment No. 25 to 27 are within the range defined in the present invention except for the Si amount, and only the Si amount is changed (steel types V to X in Table 1). When Si exceeds 0.4% (Experiment No. 27), the volume ratio of the island-like martensite phase increases, and the base material toughness does not satisfy the target value.

Experiment No. Nos. 28 to 31 are within the range defined in the present invention except for the amount of Mn, and only the amount of Mn is changed (steel types Y and Z in the table and steel types A1 and B1 in Table 2). Experiment No. with Mn exceeding 1.6% and CE _N exceeding 0.33%. In the case of No. 31 (steel type B1), it can be seen that the yield ratio YR is high, and both the weld cold crack resistance and the high heat input HAZ toughness are below the target values.

  Experiment No. Nos. 32 to 34 are obtained by changing the Cr amount on the high Mn side (steel types C1, D1, E1 in Table 2). No. 32 does not satisfy the target values of yield strength and tensile strength.

Experiment No. Nos. 35 and 36 are obtained by changing the Cr amount on the low Mn side (steel types F1 and G1 in Table 2). Cr acts effectively in reducing the yield ratio in the same tensile strength comparison, but the Cr amount is 1 Experiment No. over 3% and CE _N also over 0.33%. In the case of 36 (steel type G1), the base metal toughness, the weld cold crack resistance and the high heat input HAZ toughness do not satisfy the target values.

  Experiment No. 37 to 40 are within the range defined in the present invention except for the Al amount, and only the Al amount is changed (steel types H1 to K1 in Table 2). Experiment No. with Al over 0.08%. In the case of 40 (steel type K1), it can be seen that the base metal toughness is lower than the target value.

  Experiment No. Nos. 41 to 44 are within the range defined in the present invention except for the B amount, and the B content is changed (steel types L1 to O1 in Table 2). Experiment No. in which the total amount of B exceeded 0.003%. In the case of No. 44 (steel grade O1), the yield strength (0.2% YS) and the yield ratio (YR) exceed the target values, and the base metal toughness, welding cold crack resistance, and high heat input welding HAZ toughness are the target values. You can see that you are not satisfied.

Experiment No. Nos. 45 to 47 are within the range defined in the present invention excluding the Ni amount, and the Ni amount is changed to the upper limit side (steel types P1 to R1 in Table 2). Experiment No. with Ni amount exceeding 5%. In 47 (steel type R1), the base metal toughness and the resistance to welding cold cracking are inferior.

  Experiment No. Nos. 48 to 51 and 97 are obtained by changing the TiN amount balance (steel types S1 to V1 and S2 in Table 2). Experiment No. In the case of 48 without addition of Ti, the yield strength (YR) and the high heat input HAZ toughness are lower than the target values. On the other hand, in Experiment No. in which Ti was added excessively. In 97, the base material toughness and the HAZ toughness are lower than the target values. In addition, Experiment No. In No. 51, the N content exceeds the upper limit of the formula (1) defined in the present invention (steel type V1), and the base metal toughness and the high heat input welding HAZ toughness are lower than the target values.

  Experiment No. 52 to 55 are within the range defined by the present invention except for the V amount, and the V amount is changed (steel types W1 to Z1 in Table 2). Experiment No. In No. 55, the amount of V exceeded 0.10% (steel type Z1), and it can be seen that the high heat input HAZ toughness is below the target value.

  Experiment No. Nos. 56 to 59 are obtained by changing the Ca content based on the basic components of the present invention (steel types A2 to D2 in Table 3). When the Ca content exceeds 0.01% (Experiment No. 59), it can be seen that the high heat input welding HAZ toughness tends to deteriorate and does not satisfy the target value.

  Experiment No. 60 to 63 are those in which the basic component of the present invention contains the amount of La, which is a rare earth element, changed (steel types E2 to H2 in Table 3). When the La content exceeds 0.02% (Experiment No. 63), it can be seen that the high heat input HAZ toughness deteriorates and falls below the target value.

  Experiment No. Nos. 64 to 67 are obtained by changing the Mg content based on the basic components of the present invention (steel types I2 to L2 in Table 3). If Mg exceeds 0.0030% (Experiment No. 67), the high heat input welding HAZ toughness deteriorates and falls below the target value.

  Experiment No. Nos. 68 to 71 contain La or (La + Mg) in addition to changing the Ce amount (steel types O2 to P2 in Table 3). Experiment No. with Ce over 0.0050%. In the case of No. 71, it can be seen that the high heat input welding HAZ toughness deteriorates and falls below the target value.

  Experiment No. 72 and 73 are those in which (Mg + Ce) or (La + Ce) is added within the scope of the present invention (steel types Q2 and R2 in Table 3), the toughness of the base metal, the resistance to welding cold cracking, and the high heat input welding HAZ. It can be seen that the toughness satisfies the target value.

  Experiment No. 74 and 98 are examples of the present invention of DQ-Q'-T materials having a thickness of 45 mm and a thickness of 20 mm, respectively, and satisfy the target values in all the characteristics.

  Experiment No. 75-79 change the slab heating temperature within the range of 900-1350 degreeC with respect to the component steel grade (steel grade S2) prescribed | regulated by this invention. When the heating temperature falls below the temperature specified in the present invention (Experiment No. 75), the tensile strength (TS) is below the target value and the yield ratio (YR) does not satisfy the target of 80% or less. It becomes. When the heating temperature exceeds the range specified in the present invention (Experiment No. 79), the base material toughness is low and the target value is not satisfied. On the other hand, when the heating temperature is within the range defined by the present invention (Experiment Nos. 76 to 78), the target values are satisfied in all the characteristics.

  Experiment No. 81 to 83 are obtained by changing the rolling end temperature (rolling finishing temperature) of the slab in the temperature range near the austenite non-recrystallization temperature t with respect to the component steel type (steel type T2) defined in the present invention. When the rolling end temperature reaches 1150 ° C. (Experiment No. 83), the crystal grains of the transformation structure are also coarsened, so that the base material toughness is reduced and the target value is not satisfied. On the other hand, when the rolling end temperature is within the range specified by the present invention (Experiment Nos. 81 and 82), the target material characteristics are satisfied.

  Experiment No. No. 84 is a component steel type specified in the present invention, in which the direct quenching start temperature (cooling start temperature) after rolling is lower than the lower limit temperature specified in the present invention, but the yield strength (TS), base metal Both toughness will fall below the target material properties.

  Experiment No. Nos. 85 and 86 are obtained by changing the cooling rate after the rolling with respect to the component steel types (steel type T2) defined in the present invention. When the cooling rate falls below the lower limit specified in the present invention (Experiment No. 85), none of the yield strength, tensile strength, and base metal toughness satisfies the target characteristics. On the other hand, when the cooling rate is within the range defined by the present invention (Experiment No. 86), the target material characteristics are satisfied.

  Experiment No. Nos. 87 and 88 are obtained by changing the direct quenching stop temperature (cooling stop temperature) after rolling to the high temperature side with respect to the component steel type (steel type T2) defined in the present invention. When the quenching stop temperature exceeds the range specified in the present invention (Experiment No. 88), none of the yield strength, tensile strength, and base metal toughener satisfy the target characteristics. On the other hand, when the cooling rate is within the range defined by the present invention (Experiment No. 87), the target material characteristics are satisfied.

  Experiment No. Nos. 89 to 92 are obtained by changing the two-phase quenching temperature (Q ′) with respect to the component steel types (steel type T2) defined in the present invention. When the two-phase region quenching temperature is significantly lower than the range defined in the present invention (Experiment No. 89), the yield strength is lower than the target characteristic. In addition, when the two-phase region quenching temperature greatly exceeds the range specified in the present invention (Experiment No. 92), the yield ratio (YR) does not satisfy the target of 80% or less. On the other hand, when the two-phase quenching temperature is within the range defined by the present invention (Experiment No. 90, 91), the target material characteristics are satisfied.

  Experiment No. Nos. 93 to 96 are obtained by changing the tempering temperature (T) after the two-phase region quenching with respect to the component steel type (steel type T2) defined in the present invention. When the tempering temperature is much lower than the range defined in the present invention (Experiment No. 93), the number of ε-Cu phase cluster precipitates is small, so the yield strength (0.2% YS) is lower than the target characteristics. . When the tempering temperature greatly exceeds the range specified in the present invention (Experiment No. 96), both the yield strength and the tensile strength are lower than the target values. On the other hand, when the tempering temperature is within the range defined by the present invention (Experiment No. 94, 95), the target material characteristics are satisfied.

  Experiment No. Nos. 99 and 100 are obtained by directly quenching the component steel types defined in the present invention after online leveler correction. After rolling, experiment No. 1 was directly quenched without online leveler correction. As in the case of No. 83, a steel sheet that can satisfy the target material characteristics and does not need to be cut off due to poor flatness is obtained.

As is clear from these results, in the high-tensile steel sheet obtained by the present invention, the toughness and yield ratio of the base material satisfy the target values, and there is no gas cutting crack sensitivity and acoustic anisotropy. In addition, it can be seen that the weld crack resistance is good, and the high heat input welded HAZ toughness has an average vE ₀ of 70 J or more in all parts. It should be noted that, even in a 1RUN submerged welding corner joint between columns, since it is a high heat input welding equivalent to electroslag welding, it is needless to say that the average HAZ toughness can be secured at 70J or higher in any part. The characteristics of the invention steel can be satisfied. In addition, although it does not show in the Example about the gas cutting property of the high-tensile steel plate of this invention, all had the quality of the gas cutting surface which fully satisfied the 2nd grade or more of WES2801 (1980).

Claims (5)

C: 0.015-0.045% (meaning of mass%, the same applies hereinafter), Si: 0.4% or less (including 0%), Mn: 0.8-1.6%, Cr: 0.5 ~ 1.3%, sol. Al: 0.08% or less (including 0%), B: 0.0004 to 0.003%, Cu: 0.5 to 0.95%, Ni: 0.7 to 5.0% (however, Ni The ratio [Ni] / [Cu] ≧ 1) of the content [Ni] and the Cu content [Cu], Ti: 0.005 to 0.03% and N satisfying the following formula (1) A steel material having a chemical composition that does not substantially contain Nb and Mo and has a CE _N value in the range of 0.27 to 0.33% represented by the following formula (2) is 950 to 1300 ° C. Then, the cumulative rolling reduction in the temperature range below the austenite non-recrystallization temperature t (° C.) represented by the following formula (3) is set to 60% or less (the austenite non-recrystallization temperature t− 80 ° C.) or more and 1100 ° C. or less after rolling, and then at a cooling rate of 3 ° C./second or more from a temperature of 780 ° C. or more to 300 ° C. It is hardened directly until it becomes below, followed by reheating degree quenching in the temperature range of 760 to 840 ° C. and then tempering in the temperature range of 450 to 550 ° C. A method for producing a low-yield ratio high-tensile steel sheet having excellent heat input weld joint toughness and small acoustic anisotropy.
[Ti] × 14.0 / 47.9−0.001 ≦ [N] ≦ [Ti] × 14.0 / 47.9 + [B] × 14.0 / 10.8 (1)
However, [Ti], [N], and [B] indicate the contents (mass%) of Ti, Ni, and B, respectively.
CE _N = [C] + A (c) ・ [[Si] / 24 + [Mn] / 6 + [Cu] / 15 + [Ni] / 20 + ([Cr] + [V]) / 5 + 5 [B]] (2)
However, A (c) = 0.75 + 0.25 · tanh [20 ([C] -0.12)], [C], [Si], [Mn], [Cu], [Ni], [Cr], [ V] and [B] indicate the contents (mass%) of C, Si, Mn, Cu, Ni, Cr, V, and B, respectively.
T (℃) = 887 + 464 [C] + (732 × [V] -230 × √ [V]) + 890 × [Ti] +363 [sol.Al] -357 × [Si]
(3)
However, [C], [V], [Ti], [sol. Al] and [Si] are C, V, Ti, sol. The content (mass%) of Al and Si is shown.   The method for producing a low-yield ratio high-tensile steel sheet according to claim 1, wherein after the rolling is finished, online leveler correction is performed.   Furthermore, the manufacturing method of the low yield ratio high tension steel plate of Claim 1 or 2 which uses the steel raw material containing V: 0.005-0.10%.   Furthermore, the manufacturing method of the low yield ratio high-tensile steel plate in any one of Claims 1-3 which uses the steel raw material containing 0.0005-0.01%.   Furthermore, containing containing 1 type (s) or 2 or more types chosen from the group which consists of La: 0.002-0.02%, Ce: 0.0003-0.0050% and Mg: 0.0005-0.0030% The method for producing a low yield ratio high strength steel sheet according to any one of claims 1 to 4, wherein a steel material to be used is used.