WO2021199629A1 - Steel sheet and method for manufacturing same - Google Patents

Steel sheet and method for manufacturing same Download PDF

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Publication number
WO2021199629A1
WO2021199629A1 PCT/JP2021/002892 JP2021002892W WO2021199629A1 WO 2021199629 A1 WO2021199629 A1 WO 2021199629A1 JP 2021002892 W JP2021002892 W JP 2021002892W WO 2021199629 A1 WO2021199629 A1 WO 2021199629A1
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ferrite
temperature
toughness
steel sheet
steel
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PCT/JP2021/002892
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French (fr)
Japanese (ja)
Inventor
竜平 竹下
亮 荒尾
植田　圭治
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Ｊｆｅスチール株式会社
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Priority to CN202180015241.1A priority Critical patent/CN115135787A/en
Priority to JP2021524994A priority patent/JP7276443B2/en
Priority to KR1020227019101A priority patent/KR20220092977A/en
Publication of WO2021199629A1 publication Critical patent/WO2021199629A1/en

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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

Definitions

  • the present invention has a thickness capable of ensuring high toughness even in steel plates used for ships, marine structures, mid-to-high-rise buildings, bridges, tanks, etc., particularly in weld heat-affected zones (hereinafter, also referred to as HAZ) when welding is performed. It is about steel plates.
  • Patent Document 1 As a method for improving the toughness of HAZ (hereinafter, also referred to as high heat-affected zone HAZ) by high heat-affected zone, for example, in Patent Document 1 and Patent Document 2, coarsening of austenite grains is suppressed by a pinning effect of TiN, Al oxide, etc. A method has been proposed. Further, Patent Document 3, Patent Document 4 and Patent Document 5 show a technique for refining the structure in crystal grains by allowing a large number of ferrite transformation nuclei to exist in austenite grains. Specifically, by using TiN, MnS, Ti oxide and the like as ferrite transformation nuclei, the microstructure in the crystal grains is miniaturized and the low temperature toughness of HAZ is improved. Further, in Patent Document 6, the HAZ toughness is improved by utilizing the solid solution B and suppressing the ratio of the grain boundary ferrite. In Patent Document 7, the reproduced HAZ structure is improved by making the bainite structure in the grain finer by using the compound B.
  • the inventors focused on a coarse ferrite side plate, which is a low toughness structure produced by high heat input welding.
  • the coarse ferrite side plate (hereinafter referred to as FSP) is a structure formed by extending ferrite into grains starting from the coarse grain boundary ferrite generated from the coarse austenite grain boundaries as described above.
  • the roughness of this FSP structure is the main factor of low toughness. Therefore, the inventors considered that by refining the coarse grain boundary ferrite, the formation of coarse FSP is suppressed and the low temperature toughness of the large heat-affected zone HAZ is improved.
  • the SB defined by the following equation (1) satisfies the predetermined conditions, and the temperature obtained by the following equation (2) is from Ar 3 points (transformation start temperature). It was found that by designing the component composition so that the temperature becomes high, grain boundary ferrite is nucleated from the BN precipitated at the grain boundary, and the grain boundary ferrite can be miniaturized. By refining the grain boundary ferrite, it is possible to obtain the low temperature toughness of the large heat-affected zone HAZ, which is superior to the conventional one.
  • SB [B] -0.77 x [N] + 0.22 x [Ti] ... (1)
  • T (° C.) 12000 / (4.63-log ([B] ⁇ ([N]-[Ti] /3.42)))-273 ... (2)
  • B, N, and Ti are the contents (mass%) of each element.
  • the present invention has been made based on the above findings, and its gist structure is as follows. 1. 1. By mass% C: 0.03 to 0.15%, Si: 0.01-0.50%, Mn: 1.20 to 2.00%, P: 0.020% or less, S: 0.0005 to 0.0100%, Al: 0.005 to 0.100%, Ti: 0.004 to 0.030%, B: 0.0020 to 0.0050% and N: 0.0035 to 0.0100% Is contained in the range where SB represented by the following formula (1) is ⁇ 0.0010 or more and 0.0002 or less and the temperature T represented by the following formula (2) is more than 3 points of Ar, and the balance is Fe and A steel sheet having a component composition that is an unavoidable impurity and having a metal structure having a volume fraction of processed ferrite of 50% or more.
  • each element in the formula shows the content (mass%) of the element.
  • composition of the components is further increased by mass%.
  • the steel material of the present invention is suitably used for structures such as low-temperature storage tanks for liquefied gas constructed by large heat input such as electrogas welding, submerged arc welding, and electroslag welding, and ships operated in a low temperature environment. Be done.
  • C 0.03 to 0.15% C must contain 0.03% or more in order to obtain the required strength. However, if the content exceeds 0.15%, island-like martensite increases and the toughness of the weld heat-affected zone decreases, so the upper limit is set to 0.15%.
  • the lower limit is preferably 0.045%. Further, it is preferably less than 0.10%.
  • Si 0.01-0.50% Si is a component necessary for ensuring the strength of the base material, deoxidizing, etc., and is added in an amount of 0.01% or more.
  • the upper limit is set to 0.50%.
  • the preferred lower limit is 0.10% and the preferred upper limit is 0.30%.
  • Mn 1.20 to 2.00% Mn is required to be 1.20% or more in order to secure the strength of the base material, and if it exceeds 2.00%, not only the weldability deteriorates but also the steel material cost increases. Therefore, the range of Mn is 1.20 to 2.00%.
  • the lower limit is preferably 1.40%.
  • the upper limit is preferably 1.60%.
  • P 0.020% or less
  • P is an impurity that is inevitably mixed in, and if the content exceeds 0.020%, the toughness of the base metal and welds will decrease, so the upper limit is set to 0.020%. .. In order to obtain good toughness, 0.010% or less is preferable, and 0.007% or less is more preferable. By the way, although it is not necessary to limit the lower limit, it is preferable to set it to 0.001% or more because the cost increases by performing the ultra-low P treatment.
  • S 0.0005 to 0.0100%
  • S is required to be 0.0005% or more in order to generate the required MnS in the nucleus of the composite inclusion required for ferrite nucleation, and CaS when Ca is added.
  • S is less than 0.0005%, MnS and further CaS are not sufficiently formed, and the toughness of HAZ is lowered.
  • the upper limit is preferably 0.0090%.
  • the lower limit is preferably 0.0010%.
  • Al 0.005 to 0.100% Al needs to be 0.005% or more, preferably 0.010% or more in terms of deoxidation of steel. On the other hand, if it is contained in excess of 0.100%, the toughness of the base metal is lowered and the toughness of the weld metal is deteriorated.
  • the upper limit is preferably 0.08%.
  • Ti 0.004 to 0.030% Ti precipitates as TiN during solidification of steel, and contributes to suppressing coarse-grained austenite in the weld heat-affected zone (HAZ) and becoming ferrite transformation nuclei to increase toughness. If Ti is less than 0.004%, its effect is small, while if it exceeds 0.030%, the expected effect cannot be obtained due to the coarsening of TiN particles. Therefore, the Ti content is in the range of 0.004 to 0.030%. The lower limit is preferably 0.008%. The upper limit is preferably 0.020%.
  • B 0.0020-0.0050% B is an important element for refining grain boundary ferrite and improving HAZ toughness, and is added at least 0.0020% in order to precipitate at a ferrite transformation temperature or higher. However, if a large amount is added, the toughness of the base metal deteriorates, so the upper limit is set to 0.0050%.
  • the lower limit is preferably 0.0025%.
  • the upper limit is preferably 0.0040%.
  • N 0.0035-0.0100% N is added in an amount of 0.0035% or more in order to combine with Ti to form TiN and to combine with B to form BN. That is, when N is below the lower limit of 0.0035%, BN is not formed and sufficient HAZ toughness cannot be secured. On the other hand, when the content of N increases, the solid solution N increases and the HAZ toughness decreases, so the upper limit is 0.0100%.
  • the lower limit is preferably 0.0040%.
  • the upper limit is preferably 0.0090%.
  • the steel sheet of the present invention contains each of the above components, and the balance has a component composition of Fe and unavoidable impurities.
  • B, N and Ti are contained in the above formulas (1) and (2) so as to satisfy the above-mentioned formulas (1) and (2), so that the heat cycle received by the steel sheet during large heat input welding (hereinafter, also referred to as a welding heat cycle). ), TiN remains without solidification, and BN is deposited at an early stage with this TiN as a nucleus.
  • FIG. 1 an observation image of a sample in which a steel sheet having the above-mentioned composition composition is subjected to a welding reproduction heat cycle equivalent to 10 kJ / mm of heat input is shown. It can be seen that BN is precipitated in. That is, BN is more likely to precipitate from the high temperature region.
  • the size of the composite precipitate of TiN and BN becomes larger than the size of TiN alone.
  • Increasing the size of the precipitate facilitates nucleation of ferrite.
  • the size of the core TiN is usually 15 nm or more and 200 nm or less, and when BN precipitates on TiN, the size of the BN-coated precipitate becomes 50 nm or more and 600 nm or less.
  • the fact that ferrite is easily nucleated means that many ferrite nuclei are formed at the grain boundaries, and many ferrites are formed at the grain boundaries. Since these ferrites are nucleated from different BNs and therefore have different orientations, the crystal orientations of the ferrites are randomized.
  • the grain boundary ferrite is miniaturized, and the ferrite side plate generated from the grain boundary ferrite is also miniaturized. Therefore, the HAZ toughness is improved by satisfying the formulas (1) and (2).
  • the above formula (2) shows the precipitation temperature T when BN is deposited around TiN as shown in FIG. 1, and when this T becomes Ar 3 points or less, ferrite having BN as a core is shown. As a result of difficulty in formation, miniaturization of grain boundary ferrite is not realized.
  • the heat-affected zone structure near the bond is the density of grain boundary ferrite generated at the old ⁇ grain boundaries. Is 20 pieces / mm or more.
  • the grain boundary ferrite formation density of the old ⁇ grain boundaries is measured by performing quenching treatment immediately after the start of ferrite transformation during cooling in a thermal cycle simulation simulating welding and using EBSD (electron backscatter diffraction method). Can be done.
  • the curve length along the adjacent 3 to 3 priority grain boundaries of the old ⁇ grain boundary is defined as the old ⁇ grain boundary length, and the crystals of adjacent ferrite grains generated on the old ⁇ grain boundary are used.
  • the number of ferrite grains with an orientation difference of 15 degrees or more is defined as the number of ferrites on the old ⁇ grain boundaries, and the density of grain boundary ferrites is defined by (number of ferrites on the old ⁇ grain boundaries) / (former ⁇ grain boundary length). do.
  • the density of grain boundary ferrites formed on the old ⁇ grain boundaries in the heat-affected zone structure near the bond becomes 20 grains / mm or more when the above-mentioned large heat-immersive welding is performed. Therefore, it is possible to suppress the formation of coarse ferrite side plates and realize excellent low temperature toughness in HAZ.
  • the heat-affected zone structure in the vicinity of the bond refers to a region from the boundary of the weld metal with the base steel plate to a position within about 0.5 mm on the steel plate side of the base material.
  • the density of grain boundary ferrite generated at the old ⁇ grain boundaries is determined by controlling the addition amounts of N, B and Ti within the specified range according to the above formulas (1) and (2), for example, the heat input amount is 5 kJ / mm.
  • the density of grain boundary ferrites when the above-mentioned large heat input welding is performed can be 20 pieces / mm or more. That is, the formation of coarse ferrite side plates is suppressed, and excellent toughness can be obtained in the heat-affected zone.
  • the metal structure of the steel sheet according to the present invention has a volume fraction of 50% or more of processed ferrite in the structure.
  • the processed ferrite refers to a ferrite having a dislocation density ⁇ of 1.0 ⁇ 10 14 m- 2 or more, which is determined by X-ray diffraction (XRD). That is, in the processed ferrite, high-density dislocations are introduced, and the dislocations interact with each other to hinder each other's movements, thereby increasing the strength. Then, by setting the volume fraction of the processed ferrite to 50% or more, the strength is increased.
  • the volume fraction of processed ferrite in the metal structure is preferably 60% or more.
  • the upper limit of the amount of processed ferrite is not particularly limited and may be 100%, but from the viewpoint of the capacity of the rolling mill, it is preferably 90% or less.
  • the remaining structure at that time is preferably one or more hard phases of pearlite, bainite and martensite.
  • two or more kinds can be arbitrarily contained.
  • Cu 0.01-0.50%
  • Cu is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal. Therefore, the Cu content is preferably 0.01 to 0.50%.
  • Ni 0.01-1.50%
  • Ni is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal, and also increases the cost of the alloy. Therefore, the Ni content is preferably 0.01 to 1.50%.
  • Nb 0.005 to 0.040%
  • Nb is an element effective for ensuring the strength, toughness and joint strength of the base metal. The effect is exhibited when the content is 0.005% or more. On the other hand, if it is contained in excess of 0.040%, the toughness deteriorates due to the formation of island-shaped martensite in the weld heat affected zone. Therefore, when Nb is contained, the Nb content is preferably 0.005 to 0.040%.
  • V acts as a ferrite nucleation nucleus as a VN and improves the strength and toughness of the base metal. This effect is exhibited by containing 0.005% or more of V. On the other hand, if V is contained in an amount of more than 0.100%, the toughness of the base metal is rather lowered. Therefore, when V is contained, the V content is preferably 0.005 to 0.100%.
  • Cr 0.01-0.50%
  • Cr is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal. Therefore, the Cr content is preferably 0.01 to 0.50%.
  • Mo 0.01-0.50%
  • Mo is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal. Therefore, the Mo content is preferably 0.01 to 0.50%.
  • Ca 0.0005 to 0.0030% Ca is an element useful for improving the toughness of the base metal by fixing S, but the effect is saturated when the content exceeds 0.0030%, so Ca should be contained at 0.0030% or less. .. On the other hand, if the content is less than 0.0005%, the fixation of S becomes insufficient. Therefore, the Ca content is preferably 0.0005% or more and 0.0030% or less.
  • Mg 0.0002 to 0.0050% REM: 0.0010 to 0.1000%
  • Both Mg and REM have a strong deoxidizing power in molten steel and have a function of assisting the formation of fine oxides, and therefore are added as necessary.
  • the addition amounts showing the deoxidizing effect are Mg: 0.0002% or more and REM: 0.0010% or more, respectively.
  • Mg 0.0002% or more
  • REM 0.0010% or more
  • a steel material having the above composition is heated to a temperature of 1050 ° C. or higher and 1200 ° C. or lower, cooled to a temperature of 900 ° C. or lower at a cooling rate of 7 ° C./s or lower, and then cooled to 850 ° C. or lower.
  • Hot rolling is performed in which the cumulative reduction rate of ferrite-austenite in the two-phase temperature range is 60% or more and the finishing temperature is 650 ° C. or more.
  • the heating temperature of the steel material for example, the slab, needs to be 1050 ° C. or higher and 1200 ° C. or lower.
  • the reason for this is that heating below 1050 ° C. may leave coarse inclusions that adversely affect the toughness produced during solidification undissolved.
  • the precipitates formed by controlling the cooling rate described later may be redissolved.
  • 1200 ° C. or lower is sufficient as the heating temperature in the sense of completing the phase transformation. It should be noted that the coarsening of crystal grains that is considered to occur at that time can also be prevented in advance by the pinning effect of TiN described above. From the above, the heating temperature was limited to 1050 ° C. or higher and 1200 ° C. or lower.
  • the cooling rate is preferably 1 ° C./s or higher from the viewpoint of production efficiency.
  • the cumulative reduction rate in the two-phase temperature range is 60% or more, dislocations are added to the ferrite in the two-phase temperature range, and as a result, the strength can be improved.
  • the cumulative reduction rate is 60% or more, the rolled texture of ferrite develops, which contributes to the improvement of low temperature toughness.
  • the cumulative reduction rate of ferrite + austenite in the two-phase temperature range of 850 ° C. or lower was limited to 60% or more.
  • the cumulative rolling reduction ratio is preferably 90% or less from the viewpoint of rolling functional force.
  • the finishing temperature in hot rolling is set to 650 ° C. or higher. This is because if the finish rolling is performed at a temperature lower than 650 ° C., the ferrite produced by the phase transformation is distorted more than necessary, and the toughness is lowered.
  • cooling from a temperature of 650 ° C. or higher to a temperature range of 600 ° C. or lower and 300 ° C. or higher at a cooling rate of 5 ° C./s or higher is used to increase the strength of the base metal.
  • the steel material is cooled from a temperature of 650 ° C. or higher to a cooling rate of 5 ° C./s or higher to a temperature range of 600 ° C. or lower and 300 ° C. or higher after hot rolling is completed at 650 ° C. or higher. That is, the reason for cooling from 650 ° C.
  • cooling is started at a temperature lower than 650 ° C., the hardenability becomes insufficient and the required strength may not be obtained. Further, if the cooling rate is less than 5 ° C./s, it becomes difficult to obtain a steel having a uniform microstructure. Further, it is preferable to cool to a temperature range of 600 ° C. or lower and 300 ° C. or higher. This is because it is difficult to secure sufficient strength from the viewpoint of hardenability when cooling is stopped at a temperature exceeding 600 ° C. In addition, stopping cooling at a temperature of less than 300 ° C. does not significantly change the characteristics of the steel material, so that only the operational load increases.
  • the steel pieces are cooled from a temperature of 650 ° C. or higher to a cooling rate of 5 ° C./s or higher to 600 ° C. or lower and 300 ° C. or higher after completing hot ductility at 650 ° C. or higher.
  • the cooling rate is preferably 50 ° C./s or higher from the viewpoint of ensuring the toughness of the base material.
  • the cooling rate during slab casting is set to 0.3 m / min or more and 1.0 m / min or less. If the casting speed is less than 0.3 m / min, the size of TiN of the base metal (steel plate) becomes large. As the TiN size increases, the TiN density of the base material (steel plate) may decrease and the amount of BN composite precipitates may decrease. As a result, the ferrite cannot be sufficiently miniaturized, and the HAZ toughness may deteriorate.
  • the size of the core TiN is 15 nm or more and 200 nm or less.
  • the steel sheet thus produced has a structure having a volume fraction of processed ferrite of 50% or more in addition to the above-mentioned component composition.
  • the main phase contains a soft phase made of ferrite, and the balance is a structure made of one or more hard phases of pearlite, bainite and martensite.
  • the main phase when the main phase is ferrite, it means that ferrite has a volume fraction of 60% or more. That is, the ferrite may be 100%, but it is preferably 90% or less from the viewpoint of rollability.
  • the remaining portion at that time does not need to be particularly limited, and is as described above, for example. What is important here is the ratio of processed ferrite to the structure among the ferrites, and the ratio should be 50% or more in terms of volume fraction. Therefore, ferrites other than processed ferrites, that is, ferrites having a dislocation density ⁇ of less than 1.0 ⁇ 10 14 m- 2 may be contained.
  • the yield stress is 325 MPa or more. Further, it is desirable that the Charpy impact absorption energy of the base material at ⁇ 70 ° C. is 200 J or more. Further, it is desirable that the Charpy impact absorption energy at ⁇ 70 ° C. of the joint subjected to the large heat input welding is 80 J or more.
  • a steel slab (steel material) adjusted to the composition shown in Table 1 is cooled after heating the slab according to various conditions shown in Table 2, and then hot-rolled and cooled to obtain a thick steel sheet having a thickness of 20 mm. And said.
  • Tensile test pieces conforming to JIS Z2241 were collected from each of the thick steel sheets thus obtained, and a tensile test conforming to JIS Z2241 was performed to measure the yield stress.
  • JIS Z2242 compliant test pieces are collected from each thick steel plate, V groove is processed on each test piece, and a Charpy impact test compliant with JIS Z2242 is performed to measure Charpy impact absorption energy at -70 ° C. bottom.
  • a test piece for welding a welded joint was collected from each of the obtained thick steel plates, a V groove was machined on the test piece, and a welded joint was manufactured by submerged arc welding (welding heat input: 102 kJ / cm).
  • a JIS No. 4 impact test piece having a notch position as a bond portion was collected from these welded joints, a Charpy impact test was carried out, and the Charpy impact absorption energy at ⁇ 70 ° C. was measured.

Abstract

The present invention provides a steel sheet in which, in particular, the low-temperature toughness of a high heat input HAZ has been improved. The present invention has a component composition that contains, in mass% and in prescribed relationships, C: 0.03-0.15%, Si: 0.01-0.50%, Mn: 1.20-2.00%, P: 0.020% or less, S: 0.0005-0.0100%, Al: 0.005-0.100%, Ti: 0.004-0.030%, B: 0.0020-0.0050%, and N: 0.0035-0.0100%, and the processed ferrite volume fraction in the metal structure is set to 50% or higher.

Description

鋼板およびその製造方法Steel plate and its manufacturing method
　本発明は、船舶、海洋構造物、中高層ビル、橋梁、タンクなどに使用される鋼板、特に、溶接を行った際の溶接熱影響部（以下、ＨＡＺとも称する）においても高い靭性を確保できる厚鋼板に関するものである。 INDUSTRIAL APPLICABILITY The present invention has a thickness capable of ensuring high toughness even in steel plates used for ships, marine structures, mid-to-high-rise buildings, bridges, tanks, etc., particularly in weld heat-affected zones (hereinafter, also referred to as HAZ) when welding is performed. It is about steel plates.
　近年、船舶、海洋構造物、中高層ビル、橋梁、タンクなどの構造物に使用される、溶接用鋼材の材質特性に対する要望は厳しさを増している。さらに、そのような構造物は、短期間で製造するべく、サブマージアーク溶接法、エレクトロガス溶接法、エレクトロスラグ溶接法などに代表される、大入熱溶接法の運用が希望されていることから、鋼材自身の靭性と同様に、ＨＡＺの靭性への要求も厳しさを増している。しかし、一般に、溶接入熱量が大きくなると、ＨＡＺの組織が粗大化し、ＨＡＺの靭性は低下することが知られている。このような大入熱溶接による靭性の低下に対して、これまでも多くの対策が提案されてきた。 In recent years, there have been increasing demands for the material characteristics of welding steel materials used in structures such as ships, marine structures, mid-to-high-rise buildings, bridges, and tanks. Furthermore, in order to manufacture such structures in a short period of time, it is desired to operate a large heat-affected welding method represented by a submerged arc welding method, an electrogas welding method, an electroslag welding method, and the like. As with the toughness of steel itself, the demand for toughness of HAZ is becoming more stringent. However, it is generally known that when the amount of heat input to welding increases, the structure of HAZ becomes coarse and the toughness of HAZ decreases. Many measures have been proposed to deal with the decrease in toughness caused by such large heat input welding.
　大入熱溶接によるＨＡＺ（以下、大入熱ＨＡＺともいう）の靱性を改善する方法として、例えば特許文献１および特許文献２では、ＴｉＮ、Ａｌオキサイド等のピンニング効果によりオーステナイト粒の粗大化を抑制する方法が提案されている。また、特許文献３、特許文献４および特許文献５では、オーステナイト粒内にフェライト変態核を多数存在させることにより結晶粒内組織の微細化を図る技術が示されている。具体的には、ＴｉＮ、ＭｎＳ、Ｔｉオキサイド等をフェライト変態核として利用することにより、結晶粒内組織の微細化を達成し、ＨＡＺの低温靱性の改善を図っている。また、特許文献６では、固溶Ｂを活用し粒界フェライトの割合を抑えることでＨＡＺ靱性の改善を図っている。特許文献７では、Ｂの化合物を用いて、粒内のベイナイト組織を細粒化することで再現ＨＡＺ組織の改善を図っている。 As a method for improving the toughness of HAZ (hereinafter, also referred to as high heat-affected zone HAZ) by high heat-affected zone, for example, in Patent Document 1 and Patent Document 2, coarsening of austenite grains is suppressed by a pinning effect of TiN, Al oxide, etc. A method has been proposed. Further, Patent Document 3, Patent Document 4 and Patent Document 5 show a technique for refining the structure in crystal grains by allowing a large number of ferrite transformation nuclei to exist in austenite grains. Specifically, by using TiN, MnS, Ti oxide and the like as ferrite transformation nuclei, the microstructure in the crystal grains is miniaturized and the low temperature toughness of HAZ is improved. Further, in Patent Document 6, the HAZ toughness is improved by utilizing the solid solution B and suppressing the ratio of the grain boundary ferrite. In Patent Document 7, the reproduced HAZ structure is improved by making the bainite structure in the grain finer by using the compound B.
特公昭５５－０２６１６４号公報Special Publication No. 55-026164 特許第２９５００７６号Patent No. 2950076 特公平０７－０６８５７７号公報Special Fair 07-068577 Gazette 特公平０５－０１７３００号公報Special Fair 05-017300 Gazette 特許第３７３３８９８号Patent No. 3733898 特開２００５－３３６６０２号公報Japanese Unexamined Patent Publication No. 2005-336602 特許第４３３２０６４号Patent No. 4332064
　しかしながら、上記の析出物を利用してＨＡＺを微細化する諸技術を適用しても、大入熱溶接を施す場合には、ＨＡＺ組織の粗大化は不可避であり、例えば－４０℃を下回る環境下では低温靭性の劣化が生じる。近年、船舶やタンク等においては、従来よりも低温の環境での運用が検討されており、上記の特許各文献に記載の技術が対象としている鋼材よりも飛躍的に溶接熱影響部の低温靭性を向上させた鋼材が必要とされるようになっている。そこで、本発明は、上記実情を鑑み、特に大入熱ＨＡＺの低温靱性が向上された鋼板を提供することを目的とするものである。 However, even if various techniques for refining HAZ using the above precipitates are applied, coarsening of the HAZ structure is unavoidable when performing large heat-affected welding, for example, an environment below −40 ° C. Below, deterioration of low temperature toughness occurs. In recent years, in ships and tanks, operation in a lower temperature environment than before has been studied, and the low temperature toughness of the weld heat affected zone is dramatically higher than that of the steel materials covered by the technologies described in the above patent documents. There is a growing need for steel materials with improved quality. Therefore, in view of the above circumstances, it is an object of the present invention to provide a steel sheet in which the low temperature toughness of the large heat-affected zone HAZ is particularly improved.
　発明者らは、上記課題を解決するために、大入熱ＨＡＺの低温靱性を向上するための手法について鋭意研究を重ねた結果、以下の知見を得るに到った。
　まず、発明者らは、大入熱溶接により生成する低靭性組織である粗大なフェライトサイドプレートに着目した。大入熱溶接を施したとき、オーステナイト粒が粗大に成長すると、そこから生成する組織も粗大となる。粗大なフェライトサイドプレート（以下、ＦＳＰと示す）は、上記のような粗大なオーステナイト粒界から生成した粗大な粒界フェライトを起点として、フェライトが粒内に伸長して形成される組織である。このＦＳＰ組織の粗さが低靭性の主要因である。そこで、発明者らは、粗大な粒界フェライトを微細化することにより、粗大なＦＳＰの生成が抑制されて、大入熱ＨＡＺの低温靭性が向上すると考えた。 As a result of intensive research on a method for improving the low temperature toughness of the large heat-affected zone HAZ in order to solve the above problems, the inventors have obtained the following findings.
First, the inventors focused on a coarse ferrite side plate, which is a low toughness structure produced by high heat input welding. When austenite grains grow coarsely when high heat welding is performed, the structure formed from the austenite grains also becomes coarse. The coarse ferrite side plate (hereinafter referred to as FSP) is a structure formed by extending ferrite into grains starting from the coarse grain boundary ferrite generated from the coarse austenite grain boundaries as described above. The roughness of this FSP structure is the main factor of low toughness. Therefore, the inventors considered that by refining the coarse grain boundary ferrite, the formation of coarse FSP is suppressed and the low temperature toughness of the large heat-affected zone HAZ is improved.
　さらに、発明者らが鋭意検討した結果、次の(１)式で定義されるＳＢが所定の条件を満足し、次の（２）式で得られる温度がＡｒ３点（変態開始温度）よりも高い温度となる、成分組成に設計することによって、粒界に析出したＢＮから粒界フェライトが核生成し、粒界フェライトの微細化が達成できることを見出した。この粒界フェライトの微細化により、従来よりも優れた大入熱ＨＡＺの低温靱性を得ることができる。
　ＳＢ＝［Ｂ］－０．７７×［Ｎ］＋０．２２×［Ｔｉ］　　…（１）
　Ｔ（℃）＝１２０００／（４．６３―ｌｏｇ（［Ｂ］×（［Ｎ］－［Ｔｉ］／３．４２）））－２７３　　…（２）
　ただし、Ｂ、Ｎ、Ｔｉは、各元素の含有量（質量％）である。 Furthermore, as a result of diligent studies by the inventors, the SB defined by the following equation (1) satisfies the predetermined conditions, and the temperature obtained by the following equation (2) is from Ar 3 points (transformation start temperature). It was found that by designing the component composition so that the temperature becomes high, grain boundary ferrite is nucleated from the BN precipitated at the grain boundary, and the grain boundary ferrite can be miniaturized. By refining the grain boundary ferrite, it is possible to obtain the low temperature toughness of the large heat-affected zone HAZ, which is superior to the conventional one.
SB = [B] -0.77 x [N] + 0.22 x [Ti] ... (1)
T (° C.) = 12000 / (4.63-log ([B] × ([N]-[Ti] /3.42)))-273 ... (2)
However, B, N, and Ti are the contents (mass%) of each element.
　本発明は、上記の知見に基づいてなされたものであり、その要旨構成は以下の通りである。
１．質量％で、
　Ｃ：０．０３～０．１５％、
　Ｓｉ：０．０１～０．５０％、
　Ｍｎ：１．２０～２．００％、
　Ｐ：０．０２０％以下、
　Ｓ：０．０００５～０．０１００％、
　Ａｌ：０．００５～０．１００％、
　Ｔｉ：０．００４～０．０３０％、
　Ｂ：０．００２０～０．００５０％および
　Ｎ：０．００３５～０．０１００％
を、次式（１）で示されるＳＢが－０．００１０以上０．０００２以下および、次式（２）で示される温度ＴがＡｒ３点超である範囲にて含有し、残部はＦｅおよび不可避不純物である成分組成を有し、加工フェライトの体積分率が５０％以上である金属組織を有する、鋼板。
　ＳＢ＝［Ｂ］－０．７７×［Ｎ］＋０．２２×［Ｔｉ］　　…（１）
　Ｔ（℃）＝１２０００／（４．６３―ｌｏｇ（［Ｂ］×（［Ｎ］－［Ｔｉ］／３．４２）））－２７３　…（２）
　但し、前記式（１）および（２）における、［Ｂ］、［Ｎ］および［Ｔｉ］は各成分の含有量（質量％）を表す。 The present invention has been made based on the above findings, and its gist structure is as follows.
1. 1. By mass%
C: 0.03 to 0.15%,
Si: 0.01-0.50%,
Mn: 1.20 to 2.00%,
P: 0.020% or less,
S: 0.0005 to 0.0100%,
Al: 0.005 to 0.100%,
Ti: 0.004 to 0.030%,
B: 0.0020 to 0.0050% and N: 0.0035 to 0.0100%
Is contained in the range where SB represented by the following formula (1) is −0.0010 or more and 0.0002 or less and the temperature T represented by the following formula (2) is more than 3 points of Ar, and the balance is Fe and A steel sheet having a component composition that is an unavoidable impurity and having a metal structure having a volume fraction of processed ferrite of 50% or more.
SB = [B] -0.77 x [N] + 0.22 x [Ti] ... (1)
T (° C.) = 12000 / (4.63-log ([B] × ([N]-[Ti] /3.42)))-273 ... (2)
However, [B], [N] and [Ti] in the formulas (1) and (2) represent the content (mass%) of each component.
　ここで、前記Ａｒ３点は、例えば、
　Ａｒ３（℃）＝９１０－２７３×Ｃ－７４×Ｍｎ－５７×Ｎｉ－１６×Ｃｒ－９×Ｍｏ－５×Ｃｕ
で求めることが可能である。なお、式における各元素は、該元素の含有量（質量％）を示す。 Here, the three Ar points are, for example,
Ar 3 (° C.) = 910-273 x C-74 x Mn-57 x Ni-16 x Cr-9 x Mo-5 x Cu
It is possible to find it at. In addition, each element in the formula shows the content (mass%) of the element.
２．前記成分組成は、さらに、質量％で、
　Ｃｕ：０．０１～０．５０％、
　Ｎｉ：０．０１～１．５０％、
　Ｎｂ：０．００５～０．０４０％、
　Ｖ：０．００５～０．１００％、
　Ｃｒ：０．０１～０．５０％、
　Ｍｏ：０．０１～０．５０％、
　Ｃａ：０．０００５～０．００３０％、
　Ｍｇ：０．０００２～０．００５０％および
　ＲＥＭ：０．００１０～０．１０００％
の１種または２種以上を含有する前記１に記載の鋼板。 2. The composition of the components is further increased by mass%.
Cu: 0.01-0.50%,
Ni: 0.01-1.50%,
Nb: 0.005 to 0.040%,
V: 0.005 to 0.100%,
Cr: 0.01-0.50%,
Mo: 0.01-0.50%,
Ca: 0.0005 to 0.0030%,
Mg: 0.0002 to 0.0050% and REM: 0.0010 to 0.1000%
The steel sheet according to 1 above, which contains one or more of the above.
３．前記１または２に記載の成分組成を有する鋼素材を１０５０℃以上１２００℃以下の温度に加熱後、９００℃以下の温度まで７℃／ｓ以下の冷却速度で冷却後、８５０℃以下のフェライト＋オーステナイトの二相温度域における累積圧下率が６０％以上および、仕上温度が６５０℃以上である熱間圧延を施す、鋼板の製造方法。 3. 3. After heating the steel material having the component composition described in 1 or 2 to a temperature of 1050 ° C. or higher and 1200 ° C. or lower, cooling to a temperature of 900 ° C. or lower at a cooling rate of 7 ° C./s or lower, ferrite + of 850 ° C. or lower A method for producing a steel sheet, which is subjected to hot rolling in which the cumulative reduction rate of austenite in a two-phase temperature range is 60% or more and the finishing temperature is 650 ° C. or more.
４．前記熱間圧延を施した後、６５０℃以上の温度から５℃／ｓ以上の冷却速度で６００℃以下３００℃以上の温度域まで冷却する、前記３に記載の鋼板の製造方法。 4. The method for producing a steel sheet according to the above 3, wherein after the hot rolling is performed, the steel sheet is cooled from a temperature of 650 ° C. or higher to a temperature range of 600 ° C. or lower and 300 ° C. or higher at a cooling rate of 5 ° C./s or higher.
　本発明により、大入熱溶接を施しても溶接熱影響部の低温靱性に優れる鋼材を得ることができる。したがって、本発明の鋼材は、エレクトロガス溶接、サブマージアーク溶接、エレクトロスラグ溶接などの大入熱により施工される液化ガスの低温貯蔵タンクや低温環境で運用される船舶等の構造物に好適に用いられる。 According to the present invention, it is possible to obtain a steel material having excellent low temperature toughness in the weld heat affected zone even if large heat input welding is performed. Therefore, the steel material of the present invention is suitably used for structures such as low-temperature storage tanks for liquefied gas constructed by large heat input such as electrogas welding, submerged arc welding, and electroslag welding, and ships operated in a low temperature environment. Be done.
入熱１０kJ/mm相当のサブマージアーク溶接の溶融線（ＦＬ）近傍の熱履歴を付与した再現ＨＡＺ部の析出物の観察画像を示す透過電子顕微鏡写真である。It is a transmission electron micrograph which shows the observation image of the deposit of the reproduced HAZ part which gave the heat history in the vicinity of the fusion line (FL) of the submerged arc welding corresponding to the heat input of 10 kJ / mm.
　以下に、本発明を実施するための形態について説明する。まず、本発明において化学成分を限定した意義について説明する。なお、本発明において、化学成分に関する「％」表示は、特に断らない限り「質量％」を意味している。 Hereinafter, a mode for carrying out the present invention will be described. First, the significance of limiting the chemical composition in the present invention will be described. In the present invention, the "%" indication regarding the chemical composition means "mass%" unless otherwise specified.
Ｃ：０．０３～０．１５％
　Ｃは、必要な強度を得るために０．０３％以上の含有を必須とする。しかしながら、０．１５％を超えて含有すると、島状マルテンサイトが増加して溶接熱影響部の靱性が低下するため、上限を０．１５％とする。下限は、好ましくは０．０４５％である。また、０．１０％未満であることが好ましい。 C: 0.03 to 0.15%
C must contain 0.03% or more in order to obtain the required strength. However, if the content exceeds 0.15%, island-like martensite increases and the toughness of the weld heat-affected zone decreases, so the upper limit is set to 0.15%. The lower limit is preferably 0.045%. Further, it is preferably less than 0.10%.
Ｓｉ：０．０１～０．５０％
　Ｓｉは、母材の強度確保および脱酸などに必要な成分であり、０．０１％以上で添加する。一方、０．５０％を超えると、ＨＡＺが硬化してＨＡＺの靭性が低下するため、上限を０．５０％とする。さらに、好ましい下限は０．１０％であり、好ましい上限は０．３０％である。 Si: 0.01-0.50%
Si is a component necessary for ensuring the strength of the base material, deoxidizing, etc., and is added in an amount of 0.01% or more. On the other hand, if it exceeds 0.50%, the HAZ is hardened and the toughness of the HAZ is lowered, so the upper limit is set to 0.50%. Further, the preferred lower limit is 0.10% and the preferred upper limit is 0.30%.
Ｍｎ：１．２０～２．００％
　Ｍｎは、母材の強度を確保するために、１．２０％以上必要であり、２．００％を超えると溶接性が劣化するだけでなく鋼材コストも上昇する。したがって、Ｍｎの範囲は、１．２０～２．００％とする。下限は、好ましくは１．４０％である。上限は、好ましくは１．６０％である。 Mn: 1.20 to 2.00%
Mn is required to be 1.20% or more in order to secure the strength of the base material, and if it exceeds 2.00%, not only the weldability deteriorates but also the steel material cost increases. Therefore, the range of Mn is 1.20 to 2.00%. The lower limit is preferably 1.40%. The upper limit is preferably 1.60%.
Ｐ：０．０２０％以下
　Ｐは、不可避的に混入する不純物であり、含有量が０．０２０％を超えると、母材および溶接部の靱性を低下させるため、上限を０．０２０％とする。なお、良好な靱性を得るためには、０．０１０％以下が好ましく、０．００７％以下であることがさらに好ましい。ちなみに、下限は限定する必要はないが、極低Ｐ化処理を施すことでコストが増加してしまうため、０．００１％以上とすることが好ましい。 P: 0.020% or less P is an impurity that is inevitably mixed in, and if the content exceeds 0.020%, the toughness of the base metal and welds will decrease, so the upper limit is set to 0.020%. .. In order to obtain good toughness, 0.010% or less is preferable, and 0.007% or less is more preferable. By the way, although it is not necessary to limit the lower limit, it is preferable to set it to 0.001% or more because the cost increases by performing the ultra-low P treatment.
Ｓ：０．０００５～０．０１００％
　Ｓは、フェライト核生成に必要な複合介在物の核に所要のＭｎＳ、さらにＣａを添加する場合はＣａＳを生成させるために、０．０００５％以上は必要である。Ｓが０．０００５％未満となると、ＭｎＳ、さらにはＣａＳが十分に形成されず、ＨＡＺの靭性が低下する。一方、０．０１００％を超えると、母材の靱性を劣化させる。上限は好ましくは０．００９０％である。下限は、好ましくは０．００１０％である。 S: 0.0005 to 0.0100%
S is required to be 0.0005% or more in order to generate the required MnS in the nucleus of the composite inclusion required for ferrite nucleation, and CaS when Ca is added. When S is less than 0.0005%, MnS and further CaS are not sufficiently formed, and the toughness of HAZ is lowered. On the other hand, if it exceeds 0.0100%, the toughness of the base metal is deteriorated. The upper limit is preferably 0.0090%. The lower limit is preferably 0.0010%.
Ａｌ：０．００５～０．１００％
　Ａｌは、鋼の脱酸上０．００５％以上、好ましくは０．０１０％以上が必要である。一方、０．１００％を超えて含有すると、母材の靱性を低下させると共に溶接金属の靱性を劣化させる。上限は、好ましくは０．０８％である。 Al: 0.005 to 0.100%
Al needs to be 0.005% or more, preferably 0.010% or more in terms of deoxidation of steel. On the other hand, if it is contained in excess of 0.100%, the toughness of the base metal is lowered and the toughness of the weld metal is deteriorated. The upper limit is preferably 0.08%.
Ｔｉ：０．００４～０．０３０％
　Ｔｉは、鋼の凝固時にＴｉＮとなって析出し、溶接熱影響部（ＨＡＺ）でのオーステナイトの粗粒化抑制や、フェライト変態核となって高靱性化に寄与する。Ｔｉは、０．００４％に満たないとその効果は少なく、一方０．０３０％を超えるとＴｉＮ粒子の粗大化によって期待する効果が得られなくなる。したがって、Ｔｉの含有量は、０．００４～０．０３０％の範囲とする。下限は好ましくは０．００８％である。上限は、好ましくは０．０２０％である。 Ti: 0.004 to 0.030%
Ti precipitates as TiN during solidification of steel, and contributes to suppressing coarse-grained austenite in the weld heat-affected zone (HAZ) and becoming ferrite transformation nuclei to increase toughness. If Ti is less than 0.004%, its effect is small, while if it exceeds 0.030%, the expected effect cannot be obtained due to the coarsening of TiN particles. Therefore, the Ti content is in the range of 0.004 to 0.030%. The lower limit is preferably 0.008%. The upper limit is preferably 0.020%.
Ｂ：０．００２０～０．００５０％
　Ｂは、粒界フェライトを微細化させＨＡＺ靭性を向上させる上で重要な元素であり、フェライト変態温度以上で析出させるために少なくとも０．００２０％添加する。しかし、多量に添加すると母材靱性を劣化させるため、上限を０．００５０％とする。下限は好ましくは０．００２５％である。上限は、好ましくは０．００４０％である。 B: 0.0020-0.0050%
B is an important element for refining grain boundary ferrite and improving HAZ toughness, and is added at least 0.0020% in order to precipitate at a ferrite transformation temperature or higher. However, if a large amount is added, the toughness of the base metal deteriorates, so the upper limit is set to 0.0050%. The lower limit is preferably 0.0025%. The upper limit is preferably 0.0040%.
Ｎ：０．００３５～０．０１００％
　Ｎは、Ｔｉと結合してＴｉＮを形成し、かつＢと結合してＢＮを形成するために、０．００３５％以上で添加する。すなわち、Ｎが０．００３５％の下限を下回ると、ＢＮが形成されず十分なＨＡＺ靭性を確保できなくなる。一方、Ｎの含有量が増えると、固溶Ｎが増大しＨＡＺ靱性の低下を招くことから、０．０１００％を上限とした。下限は好ましくは０．００４０％である。上限は、好ましくは０．００９０％である。 N: 0.0035-0.0100%
N is added in an amount of 0.0035% or more in order to combine with Ti to form TiN and to combine with B to form BN. That is, when N is below the lower limit of 0.0035%, BN is not formed and sufficient HAZ toughness cannot be secured. On the other hand, when the content of N increases, the solid solution N increases and the HAZ toughness decreases, so the upper limit is 0.0100%. The lower limit is preferably 0.0040%. The upper limit is preferably 0.0090%.
　本発明の鋼板は、以上の各成分を含み、残部はＦｅおよび不可避不純物である成分組成を有する。この成分組成において、さらに、次式（１）で示されるＳＢが－０．００１０以上０．０００２以下および、次式（２）で示される温度ＴがＡｒ３点超であることが肝要である。
　ＳＢ＝［Ｂ］－０．７７×［Ｎ］＋０．２２×［Ｔｉ］　　…（１）
　Ｔ（℃）＝１２０００／（４．６３―ｌｏｇ（［Ｂ］×（［Ｎ］－［Ｔｉ］／３．４２）））－２７３　…（２）
　但し、上記式（１）および（２）における、［Ｂ］、［Ｎ］および［Ｔｉ］は各成分の含有量（質量％）を表す。 The steel sheet of the present invention contains each of the above components, and the balance has a component composition of Fe and unavoidable impurities. In this component composition, it is further important that the SB represented by the following formula (1) is −0.0010 or more and 0.0002 or less, and the temperature T represented by the following formula (2) is Ar 3 points or more. ..
SB = [B] -0.77 x [N] + 0.22 x [Ti] ... (1)
T (° C.) = 12000 / (4.63-log ([B] × ([N]-[Ti] /3.42)))-273 ... (2)
However, [B], [N] and [Ti] in the above formulas (1) and (2) represent the content (mass%) of each component.
　本発明において、Ｂ、ＮおよびＴｉは、上記した式（１）および式（２）を満足するように含有させることにより、大入熱溶接時に鋼板が受ける熱サイクル（以下、溶接熱サイクルともいう）においてもＴｉＮが固溶することなく残存し、このＴｉＮを核としてＢＮが早期に析出するようになる。ここに、図１に、上記した成分組成の鋼板に入熱１０kJ/mm相当の溶接再現熱サイクルを付与したサンプルの観察画像を示すように、溶接熱サイクルの冷却過程の初期段階においてＴｉＮの周囲にＢＮが析出していることがわかる。すなわち、ＢＮはより高温域から析出しやすくなる。かようにＢＮがＴｉＮから析出すると、ＴｉＮとＢＮの複合析出物のサイズは、ＴｉＮ単独のサイズよりも大きくなる。析出物のサイズが大きくなることによって、フェライトが核生成しやすくなる。なお、核となるＴｉＮのサイズは１５nm以上２００nm以下が通常であり、ＢＮがＴｉＮに析出するとＢＮ被覆析出物のサイズは５０nm以上６００nm以下となる。フェライトが核生成しやすいということは、粒界に多くのフェライト核が生成するということであり、粒界に多くのフェライトができる。これらのフェライトは異なるＢＮから核生成していることから方位が異なるため、フェライトの結晶方位はランダム化する。この結晶方位のランダム化により、隣接するフェライト同士が合体しなくなる。その結果、粒界フェライトが微細化し、そこから生成するフェライトサイドプレートも微細化する。従って、式（１）および式（２）を満足することでＨＡＺ靭性が向上することになる。 In the present invention, B, N and Ti are contained in the above formulas (1) and (2) so as to satisfy the above-mentioned formulas (1) and (2), so that the heat cycle received by the steel sheet during large heat input welding (hereinafter, also referred to as a welding heat cycle). ), TiN remains without solidification, and BN is deposited at an early stage with this TiN as a nucleus. Here, as shown in FIG. 1, an observation image of a sample in which a steel sheet having the above-mentioned composition composition is subjected to a welding reproduction heat cycle equivalent to 10 kJ / mm of heat input is shown. It can be seen that BN is precipitated in. That is, BN is more likely to precipitate from the high temperature region. When BN is precipitated from TiN in this way, the size of the composite precipitate of TiN and BN becomes larger than the size of TiN alone. Increasing the size of the precipitate facilitates nucleation of ferrite. The size of the core TiN is usually 15 nm or more and 200 nm or less, and when BN precipitates on TiN, the size of the BN-coated precipitate becomes 50 nm or more and 600 nm or less. The fact that ferrite is easily nucleated means that many ferrite nuclei are formed at the grain boundaries, and many ferrites are formed at the grain boundaries. Since these ferrites are nucleated from different BNs and therefore have different orientations, the crystal orientations of the ferrites are randomized. Due to this randomization of crystal orientation, adjacent ferrites do not coalesce. As a result, the grain boundary ferrite is miniaturized, and the ferrite side plate generated from the grain boundary ferrite is also miniaturized. Therefore, the HAZ toughness is improved by satisfying the formulas (1) and (2).
　すなわち、前記ＳＢの値が０．０００２を超えると、固溶Ｂが増加し該固溶Ｂによって焼き入れ性が上がり、島状マルテンサイトが形成される結果、低温での靱性を十分に確保できなくなる。また、前記ＳＢの値が－０．００１０を下回ると、ＢＮの析出が不十分になって、粒界フェライトを微細化できない。 That is, when the value of SB exceeds 0.0002, the solid solution B increases, the hardenability is improved by the solid solution B, and island-shaped martensite is formed. As a result, sufficient toughness at low temperature can be secured. It disappears. Further, when the value of SB is less than −0.0010, the precipitation of BN becomes insufficient and the grain boundary ferrite cannot be refined.
　さらに、前記式（２）は、図１に示したようにＴｉＮの周囲にＢＮが析出する際の析出温度Ｔを示しており、このＴがＡｒ３点以下になると、ＢＮを核としたフェライト生成が難しくなる結果、粒界フェライトの微細化が実現しない。 Further, the above formula (2) shows the precipitation temperature T when BN is deposited around TiN as shown in FIG. 1, and when this T becomes Ar 3 points or less, ferrite having BN as a core is shown. As a result of difficulty in formation, miniaturization of grain boundary ferrite is not realized.
　上記した成分組成を有する鋼板において、例えば、入熱量が５kJ／mm以上の大入熱溶接を行った際の、ボンド近傍の熱影響部組織は、旧γ粒界に生成する粒界フェライトの密度が２０個／mm以上となる。ここで、旧γ粒界の粒界フェライト生成密度は、溶接を模擬した熱サイクルシミュレーションの冷却途中でフェライト変態開始直後から急冷処理を行い、ＥＢＳＤ（電子線後方回折法）を用いて計測することができる。本発明において、旧γ粒界の隣接する３重点から３重点までの粒界に沿った曲線長さを旧γ粒界長さとして、その旧γ粒界上に生成した隣り合うフェライト粒の結晶方位差が１５度以上となるフェライト粒の個数を旧γ粒界上のフェライト数として、（旧γ粒界上のフェライト数）／（旧γ粒界長さ）により粒界フェライトの密度を定義する。 In a steel sheet having the above-mentioned composition, for example, when a large heat input welding with a heat input amount of 5 kJ / mm or more is performed, the heat-affected zone structure near the bond is the density of grain boundary ferrite generated at the old γ grain boundaries. Is 20 pieces / mm or more. Here, the grain boundary ferrite formation density of the old γ grain boundaries is measured by performing quenching treatment immediately after the start of ferrite transformation during cooling in a thermal cycle simulation simulating welding and using EBSD (electron backscatter diffraction method). Can be done. In the present invention, the curve length along the adjacent 3 to 3 priority grain boundaries of the old γ grain boundary is defined as the old γ grain boundary length, and the crystals of adjacent ferrite grains generated on the old γ grain boundary are used. The number of ferrite grains with an orientation difference of 15 degrees or more is defined as the number of ferrites on the old γ grain boundaries, and the density of grain boundary ferrites is defined by (number of ferrites on the old γ grain boundaries) / (former γ grain boundary length). do.
　本発明の鋼板では、上記のような大入熱溶接を施したときに、ボンド近傍の熱影響部組織の、旧γ粒界上に生成する粒界フェライトの密度が２０個／mm以上となることから、粗大なフェライトサイドプレートの生成抑制が可能となり、ＨＡＺにおいて優れた低温靱性を実現する。ここで、ボンド近傍の熱影響部組織とは、溶接金属の母材鋼板との境界からおよそ０．５mm母材である鋼板側に入った位置までの領域をいう。旧γ粒界に生成する粒界フェライトの密度は、上記した式（１）および（２）に従ってＮ、ＢおよびＴｉの添加量を規定範囲内に制御することにより、例えば入熱量が５kJ／mm以上の大入熱溶接を行った際の、粒界フェライトの密度を２０個／mm以上とすることができる。すなわち、粗大なフェライトサイドプレートの生成が抑制されて、熱影響部において優れた靱性を得ることができる。 In the steel sheet of the present invention, the density of grain boundary ferrites formed on the old γ grain boundaries in the heat-affected zone structure near the bond becomes 20 grains / mm or more when the above-mentioned large heat-immersive welding is performed. Therefore, it is possible to suppress the formation of coarse ferrite side plates and realize excellent low temperature toughness in HAZ. Here, the heat-affected zone structure in the vicinity of the bond refers to a region from the boundary of the weld metal with the base steel plate to a position within about 0.5 mm on the steel plate side of the base material. The density of grain boundary ferrite generated at the old γ grain boundaries is determined by controlling the addition amounts of N, B and Ti within the specified range according to the above formulas (1) and (2), for example, the heat input amount is 5 kJ / mm. The density of grain boundary ferrites when the above-mentioned large heat input welding is performed can be 20 pieces / mm or more. That is, the formation of coarse ferrite side plates is suppressed, and excellent toughness can be obtained in the heat-affected zone.
　本発明において、上記成分組成を満足することにより低温靭性の向上を達成できるが、一方で、母材および継手の強度の確保が難しくなる。そこで、強度の確保を目的として、本発明に係る鋼板の金属組織は、該組織に占める加工フェライトの割合を体積分率で５０％以上とすることが肝要である。ここで、加工フェライトとは、Ｘ線回析（XRD）により求められる、転位密度ρの値が、１．０×１０１４ｍ－２以上のフェライトを指す。
　すなわち、加工フェライトは、高密度の転位が導入されており、転位同士が相互作用を起こし互いの運動を妨げあうことで強度が上昇する。そして、この加工フェライトの体積分率を５０％以上とすることによって、強度が上昇することになる。 In the present invention, improvement of low temperature toughness can be achieved by satisfying the above component composition, but on the other hand, it becomes difficult to secure the strength of the base material and the joint. Therefore, for the purpose of ensuring strength, it is important that the metal structure of the steel sheet according to the present invention has a volume fraction of 50% or more of processed ferrite in the structure. Here, the processed ferrite refers to a ferrite having a dislocation density ρ of 1.0 × 10 14 m- 2 or more, which is determined by X-ray diffraction (XRD).
That is, in the processed ferrite, high-density dislocations are introduced, and the dislocations interact with each other to hinder each other's movements, thereby increasing the strength. Then, by setting the volume fraction of the processed ferrite to 50% or more, the strength is increased.
　なお、金属組織に占める加工フェライトの体積分率は、好ましくは６０％以上である。一方、加工フェライト量の上限は特に限定する必要はなく、１００％であってもよいが、圧延機の能力の観点からは、９０％以下であることが好ましい。その際の残部組織は、パーライト、ベイナイトおよびマルテンサイトのうちの１種類以上の硬質相であることが好ましい。 The volume fraction of processed ferrite in the metal structure is preferably 60% or more. On the other hand, the upper limit of the amount of processed ferrite is not particularly limited and may be 100%, but from the viewpoint of the capacity of the rolling mill, it is preferably 90% or less. The remaining structure at that time is preferably one or more hard phases of pearlite, bainite and martensite.
　本発明の他の実施形態においては、さらに特性を向上させるため、上記成分組成に加えて、Ｃｕ、Ｎｉ、Ｎｂ、Ｖ、Ｃｒ、Ｍｏ、Ｃａ、ＭｇおよびＲＥＭからなる群より選択される１種または２種以上を任意に含有することが可能である。 In another embodiment of the present invention, in order to further improve the characteristics, in addition to the above component composition, one selected from the group consisting of Cu, Ni, Nb, V, Cr, Mo, Ca, Mg and REM. Alternatively, two or more kinds can be arbitrarily contained.
Ｃｕ：０．０１～０．５０％
　Ｃｕは、鋼の焼き入れ性を高める元素であり、圧延後の母材の強度向上に加え、靱性、高温強度、耐候性などの機能向上に寄与する。これらの効果は、０．０１％以上の含有によって発揮される。一方、過度の含有は母材の靱性や溶接性をかえって劣化させる。そのため、Ｃｕ含有量は０．０１～０．５０％とすることが好ましい。 Cu: 0.01-0.50%
Cu is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal. Therefore, the Cu content is preferably 0.01 to 0.50%.
Ｎｉ：０．０１～１．５０％
　Ｎｉは、鋼の焼き入れ性を高める元素であり、圧延後の母材の強度向上に加え、靱性、高温強度、耐候性などの機能向上に寄与する。これらの効果は、０．０１％以上の含有によって発揮される。一方で、過度の含有は母材の靱性や溶接性をかえって劣化させることに加え、合金のコスト増加を招く。そのため、Ｎｉ含有量は０．０１～１．５０％とすることが好ましい。 Ni: 0.01-1.50%
Ni is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal, and also increases the cost of the alloy. Therefore, the Ni content is preferably 0.01 to 1.50%.
Ｎｂ：０．００５～０．０４０％
　Ｎｂは、母材の強度、靭性および継手の強度を確保するのに有効な元素である。その効果は０．００５％以上の含有により発揮される。一方、０．０４０％を超えて含有すると、溶接熱影響部に島状マルテンサイトを形成することにより靭性が劣化する。そのため、Ｎｂを含有する場合、Ｎｂ含有量を０．００５～０．０４０％とすることが好ましい。 Nb: 0.005 to 0.040%
Nb is an element effective for ensuring the strength, toughness and joint strength of the base metal. The effect is exhibited when the content is 0.005% or more. On the other hand, if it is contained in excess of 0.040%, the toughness deteriorates due to the formation of island-shaped martensite in the weld heat affected zone. Therefore, when Nb is contained, the Nb content is preferably 0.005 to 0.040%.
Ｖ：０．００５～０．１００％
　Ｖは、母材の強度・靭性の向上およびＶＮとしてフェライト生成核として働く。この効果はＶを０．００５％以上含有させることにより発揮される。一方、Ｖは０．１００％を超えて含有すると、かえって母材の靱性が低下する。このため、Ｖを含有させる場合には、Ｖ含有量を０．００５～０．１００％とすることが好ましい。 V: 0.005 to 0.100%
V acts as a ferrite nucleation nucleus as a VN and improves the strength and toughness of the base metal. This effect is exhibited by containing 0.005% or more of V. On the other hand, if V is contained in an amount of more than 0.100%, the toughness of the base metal is rather lowered. Therefore, when V is contained, the V content is preferably 0.005 to 0.100%.
Ｃｒ：０．０１～０．５０％
　Ｃｒは、Ｃｕと同様に、鋼の焼き入れ性を高める元素であり、圧延後の母材の強度向上に加え、靱性、高温強度、耐候性などの機能向上に寄与する。これらの効果は、０．０１％以上の含有によって発揮される。一方、過度の含有は母材の靱性や溶接性をかえって劣化させる。そのため、Ｃｒ含有量は０．０１～０．５０％とすることが好ましい。 Cr: 0.01-0.50%
Like Cu, Cr is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal. Therefore, the Cr content is preferably 0.01 to 0.50%.
Ｍｏ：０．０１～０．５０％
　Ｍｏは、ＣｕやＣｒと同様に、鋼の焼き入れ性を高める元素であり、圧延後の母材の強度向上に加え、靱性、高温強度、耐候性などの機能向上に寄与する。これらの効果は、０．０１％以上の含有によって発揮される。一方、過度の含有は母材の靱性や溶接性をかえって劣化させる。そのため、Ｍｏ含有量は０．０１～０．５０％とすることが好ましい。 Mo: 0.01-0.50%
Like Cu and Cr, Mo is an element that enhances the hardenability of steel, and contributes to the improvement of functions such as toughness, high temperature strength, and weather resistance in addition to the improvement of the strength of the base metal after rolling. These effects are exhibited by the content of 0.01% or more. On the other hand, excessive content deteriorates the toughness and weldability of the base metal. Therefore, the Mo content is preferably 0.01 to 0.50%.
Ｃａ：０．０００５～０．００３０％
　Ｃａは、Ｓの固定による母材の靭性向上に有用な元素であるが、含有量が０．００３０％を超えるとその効果は飽和するので、Ｃａは０．００３０％以下で含有させるものとする。一方、含有量が０．０００５％未満であると、Ｓの固定が不十分となる。そのため、Ｃａの含有量は、０．０００５％以上０．００３０％以下とすることが好ましい。 Ca: 0.0005 to 0.0030%
Ca is an element useful for improving the toughness of the base metal by fixing S, but the effect is saturated when the content exceeds 0.0030%, so Ca should be contained at 0.0030% or less. .. On the other hand, if the content is less than 0.0005%, the fixation of S becomes insufficient. Therefore, the Ca content is preferably 0.0005% or more and 0.0030% or less.
Ｍｇ：０．０００２～０．００５０％
ＲＥＭ：０．００１０～０．１０００％
　ＭｇおよびＲＥＭは、いずれも溶鋼中で強い脱酸力を有し、微細酸化物形成を補助する働きがあることから、必要に応じて添加する。脱酸効果を示す添加量はそれぞれ、Ｍｇ：０．０００２％以上、ＲＥＭ：０．００１０％以上であるが、多量に添加すると、粗大な介在物ができて母材特性を損ねることから、添加の上限をそれぞれＭｇ：０．００５０％およびＲＥＭ：０．１０００％とすることが好ましい。 Mg: 0.0002 to 0.0050%
REM: 0.0010 to 0.1000%
Both Mg and REM have a strong deoxidizing power in molten steel and have a function of assisting the formation of fine oxides, and therefore are added as necessary. The addition amounts showing the deoxidizing effect are Mg: 0.0002% or more and REM: 0.0010% or more, respectively. However, if a large amount is added, coarse inclusions are formed and the properties of the base material are impaired. It is preferable that the upper limits of Mg: 0.0050% and REM: 0.1000%, respectively.
　本発明の製造方法は、以上の成分組成を有する鋼素材を１０５０℃以上１２００℃以下の温度に加熱後、９００℃以下の温度まで７℃／ｓ以下の冷却速度で冷却後、８５０℃以下のフェライト－オーステナイトの二相温度域における累積圧下率が６０％以上および、仕上温度が６５０℃以上である、熱間圧延を施す。次に、本発明における製造条件の限定理由について説明する。 In the production method of the present invention, a steel material having the above composition is heated to a temperature of 1050 ° C. or higher and 1200 ° C. or lower, cooled to a temperature of 900 ° C. or lower at a cooling rate of 7 ° C./s or lower, and then cooled to 850 ° C. or lower. Hot rolling is performed in which the cumulative reduction rate of ferrite-austenite in the two-phase temperature range is 60% or more and the finishing temperature is 650 ° C. or more. Next, the reason for limiting the production conditions in the present invention will be described.
［鋼素材加熱温度および９００℃以下の温度までの冷却速度］
　まず、鋼素材、例えばスラブの加熱温度は、１０５０℃以上１２００℃以下であることが必要である。この理由は、１０５０℃未満の加熱では、凝固中に生成した靱性に悪影響を及ぼす粗大な介在物が溶けずに残る可能性があるためである。一方、高温で加熱すると、後述する冷却速度を制御して造りこんだ析出物を再溶解させてしまう可能性がある。これを踏まえると、相変態を完了させる意味での加熱温度としては１２００℃以下で十分である。なお、そのときに生じると考えられる結晶粒の粗大化も、上記したＴｉＮのピンニング効果により、あらかじめ防ぐことができる。以上より、加熱温度を１０５０℃以上１２００℃以下に限定した。 [Steel material heating temperature and cooling rate up to 900 ° C or less]
First, the heating temperature of the steel material, for example, the slab, needs to be 1050 ° C. or higher and 1200 ° C. or lower. The reason for this is that heating below 1050 ° C. may leave coarse inclusions that adversely affect the toughness produced during solidification undissolved. On the other hand, when heated at a high temperature, there is a possibility that the precipitates formed by controlling the cooling rate described later may be redissolved. Based on this, 1200 ° C. or lower is sufficient as the heating temperature in the sense of completing the phase transformation. It should be noted that the coarsening of crystal grains that is considered to occur at that time can also be prevented in advance by the pinning effect of TiN described above. From the above, the heating temperature was limited to 1050 ° C. or higher and 1200 ° C. or lower.
　次いで、９００℃以下の温度まで７℃／ｓ以下の冷却速度で冷却する必要がある。この理由は、７℃／ｓを超える冷却速度では、ＢがＢＮとして析出せずに、粒界に固溶Ｂとして残存することで、粒界フェライトの生成が抑制され、母材の低温靭性を十分に確保することが困難となるためである。なお、冷却速度は、製造能率の観点から、１℃／ｓ以上とすることが好ましい。 Next, it is necessary to cool to a temperature of 900 ° C or lower at a cooling rate of 7 ° C / s or lower. The reason for this is that at a cooling rate exceeding 7 ° C./s, B does not precipitate as BN but remains as a solid solution B at the grain boundaries, so that the formation of grain boundary ferrite is suppressed and the low temperature toughness of the base metal is improved. This is because it is difficult to secure a sufficient amount. The cooling rate is preferably 1 ° C./s or higher from the viewpoint of production efficiency.
［熱間圧延条件］
　８５０℃以下のフェライト＋オーステナイトの二相温度域において累積圧下率６０％以上の熱間圧延を行う必要がある。その理由として、二相温度域における圧下量の増加は、圧延中のフェライトの加工による転位強化に伴う強度向上と加工によるサブグレインの形成を通じた細粒化の効果による靱性の向上の効果があるからである。 [Hot rolling conditions]
It is necessary to perform hot rolling with a cumulative rolling reduction of 60% or more in a two-phase temperature range of ferrite + austenite at 850 ° C or lower. The reason is that the increase in the amount of reduction in the two-phase temperature range has the effect of improving the strength due to dislocation strengthening due to the processing of ferrite during rolling and the effect of improving toughness due to the effect of fine granulation through the formation of subgrains during processing. Because.
　すなわち、二相温度域における累積圧下率を６０％以上に増加することにより、二相域の温度領域のフェライトに転位が加わる結果、強度の向上を図ることができる。特に、累積圧下率を６０％以上に高めることにより、加工フェライト分率を５０％以上確保することが可能となる。さらに、フェライト＋オーステナイトの二相温度域での累積圧下率が６０％以上となることにより、フェライトの圧延集合組織が発達し、低温靱性の向上に寄与することになる。
　以上のことから、８５０℃以下のフェライト＋オーステナイトの二相温度域における累積圧下率を６０％以上に限定した。なお、累積圧下率は、圧延機能力の観点から、９０％以下とすることが好ましい。 That is, by increasing the cumulative reduction rate in the two-phase temperature range to 60% or more, dislocations are added to the ferrite in the two-phase temperature range, and as a result, the strength can be improved. In particular, by increasing the cumulative reduction rate to 60% or more, it is possible to secure a processed ferrite fraction of 50% or more. Further, when the cumulative reduction rate of ferrite + austenite in the two-phase temperature range is 60% or more, the rolled texture of ferrite develops, which contributes to the improvement of low temperature toughness.
From the above, the cumulative reduction rate of ferrite + austenite in the two-phase temperature range of 850 ° C. or lower was limited to 60% or more. The cumulative rolling reduction ratio is preferably 90% or less from the viewpoint of rolling functional force.
　さらに、熱間圧延における仕上温度を、６５０℃以上とする。なぜなら、６５０℃未満で仕上圧延を行うと、相変態により生成したフェライトに必要以上に歪を与えることになり、靱性が低下してしまうためである。 Further, the finishing temperature in hot rolling is set to 650 ° C. or higher. This is because if the finish rolling is performed at a temperature lower than 650 ° C., the ferrite produced by the phase transformation is distorted more than necessary, and the toughness is lowered.
　さらに、上記の熱間圧延を施した後、６５０℃以上の温度から５℃／ｓ以上の冷却速度で６００℃以下３００℃以上の温度域まで冷却することが、母材の強度を高める上で好ましい。
［熱間圧延後冷却条件］
　また、前記鋼素材は６５０℃以上で熱間圧延を完了させた後、６５０℃以上の温度から５℃／ｓ以上の冷却速度で６００℃以下３００℃以上の温度域まで冷却することが好ましい。すなわち、６５０℃以上から冷却する理由として、６５０℃未満にて冷却を開始すると、焼き入れ性が不十分となり、所要の強度が得られない可能性があるためである。また、冷却速度が５℃／ｓ未満では均一なミクロ組織を有する鋼を得ることが難しくなる。さらに、６００℃以下３００℃以上の温度域まで冷却することが好ましい。なぜなら、６００℃を超える温度での冷却停止では、焼き入れ性の観点から、十分な強度確保が困難となるためである。また、３００℃未満の温度での冷却停止は、鋼材特性に大きな変化を与えないことから、操業上の負荷のみが大きくなるためである。上記の理由により、鋼片は６５０℃以上で熱間延性を完了させた後、６５０℃以上の温度から５℃／ｓ以上の冷却速度で６００℃以下３００℃以上まで冷却することが好ましい。なお、冷却速度は、母材の靭性確保の観点から、５０℃／ｓ以上とすることが好ましい。 Further, after performing the above hot rolling, cooling from a temperature of 650 ° C. or higher to a temperature range of 600 ° C. or lower and 300 ° C. or higher at a cooling rate of 5 ° C./s or higher is used to increase the strength of the base metal. preferable.
[Cooling conditions after hot rolling]
Further, it is preferable that the steel material is cooled from a temperature of 650 ° C. or higher to a cooling rate of 5 ° C./s or higher to a temperature range of 600 ° C. or lower and 300 ° C. or higher after hot rolling is completed at 650 ° C. or higher. That is, the reason for cooling from 650 ° C. or higher is that if cooling is started at a temperature lower than 650 ° C., the hardenability becomes insufficient and the required strength may not be obtained. Further, if the cooling rate is less than 5 ° C./s, it becomes difficult to obtain a steel having a uniform microstructure. Further, it is preferable to cool to a temperature range of 600 ° C. or lower and 300 ° C. or higher. This is because it is difficult to secure sufficient strength from the viewpoint of hardenability when cooling is stopped at a temperature exceeding 600 ° C. In addition, stopping cooling at a temperature of less than 300 ° C. does not significantly change the characteristics of the steel material, so that only the operational load increases. For the above reasons, it is preferable that the steel pieces are cooled from a temperature of 650 ° C. or higher to a cooling rate of 5 ° C./s or higher to 600 ° C. or lower and 300 ° C. or higher after completing hot ductility at 650 ° C. or higher. The cooling rate is preferably 50 ° C./s or higher from the viewpoint of ensuring the toughness of the base material.
　上記した鋼素材はスラブを用いることができるが、これら鋼素材を鋳造によって製造する場合、その鋳造条件に関して、以下の条件を満たすことが好ましい。すなわち、スラブ鋳造時の冷却速度を、０．３ｍ/min以上かつ１．０ｍ/min以下にする。鋳造速度が０．３ｍ/min未満になると、母材（鋼板）のＴｉＮのサイズが大きくなってしまう。ＴｉＮサイズが大きくなると、母材（鋼板）のＴｉＮ密度が低下してＢＮ複合析出物の量が減少する可能性がある。その結果、十分にフェライトが微細化できず、ＨＡＺ靭性が劣化する、おそれがある。なお、核となるＴｉＮのサイズは１５nm以上２００nm以下である。一方、鋳造速度が１．０ｍ/minを超えると、ＴｉＮの密度は増加するが、ＴｉＮのサイズが小さくなってしまい、溶接時に固溶してしまう可能性がある。その結果、オーステナイト粒径が粗大化して、ＨＡＺ靭性が劣化する、おそれがある。 Although slabs can be used for the above-mentioned steel materials, when these steel materials are manufactured by casting, it is preferable to satisfy the following conditions regarding the casting conditions. That is, the cooling rate during slab casting is set to 0.3 m / min or more and 1.0 m / min or less. If the casting speed is less than 0.3 m / min, the size of TiN of the base metal (steel plate) becomes large. As the TiN size increases, the TiN density of the base material (steel plate) may decrease and the amount of BN composite precipitates may decrease. As a result, the ferrite cannot be sufficiently miniaturized, and the HAZ toughness may deteriorate. The size of the core TiN is 15 nm or more and 200 nm or less. On the other hand, when the casting speed exceeds 1.0 m / min, the density of TiN increases, but the size of TiN becomes small, and there is a possibility that the TiN will dissolve in solid solution during welding. As a result, the austenite particle size may become coarse and the HAZ toughness may deteriorate.
　かくして製造される鋼板は、上記した成分組成に加えて、加工フェライトの体積分率が５０％以上の組織を有することになる。好ましくは、フェライトからなる軟質相を含む主相とし、残部がパーライト、ベイナイトおよびマルテンサイトのうちの１種類以上の硬質相からなる組織である。ここで、主相がフェライトとは、フェライトが体積分率で６０％以上であることを意味する。すなわち、フェライトは１００％であってもよいが、圧延性の観点からは９０％以下であることが好ましい。その際の残部は特に限定する必要はなく、例えば上記した通りである。ここで重要なのは、フェライトのうちの加工フェライトが組織に占める比率であり、該比率を体積分率で５０％以上とすることである。従って、加工フェライト以外のフェライト、すなわち転位密度ρの値が１．０×１０１４ｍ－２未満のフェライトが含まれていてもよい。 The steel sheet thus produced has a structure having a volume fraction of processed ferrite of 50% or more in addition to the above-mentioned component composition. Preferably, the main phase contains a soft phase made of ferrite, and the balance is a structure made of one or more hard phases of pearlite, bainite and martensite. Here, when the main phase is ferrite, it means that ferrite has a volume fraction of 60% or more. That is, the ferrite may be 100%, but it is preferably 90% or less from the viewpoint of rollability. The remaining portion at that time does not need to be particularly limited, and is as described above, for example. What is important here is the ratio of processed ferrite to the structure among the ferrites, and the ratio should be 50% or more in terms of volume fraction. Therefore, ferrites other than processed ferrites, that is, ferrites having a dislocation density ρ of less than 1.0 × 10 14 m- 2 may be contained.
　また、上記した用途の鋼板としては、低温靭性が高いことに加えて、とくに降伏応力が３２５MPa以上であることが好ましい。また、母材の－７０℃でのシャルピー衝撃吸収エネルギーは２００Ｊ以上であることが望ましい。さらに、大入熱溶接施工された継手の－７０℃でのシャルピー衝撃吸収エネルギーは８０Ｊ以上であることが望ましい。 Further, as the steel sheet for the above-mentioned applications, in addition to high low temperature toughness, it is particularly preferable that the yield stress is 325 MPa or more. Further, it is desirable that the Charpy impact absorption energy of the base material at −70 ° C. is 200 J or more. Further, it is desirable that the Charpy impact absorption energy at −70 ° C. of the joint subjected to the large heat input welding is 80 J or more.
　次に、本発明を実施例に基づいて具体的に説明する。
　表１に示す成分組成に調整した鋼スラブ（鋼素材）に対して、表２に示す種々の条件に従って、スラブ加熱後に冷却し、次いで熱間圧延そして冷却処理を施して板厚２０mmの厚鋼板とした。 Next, the present invention will be specifically described based on examples.
A steel slab (steel material) adjusted to the composition shown in Table 1 is cooled after heating the slab according to various conditions shown in Table 2, and then hot-rolled and cooled to obtain a thick steel sheet having a thickness of 20 mm. And said.
　かくして得られた各厚鋼板からJIS Ｚ2241に準拠した引張試験片を採取し、JIS Ｚ2241に準拠した引張試験を行って降伏応力を測定した。また、各厚鋼板からJIS Ｚ2242に準拠した試験片を採取し、各試験片にＶ開先を加工し、JIS Ｚ2242に準拠したシャルピー衝撃試験を行って－７０℃でのシャルピー衝撃吸収エネルギーを測定した。さらに、得られた各厚鋼板から、溶接継手作製用試験片を採取し、試験片にＶ開先を加工し、サブマージアーク溶接（溶接入熱量：１０２ｋＪ／ｃｍ）により、溶接継手を作製した。これら溶接継手から切欠き位置をボンド部とするJIS４号衝撃試験片を採取し、シャルピー衝撃試験を実施し、－７０℃でのシャルピー衝撃吸収エネルギーを測定した。 Tensile test pieces conforming to JIS Z2241 were collected from each of the thick steel sheets thus obtained, and a tensile test conforming to JIS Z2241 was performed to measure the yield stress. In addition, JIS Z2242 compliant test pieces are collected from each thick steel plate, V groove is processed on each test piece, and a Charpy impact test compliant with JIS Z2242 is performed to measure Charpy impact absorption energy at -70 ° C. bottom. Further, a test piece for welding a welded joint was collected from each of the obtained thick steel plates, a V groove was machined on the test piece, and a welded joint was manufactured by submerged arc welding (welding heat input: 102 kJ / cm). A JIS No. 4 impact test piece having a notch position as a bond portion was collected from these welded joints, a Charpy impact test was carried out, and the Charpy impact absorption energy at −70 ° C. was measured.
　また、各厚鋼板の板厚1/4位置から、圧延方向に垂直な方向の断面を切り出し、金属組織観察用のサンプルとした。組織サンプルは鏡面まで研磨し、ナイタール（３％硝酸－エタノール溶液）で腐食して組織を現出して、１００倍の光学顕微鏡により組織を観察した。５視野で確認して、フェライトの分率を測定した。その際、フェライトに対してＸ線回析により転位密度ρを測定し、転位密度ρの値が１．０×１０１４ｍ－２以上となる加工フェライトの体積分率を測定した。この測定結果を、加工フェライトの体積分率とした。
　以上の各測定結果を、表３に示す。 In addition, a cross section in the direction perpendicular to the rolling direction was cut out from the plate thickness 1/4 position of each thick steel plate to prepare a sample for observing the metallographic structure. The tissue sample was polished to a mirror surface, corroded with nital (3% nitric acid-ethanol solution) to reveal the tissue, and the tissue was observed with a 100x optical microscope. The fraction of ferrite was measured after confirming with 5 fields of view. At that time, the dislocation density ρ was measured for the ferrite by X-ray diffraction, and the volume fraction of the processed ferrite having a dislocation density ρ value of 1.0 × 10 14 m- 2 or more was measured. This measurement result was used as the volume fraction of processed ferrite.
The results of each of the above measurements are shown in Table 3.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000003

Claims (4)

  1. 　質量％で、
    　Ｃ：０．０３～０．１５％、
    　Ｓｉ：０．０１～０．５０％、
    　Ｍｎ：１．２０～２．００％、
    　Ｐ：０．０２０％以下、
    　Ｓ：０．０００５～０．０１００％、
    　Ａｌ：０．００５～０．１００％、
    　Ｔｉ：０．００４～０．０３０％、
    　Ｂ：０．００２０～０．００５０％および
    　Ｎ：０．００３５～０．０１００％
    を、次式（１）で示されるＳＢが－０．００１０以上０．０００２以下および、次式（２）で示される温度ＴがＡｒ３点超となる範囲にて含有し、残部はＦｅおよび不可避不純物である成分組成を有し、加工フェライトの体積分率が５０％以上である金属組織を有する、鋼板。
    　ＳＢ＝［Ｂ］－０．７７×［Ｎ］＋０．２２×［Ｔｉ］　　…（１）
    　Ｔ（℃）＝１２０００／（４．６３―ｌｏｇ（［Ｂ］×（［Ｎ］－［Ｔｉ］／３．４２）））－２７３　…（２）
    　但し、前記式（１）および（２）における、［Ｂ］、［Ｎ］および［Ｔｉ］は各成分の含有量（質量％）を表す。     By mass%
    C: 0.03 to 0.15%,
    Si: 0.01-0.50%,
    Mn: 1.20 to 2.00%,
    P: 0.020% or less,
    S: 0.0005 to 0.0100%,
    Al: 0.005 to 0.100%,
    Ti: 0.004 to 0.030%,
    B: 0.0020 to 0.0050% and N: 0.0035 to 0.0100%
    Is contained in a range where the SB represented by the following formula (1) is −0.0010 or more and 0.0002 or less and the temperature T represented by the following formula (2) is more than 3 points of Ar, and the balance is Fe and A steel sheet having a component composition that is an unavoidable impurity and having a metal structure having a volume fraction of processed ferrite of 50% or more.
    SB = [B] -0.77 x [N] + 0.22 x [Ti] ... (1)
    T (° C.) = 12000 / (4.63-log ([B] × ([N]-[Ti] /3.42)))-273 ... (2)
    However, [B], [N] and [Ti] in the formulas (1) and (2) represent the content (mass%) of each component.
  2. 　前記成分組成は、さらに、質量％で、
    　Ｃｕ：０．０１～０．５０％、
    　Ｎｉ：０．０１～１．５０％、
    　Ｎｂ：０．００５～０．０４０％、
    　Ｖ：０．００５～０．１００％、
    　Ｃｒ：０．０１～０．５０％、
    　Ｍｏ：０．０１～０．５０％、
    　Ｃａ：０．０００５～０．００３０％、
    　Ｍｇ：０．０００２～０．００５０％および
    　ＲＥＭ：０．００１０～０．１０００％
    の１種または２種以上を含有する請求項１に記載の鋼板。     The composition of the components is further increased by mass%.
    Cu: 0.01-0.50%,
    Ni: 0.01-1.50%,
    Nb: 0.005 to 0.040%,
    V: 0.005 to 0.100%,
    Cr: 0.01-0.50%,
    Mo: 0.01-0.50%,
    Ca: 0.0005 to 0.0030%,
    Mg: 0.0002 to 0.0050% and REM: 0.0010 to 0.1000%
    The steel sheet according to claim 1, which contains one or more of the above.
  3. 　請求項１または請求項２に記載の成分組成を有する鋼素材を１０５０℃以上１２００℃以下の温度に加熱後、９００℃以下の温度まで７℃／ｓ以下の冷却速度で冷却後、８５０℃以下のフェライト－オーステナイトの二相温度域における累積圧下率が６０％以上および、仕上温度が６５０℃以上である、熱間圧延を施す、鋼板の製造方法。 The steel material having the component composition according to claim 1 or 2 is heated to a temperature of 1050 ° C. or higher and 1200 ° C. or lower, cooled to a temperature of 900 ° C. or lower at a cooling rate of 7 ° C./s or lower, and then 850 ° C. or lower. A method for producing a steel sheet by hot rolling, wherein the cumulative reduction rate of ferrite-austenite in the two-phase temperature range is 60% or more and the finishing temperature is 650 ° C. or more.
  4. 　前記熱間圧延を施した後、６５０℃以上の温度から５℃／ｓ以上の冷却速度で６００℃以下３００℃以上の温度域まで冷却する、請求項３に記載の鋼板の製造方法。
          The method for producing a steel sheet according to claim 3, wherein after the hot rolling is performed, the steel sheet is cooled from a temperature of 650 ° C. or higher to a temperature range of 600 ° C. or lower and 300 ° C. or higher at a cooling rate of 5 ° C./s or higher.
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