Steel plate and its manufacturing method

Abstract

Provided are extra-thick steel flange materials with excellent fatigue properties and low-temperature impact toughness, and methods for producing the same. [Solution] The microstructure has a surface layer region from the surface to the 1 / 8 point in the thickness direction, which contains 80% or more polygonal ferrite and the remainder bainite, in terms of area %, and a central region from the 1 / 8 point to the 1 / 2 point in the thickness direction, which is the remaining region, which is the central region, which contains 80% or more of one or more of bainite and martensite, and the bainite or martensite in the central region has a high-angle grain boundary of 15° or more, the average packet size is 25 μm or less, and the difference in porosity between the region from the surface to the 3 / 8 point in the thickness direction and the region from the 3 / 8 point to the 5 / 8 point in the thickness direction is 0.100 mm 3 / g or less, and the Charpy impact absorption energy value at -50°C is 50J or more.

Description

[Technical Field]

[0001] The present invention relates to a steel plate and a manufacturing method thereof, and more particularly to an extra-heavy steel material for flanges having excellent fatigue properties and low-temperature impact toughness, and a manufacturing method thereof. [Background technology]

[0002] Wind turbines, which are gaining attention as an environmentally friendly means of generating electricity, include components such as tower flanges, bearings, and main shafts. Among these, tower flanges are necessary for connecting towers. Three to four flanges are typically used per tower, and high durability is required for installations at sea or in extremely cold regions. In particular, wind turbines are becoming larger in size to meet demands for larger-capacity energy production and higher efficiency. Accordingly, the steel materials used in these turbines are continuously required to have high fatigue properties and low-temperature impact toughness. As the thickness of a material increases, the total deformation decreases, which leads to an increase in the microstructure, and the material tends to deteriorate due to defects in the material, such as inclusions and segregations.

[0003] Therefore, in order to improve the internal and external soundness of steel, there is a tendency to reduce the concentration of impurities such as non-metallic inclusions and segregations, and to control cracks and voids on the surface and inside the material as much as possible.

[0004] In particular, in the case of extremely thick materials with a thickness exceeding 200 mm, the deformation amount at the center of the material is not large, so if the unsolidified shrinkage cavities that occur during continuous casting or casting are not sufficiently compressed during the forging process, they remain in the form of residual voids at the center of the flange.

[0005] Such residual voids can act as crack initiation points when a structure is subjected to stress in the thickness direction, eventually causing damage to the entire equipment in the form of lamellar tearing. Therefore, a process is required to sufficiently compress the central void so that no residual voids remain.

[0006] Related Patent Document 1 (Patent Document 1) describes a technology for applying reduction in thickness during rough rolling of thick plates. Specifically, the technology utilizes a technique for determining the critical reduction for each thickness at which plate jamming occurs from the reduction for each pass set to approximate the design tolerances (load and torque) of the rolling mill, a technique for allocating the reduction by adjusting the index of the thickness ratio for each pass to ensure the target thickness of the rough rolling mill, and a technique for correcting the reduction based on the critical reduction for each thickness to prevent plate jamming. This technology provides a manufacturing method that can approve an average reduction of approximately 27.5% in the final three passes of rough rolling with a standard thickness of 80 mm. However, with this rolling method, when the average reduction for the entire product thickness is measured, applying high deformation to the center where residual porosity exists is technically difficult for extremely thick materials with a maximum thickness of 200 mm or more.

[0007] One of the other methods for manufacturing extra-thick products is to utilize a forging machine, which has a higher effective deformation per pass than a rolling mill. Patent Document 2 discloses a method for producing a slab containing, by mass%, C: 0.08 to 0.20%, Si: 0.40% or less, Mn: 0.5 to 5.0%, P: 0.010% or less, S: 0.0050% or less, Cr: 3.0% or less, Ni: 0.1 to 5.0%, Al: 0.010 to 0.080%, N: 0.0070% or less, and O: 0.0025% or less, which satisfies the relationships of formulas (1) and (2), with the remainder being Fe and unavoidable impurities, by heat treatment with a cumulative reduction of 25% or more. The document clearly states that by performing cold forging, heating from the Ac3 point to 1200°C and below, and hot rolling with a cumulative reduction of 40% or more, quenching from a temperature above the Ar3 point to 350°C or below, or a lower temperature below the Ar3 point, and tempering at a temperature of 450-700°C, it is possible to produce a thick, high-toughness, high-strength material with a plate thickness of 100 mm or more, a yield strength of 620 MPa or above, and an absorbed energy of 70 J or more when evaluated for low-temperature impact toughness at -40°C. However, if the cumulative reduction in this manufacturing method is too high, localized deformation concentration can cause surface defects. In particular, if surface or subsurface defects exist in the cast slab before forging, the defects can propagate during the forging process, further deteriorating the surface quality of the product after rolling. Furthermore, if the forging reduction per pass is insufficient, it is difficult to fully close the voids remaining in the center even if the cumulative reduction is high. In addition, the rolling process is not suitable for controlling the voids and structure in the center of an extremely thick material because the effective deformation of the center is small relative to the surface deformation.

[0008] Meanwhile, Patent Document 3 discloses that a thick, high-strength steel plate with a yield strength of 620 MPa or greater and a thickness of 100 mm or greater can be manufactured by heating a material with a predetermined alloy composition to 1200-1350°C, hot forging it with a cumulative reduction of 25% or greater, heating it to between the Ac3 point and 1200°C, hot rolling it with a cumulative reduction of 40% or greater, reheating it to between the Ac3 point and 1050°C, quenching it from a temperature above the Ac3 point to the lower of either 350°C or less or the Ar3 point or less, and tempering it at a temperature of 450-700°C. However, the ultra-high-strength steel plate described above not only has a high carbon equivalent (Ceq) and hardenability index (DI), making it vulnerable to surface cracking during casting, but also has low fatigue propagation resistance due to localized low-temperature transformation structures such as martensite and bainite formed in the surface layer during the cooling process, ultimately making it difficult to ensure adequate fatigue quality.

[0009] Therefore, a forging method has been proposed to close the void in the center and improve the internal integrity of the final product, but no practical method has been presented to ensure both the appropriate material quality and excellent surface quality of the flange steel. [Prior art documents] [Patent documents]

[0010] [Patent Document 1] Korean Patent Publication No. 10-2012-0075246 [Patent Document 2] Korean Patent Publication No. 10-2017-0095307 [Patent Document 3] Korean Patent Publication No. 10-2017-0095307 Summary of the Invention [Problem to be solved by the invention]

[0011] According to one embodiment of the present invention, a steel sheet and a method for manufacturing the same are provided.

[0012] According to one embodiment of the present invention, an extra-thick steel material for flanges having excellent fatigue properties and low-temperature impact toughness and a method for manufacturing the same are provided.

[0013] The object of the present invention is not limited to the above-mentioned content, and a person skilled in the art will have no difficulty in understanding further object of the present invention from the entire content of this specification. [Means for solving the problem]

[0014] According to one embodiment of the present invention, the alloy comprises, in weight percent, carbon (C): 0.05 to 0.20%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.00 to 2.00%, aluminum (Al): 0.005 to 0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005 to 0.070%, vanadium (V): 0.001 to 0.300%, titanium (Ti): 0.001 to 0.050%, chromium (Cr): 0.01 to 0.30%, molybdenum (Mo): 0.01 to 0.12%, copper (Cu): 0.01 to 0.60%, nickel (Ni): 0.05 to 4.00%, calcium (Ca): 0.0005 to 0.0040%, and the balance being Fe and unavoidable impurities, The Ceq value defined by the following relational expression 1 is 0.40 to 0.65, The microstructure includes, in terms of area percentage, 80% or more of polygonal ferrite and the remainder of bainite in the surface layer portion, which is the region from the surface to the 1 / 8 point in the thickness direction, and 80% or more of one or more of bainite and martensite in the center portion, which is the remaining region from the 1 / 8 point to the 1 / 2 point in the thickness direction, In the central portion, the bainite or martensite has a packet average size of 25 μm or less having a high-angle grain boundary of 15° or more, The difference between the porosity of the area from the surface to the 3 / 8 point in the thickness direction and the porosity of the area from the 3 / 8 point to the 5 / 8 point in the thickness direction is 0.100 mm 3 / g or less, At -50°C, the Charpy impact energy absorption value can be 50J or more. [Equation 1] Ceq=[C]+[Mn] / 6+([Ni]+[Cu]) / 15+([Cr]+[Mo]+[V]) / 5 (In the formula, [C], [Mn], [Ni], [Cu], [Cr], [Mo], and [V] are the weight percentages of each element.)

[0015] The steel plate may further contain, by weight percent, zirconium (Zr): 0.001 to 0.150%.

[0016] The central portion may contain at least one of ferrite and pearlite as the remaining structure.

[0017] The steel sheet contains at least one of fine NbC, NbCN, VC, and CVN precipitates with a diameter of 5 to 50 nm, with a maximum size of 1 μm. 2 There can be 15 or more per prize.

[0018] The surface layer may have a hardness value of 250HB or less.

[0019] The steel plate may have a tensile strength of 590 to 820 MPa and a fatigue limit ratio (tensile strength / fatigue strength) of 0.3 or more.

[0020] The thickness of the steel plate may be 133 to 233 mm.

[0021] According to one embodiment of the present invention, the composition contains, in weight percent, carbon (C): 0.05 to 0.20%, silicon (Si): 0.05 to 0.5%, manganese (Mn): 1.00 to 2.00%, aluminum (Al): 0.005 to 0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005 to 0.070%, vanadium (V): 0.001 to 0.300%, titanium (T): 0.001 to 0.300%, and the like. i): 0.001 to 0.050%, chromium (Cr): 0.01 to 0.30%, molybdenum (Mo): 0.01 to 0.12%, copper (Cu): 0.01 to 0.60%, nickel (Ni): 0.05 to 4.00%, calcium (Ca): 0.0005 to 0.0040%, the balance being Fe and unavoidable impurities, and a Ceq value defined by the following relational formula 1 of 0.40 to 0.65 is a step of primarily reheating a steel slab; forging the primarily reheated steel slab at a cumulative reduction rate of 35 to 65% and a deformation rate of 1.0 to 4.0 / s; Secondly reheating the forged steel slab; hot rolling the second-reheated steel slab at a finish rolling temperature of Tnr-50 to Tnr+50°C and a cumulative reduction of 20.0% or more; heating the hot-rolled steel sheet to a temperature range of 820 to 900°C, holding the temperature for 10 to 40 minutes, and then primarily cooling the steel sheet to 700°C at an average cooling rate of 0.1 to 5.0°C / s based on the surface temperature of the steel sheet; Secondarily cooling the primarily cooled steel sheet to room temperature at an average cooling rate of 10.0°C / s or more based on the temperature of the steel sheet surface; The second-cooled steel sheet is heated to a temperature range of 550 to 700 ° C. and held for 5 to 60 minutes. It is possible to provide a method for producing a steel plate in which the cumulative reduction rate at the time of forging at or below the recrystallization temperature is 20% or less. [Equation 1] Ceq=[C]+[Mn] / 6+([Ni]+[Cu]) / 15+([Cr]+[Mo]+[V]) / 5 (In the formula, [C], [Mn], [Ni], [Cu], [Cr], [Mo], and [V] are the weight percentages of each element.)

[0022] The steel slab may further contain, by weight, zirconium (Zr): 0.001 to 0.150%.

[0023] The first reheating step is performed in a temperature range of 1100 to 1300°C, The second reheating step can be performed at a temperature range of 1000 to 1200°C.

[0024] During the first reheating, the thickness of the steel slab is 650 to 750 mm, After the forging, the thickness of the steel slab is 350-450 mm; After the hot rolling, the thickness of the steel plate may be 133 to 233 mm. [Effects of the Invention]

[0025] According to an embodiment of the present invention, a steel sheet and a manufacturing method thereof can be provided.

[0026] According to an embodiment of the present invention, an extra-thick steel material for flanges having excellent fatigue properties and low-temperature impact toughness and a method for manufacturing the same can be provided.

[0027] According to an embodiment of the present invention, it is possible to provide an extra-thick steel material for flanges having excellent strength and low-temperature impact toughness, which can be used in wind power towers and systems, and a method for manufacturing the same. [Brief explanation of the drawings]

[0028] [Figure 1] 1 is a CCT diagram showing a schematic cooling history of Example 1 according to an embodiment of the present invention and Comparative Example 7. DETAILED DESCRIPTION OF THE INVENTION

[0029] Preferred embodiments of the present invention will be described below. The embodiments of the present invention can be modified in various ways, and the scope of the present invention should not be construed as being limited to the embodiments described below. These embodiments are provided to further explain the present invention to those skilled in the art to which the invention pertains.

[0030] The present invention will be described in detail below.

[0031] The steel composition of the present invention will be described in detail below.

[0032] Unless otherwise specified in the present invention, the percentages indicating the contents of each element are based on weight.

[0033] The steel sheet according to one embodiment of the present invention contains, in weight percent, carbon (C): 0.05 to 0.20%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.00 to 2.00%, aluminum (Al): 0.005 to 0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005 to 0.070%, vanadium (V): 0.005 to 0.070%, and ): 0.001 to 0.300%, titanium (Ti): 0.001 to 0.050%, chromium (Cr): 0.01 to 0.30%, molybdenum (Mo): 0.01 to 0.12%, copper (Cu): 0.01 to 0.60%, nickel (Ni): 0.05 to 4.00%, calcium (Ca): 0.0005 to 0.0040%, and the balance being Fe and unavoidable impurities.

[0034] Carbon (C): 0.05~0.20% Carbon (C) is the most important element for ensuring basic strength, so it must be contained in steel within an appropriate range. To achieve this effect, 0.05% or more of carbon (C) may be added. According to one embodiment of the present invention, 0.10% or more of carbon (C) may be added. However, if the carbon (C) content exceeds a certain level, the hardenability during QT heat treatment may increase excessively, resulting in excessively high base material strength and hardness. This may lead to surface cracks during forging and reduced low-temperature impact toughness in the final product. Therefore, in the present invention, the upper limit of the carbon (C) content may be limited to 0.20%. According to one embodiment of the present invention, the upper limit may be 0.18%.

[0035] Silicon (Si): 0.05 to 0.50% Silicon (Si) is a substitutional element that improves the strength of steel through solid solution strengthening and has a strong deoxidizing effect, making it an essential element for producing clean steel. Therefore, silicon (Si) can be added in an amount of 0.05% or more. According to one embodiment of the present invention, silicon (Si) can be added in an amount of 0.20% or more. However, if a large amount of silicon (Si) is added, it can form a martensite-austenite (MA) phase, excessively increase the strength of the ferrite matrix, and cause deterioration in the surface quality of extra-thick steel products. Therefore, the upper limit of silicon content can be limited to 0.50%. According to one embodiment of the present invention, the upper limit can be 0.40%.

[0036] Manganese (Mn): 1.00-2.00% Manganese (Mn) is a useful element that improves strength through solid solution strengthening and improves hardenability by forming a low-temperature transformation phase. Therefore, to ensure a tensile strength of 590 MPa or more, 1.00% or more may be added. According to one embodiment of the present invention, the manganese (Mn) content may be 1.10% or more. However, as the manganese (Mn) content increases, manganese (Mn) may form elongated non-metallic inclusions (MnS) with S, reducing toughness and acting as an impact initiation point, which may rapidly reduce the low-temperature impact toughness of the product. Therefore, the manganese (Mn) content may be limited to 2.00% or less. According to one embodiment of the present invention, the manganese (Mn) content may be 1.50% or less.

[0037] Aluminum (Al): 0.005 to 0.100% Aluminum (Al), along with Si, is a powerful deoxidizer in the steelmaking process. To achieve this effect, 0.005% or more of Al may be added. According to one embodiment of the present invention, the lower limit of the aluminum (Al) content may be 0.010%. However, excessive aluminum (Al) content can lead to an excessive increase in the Al2O3 fraction in the oxidized inclusions formed as a result of deoxidation, resulting in coarse inclusion size, making removal of the inclusions difficult during refining and potentially reducing low-temperature impact toughness. Therefore, the aluminum (Al) content may be limited to 0.100% or less. According to one embodiment of the present invention, the aluminum (Al) content may be 0.070% or less.

[0038] Phosphorus (P): 0.0100% or less Phosphorus (P) is an element that induces embrittlement at grain boundaries or by forming coarse inclusions, so to improve resistance to brittle crack propagation, the phosphorus (P) content can be limited to 0.0100% or less. However, 0% is excluded in consideration of the level that is inevitably contained during steel manufacturing.

[0039] Sulfur (S): 0.0150% or less Sulfur (S) is an element that induces embrittlement at grain boundaries or by forming coarse inclusions, so to improve resistance to brittle crack propagation, sulfur (S) can be limited to 0.0015% or less. However, 0% is excluded in consideration of the level that is unavoidably contained during steel manufacturing.

[0040] Niobium (Nb): 0.005 to 0.070% Niobium (Nb) is an element that precipitates in the form of NbC or NbCN to improve the strength of the base material. Furthermore, Nb, which dissolves during reheating at high temperatures, precipitates very finely in a modified organic form below the recrystallization temperature during forging, suppressing the growth of austenite and thereby refining the microstructure. Therefore, niobium (Nb) may be added in an amount of 0.005% or more. According to one embodiment of the present invention, it may be added in an amount of 0.010% or more. However, if excessive niobium (Nb) is added, undissolved niobium (Nb) forms in the form of TiNb(C,N), which can reduce low-temperature impact toughness. Therefore, the upper limit of the niobium (Nb) content may be limited to 0.070%. According to one embodiment of the present invention, the niobium (Nb) content may be 0.065% or less.

[0041] Vanadium (V): 0.001 to 0.300% Vanadium (V) is almost completely dissolved during reheating, so its strengthening effect due to precipitation or dissolution during subsequent rolling is minimal. However, in the case of extra-thick forged steel, the air-cooling rate is very slow, so it precipitates as very fine carbonitrides during the cooling or tempering heat treatment, thereby improving strength. To fully achieve this effect, vanadium (V) can be added in an amount of 0.001% or more. According to one embodiment of the present invention, the lower limit of the vanadium (V) content can be 0.010%. However, excessive vanadium (V) content can excessively increase the surface hardness of the slab due to its high hardening ability, which can cause surface cracks during flanging and significantly increase manufacturing costs, making it commercially unreasonable. Therefore, the vanadium (V) content can be 0.300% or less. According to one embodiment of the present invention, the vanadium (V) content can be 0.250% or less.

[0042] Titanium (Ti): 0.001 to 0.050% Titanium (Ti) precipitates as TiN during reheating, inhibiting the growth of prior austenite grains at high temperatures and significantly improving low-temperature toughness. To achieve this effect, 0.001% or more of Ti can be added. However, excessive Ti can cause clogging of the continuous casting nozzle and reduced low-temperature toughness due to crystallization in the center. Furthermore, Ti combines with N to form coarse TiN precipitates in the center of the thickness, reducing the elongation of the product and potentially causing surface cracks during the forging process. Therefore, the Ti content can be 0.050% or less. According to one embodiment of the present invention, the Ti content can be 0.030% or less, or even 0.018% or less.

[0043] Chromium (Cr): 0.01 to 0.30% Chromium (Cr) slows the spheroidization rate of cementite, thereby preventing a decrease in strength and improving hardenability during the cooling process. To achieve this effect, 0.01% or more of chromium (Cr) can be added. However, excessive chromium (Cr) content can increase the size and fraction of Cr-rich coarse carbides, such as M23C6, resulting in a decrease in impact toughness of the product. Furthermore, the solid solubility of Nb in the product and the fraction of fine precipitates, such as NbC, can decrease, resulting in a decrease in strength of the product. Therefore, the upper limit of chromium (Cr) content can be limited to 0.30%. According to one embodiment of the present invention, the upper limit of chromium (Cr) content can be 0.25%.

[0044] Molybdenum (Mo): 0.01 to 0.12% Molybdenum (Mo) is an element that increases grain boundary strength, increases hardening ability, and dissolves in precipitates to improve strength, effectively contributing to increased strength and ductility of products. Molybdenum (Mo) also has the effect of preventing a decrease in toughness due to the grain boundary segregation of impurities such as P. For this effect, 0.01% or more of Mo can be added. However, since Mo is an expensive element and adding too much significantly increases manufacturing costs, the upper limit of its content can be limited to 0.12%.

[0045] Copper (Cu): 0.01 to 0.60% Copper (Cu) is an element that significantly improves the strength of the matrix phase by solid solution strengthening in ferrite. To achieve this effect, copper (Cu) can be contained in an amount of 0.01% or more. A more preferable copper (Cu) content is 0.03% or more. However, excessive copper (Cu) content can increase the possibility of star cracks on the surface of the steel sheet, and as an expensive element, it can significantly increase manufacturing costs. Therefore, the present invention limits the copper (Cu) content to 0.60%. According to one embodiment of the present invention, the upper limit can be 0.50%.

[0046] Nickel (Ni): 0.05-4.00% Nickel (Ni) is an element that effectively contributes to improving impact toughness by increasing stacking faults at low temperatures and facilitating cross-slip of potentials, improving hardenability, and improving solid solution hardening, thereby improving strength. To achieve these effects, 0.05% or more of nickel (Ni) may be added. According to one embodiment of the present invention, the nickel (Ni) content may be 0.10% or more. However, because excessive addition of nickel (Ni) can increase manufacturing costs due to its high cost, the upper limit of the nickel (Ni) content may be limited to 4.00%. According to one embodiment of the present invention, the upper limit of the nickel (Ni) content may be 3.50%.

[0047] Calcium (Ca): 0.0005-0.0040% When calcium (Ca) is added after deoxidation with Al, it combines with S, which forms MnS inclusions, to suppress the formation of MnS, while at the same time forming spherical CaS, which suppresses cracking due to hydrogen-induced cracking. To ensure that the impurity S is converted into CaS, calcium (Ca) can be added in amounts of 0.0005% or more. However, if the amount added is excessive, the remaining calcium (Ca) from the CaS formation combines with O to form coarse oxidized inclusions, which can be stretched and fractured during forging, reducing elongation and low-temperature impact toughness. Therefore, the upper limit of calcium (Ca) content is limited to 0.0040%.

[0048] In addition to the above-mentioned composition, the steel material of the present invention may contain the remaining iron (Fe) and inevitable impurities. The inevitable impurities cannot be excluded because they may be unintentionally mixed in during the normal manufacturing process. These impurities are known to anyone skilled in the field of normal steel manufacturing, so the details of all of them will not be specifically mentioned in this specification.

[0049] According to an embodiment of the present invention, the composition may further contain, by weight, zirconium (Zr): 0.001 to 0.150%.

[0050] Zirconium (Zr) is a strong carbide-forming element and can exist in the form of ZrC. Like VC and NbC, it can improve the strength of the matrix phase in a precipitation-strengthening form. For this effect, 0.001% or more of zirconium (Zr) can be added. However, if added in excess, the high cost can increase manufacturing costs, so the upper limit can be limited to 0.150%.

[0051] The steel sheet according to an embodiment of the present invention may have a Ceq value defined by the following Relation 1 of 0.40 to 0.65. [Equation 1] Ceq=[C]+[Mn] / 6+([Ni]+[Cu]) / 15+([Cr]+[Mo]+[V]) / 5 (In the formula, [C], [Mn], [Ni], [Cu], [Cr], [Mo], and [V] are the weight percentages of each element.)

[0052] If the Ceq value defined by Relational Formula 1 is less than 0.40, the hardenability will be insufficient, and there may be a problem in ensuring the appropriate tensile strength of 590 to 820 MPa required in the present invention. On the other hand, if the value exceeds 0.65, the strength value will be too high, and there is a risk of deterioration in low-temperature impact toughness at -50°C. Therefore, it is preferable that Ceq is 0.40 to 0.65.

[0053] The steel microstructure of the present invention will now be described in detail.

[0054] In the present invention, unless otherwise specified, the percentage representing the fraction of the microstructure is based on the area.

[0055] In the microstructure of a steel sheet according to an embodiment of the present invention, the surface layer region, which is the region from the surface to the 1 / 8 point in the thickness direction, may contain 80% or more polygonal ferrite and the remainder bainite, in terms of area%, and the remaining region, the center region, which is the region from the 1 / 8 point to the 1 / 2 point in the thickness direction, may contain 80% or more of one or more of bainite and martensite.

[0056] In the present invention, in order to improve crack propagation resistance by ensuring a soft surface structure, the surface layer portion, which is the region from the surface to the 1 / 8 point in the thickness direction, can be limited to contain 80% or more polygonal ferrite. If the polygonal ferrite fraction is less than 80%, the crack propagation resistance of the surface layer portion decreases and the fatigue limit ratio of the steel material cannot be ensured at the desired level.

[0057] The remaining region, the central region from the 1 / 8 point to the 1 / 2 point, can contain 80% or more of one or more of bainite and martensite to ensure the desired strength, and the remaining structure can contain one or more of ferrite and pearlite.

[0058] According to one embodiment of the present invention, the bainite or martensite in the central portion may have a packet average size of 25 μm or less, which has a high hardness grain boundary of 15° or more.

[0059] When the matrix structure is bainite or martensite, cracks may propagate along packet boundaries based on high-hardness grain boundaries during impact testing. Therefore, if the packet size is large, DBTT may increase, resulting in a deterioration in impact toughness. Therefore, in the present invention, the average packet size can be limited to 25 μm or less to ensure the desired impact toughness.

[0060] According to one embodiment of the present invention, the difference between the porosity of the region from the surface to the 3 / 8 point in the thickness direction and the porosity of the region from the 3 / 8 point to the 5 / 8 point in the thickness direction is 0.100 mm 3 / g or less.

[0061] That is, the porosity of the area from the surface to the 3 / 8 point in the thickness direction of the steel plate is subtracted from the porosity of the area from the 3 / 8 point to the 5 / 8 point in the thickness direction of the steel plate to be 0.100 mm 3 / g or less. Here, the porosity is defined as the density (g / mm 3 ) and measure the reciprocal (mm 3 / g) can be calculated.

[0062] The difference above is 0.100 mm 3 If the difference exceeds 0.050 mm / g, the residual voids act as crack initiation points during impact testing, making it difficult to ensure the desired low-temperature impact toughness at -50°C. 3According to one embodiment of the present invention, the difference may be 0.030 mm / g or less. 3 The lower limit of the difference does not need to be particularly limited, but in one embodiment, the lower limit is 0 mm 3 / g.

[0063] According to one embodiment of the present invention, one or more of fine NbC, NbCN, VC, and CVN precipitates having a diameter of 5 to 50 nm are present in a range of 1 μm. 2 There can be 15 or more per prize.

[0064] If there are less than 15 NbC precipitates of 5 to 50 nm, the precipitation strengthening effect weakens and the pinning effect and precipitation strengthening effect are lost, making it difficult to ensure the level of strength required in the present invention.

[0065] The steel plate according to one embodiment of the present invention has a thickness of 133 to 233 mm, a hardness value of the surface layer of 250 HB or less, a tensile strength of 590 to 820 MPa, a Charpy impact absorption energy value of 50 J or more at -50°C, and a fatigue limit ratio (tensile strength / fatigue strength) of 0.30 or more.

[0066] The steel plate according to an embodiment of the present invention may be manufactured into a flange having a maximum thickness of 200 to 500 mm, an inner diameter of 4000 to 7000 mm, and an outer diameter of 5000 to 8000 mm.

[0067] The steel production method of the present invention will be described in detail below.

[0068] The steel sheet according to an embodiment of the present invention may be manufactured by subjecting a steel slab having the alloy composition of the present invention to primary reheating, forging, secondary reheating, hot rolling, primary cooling, secondary cooling, and tempering.

[0069] 1st reheating A steel slab having a thickness of 650 to 750 mm and satisfying the alloy composition of the present invention can be primarily reheated in a temperature range of 1100 to 1300°C.

[0070] The primary reheating can be performed at a temperature range of 1100°C or higher to redissolve Ti and Nb complex carbonitrides or TiNb(C,N) coarse crystals formed during casting, homogenize the structure by heating the austenite before forging to above the recrystallization temperature and holding it there, and minimize surface cracks that may occur during the forging process by ensuring a sufficiently high forging end temperature. However, reheating steel slabs at excessively high temperatures can cause problems due to oxide scale at high temperatures, and the increased costs associated with heating and holding can lead to excessive increases in manufacturing costs, so the upper limit of the steel slab heating temperature can be limited to 1300°C.

[0071] According to an embodiment of the present invention, the thickness of the steel slab may be 650 to 750 mm during the first reheating. According to an embodiment of the present invention, the thickness of the steel slab may be 700 to 750 mm.

[0072] Meanwhile, according to one embodiment of the present invention, the slab may be manufactured using any one of a continuous casting process, a semi-continuous casting process, and an ingot casting process.

[0073] forging The above-mentioned primarily reheated steel slab can be forged at a cumulative reduction rate of 35 to 65% and a deformation rate of 1.0 to 4.0 / s.

[0074] Forging is the process of forming a heated steel slab into the final desired intermediate material shape. To fully compress the voids, high-deformation, low-speed forging is essential, so the cumulative reduction rate of forging is 35-65% and the deformation speed is 1.0 / s-4.0 / s.

[0075] In the present invention, the deformation rate means the rate of deformation per unit time, and its unit is % / s.

[0076] If the cumulative reduction rate of forging is less than 35%, the voids remaining in the steel slab may not be fully closed, resulting in residual voids, which may reduce the impact toughness of the product. The cumulative reduction rate of forging according to one embodiment of the present invention may be 40% or more. If the cumulative reduction rate exceeds 65% and the number of forging passes increases, the surface temperature may continuously decrease, potentially causing surface cracks at low temperatures.

[0077] On the other hand, if the cumulative reduction below the non-recrystallization temperature exceeds 20%, which is the temperature at which dislocation density is not recovered or compensated for by recrystallization, the uniform elongation of the surface of the steel material at approximately 850°C to 1150°C where forging is performed will exceed 20%. Therefore, the work hardening of the accumulated dislocations significantly reduces the uniform elongation of the surface, which can lead to surface cracks during the forging process. Therefore, the cumulative reduction below the recrystallization temperature should be limited to 20% or less. In one embodiment of the present invention, it can be 15% or less.

[0078] If the forging deformation rate is less than 1.0 / s, productivity may decrease excessively, while if the deformation rate exceeds 4.0 / s, internal cracks may occur during the forging process due to work hardening, which may reduce impact toughness and fatigue properties.

[0079] According to one embodiment of the present invention, the thickness of the steel slab after forging may be 350-450 mm.

[0080] 2nd reheating The forged steel slab can be secondarily reheated in the temperature range of 1000 to 1200°C.

[0081] Reheating can be performed at temperatures above 1000°C to homogenize the structure by heating and holding the austenite before hot rolling above its recrystallization temperature, and to minimize inclusion breakage during the rolling process by ensuring a sufficiently high end temperature. However, heating slabs at excessively high temperatures can cause problems due to oxide scale at high temperatures, and the increased costs associated with heating and holding can lead to excessive increases in manufacturing costs. Therefore, the upper limit of the secondary heating temperature can be limited to 1200°C.

[0082] hot rolling The secondarily reheated steel slab can be hot rolled at a finish rolling temperature of Tnr-50 to Tnr+50°C and a cumulative reduction of 20.0% or more.

[0083] If the finish rolling temperature during hot rolling is less than Tnr - 50°C, the deformation resistance increases excessively as the temperature decreases, making it difficult to sufficiently refine the austenite grains in the center of the product thickness direction, which can result in a deterioration of the low-temperature impact toughness of the final product.On the other hand, if the finish rolling temperature exceeds Tnr + 50°C, the austenite grains become excessively coarse, which can result in a deterioration of strength and impact toughness.

[0084] If the cumulative rolling reduction is less than 20.0%, there is a problem in that the prior austenite grain size cannot be sufficiently refined.

[0085] According to one embodiment of the present invention, the thickness of the steel sheet after hot rolling may be 133 to 233 mm. [formula] Tnr=887+464[C]+890[Ti]+363[Al]-357[Si]+6445[Nb]-644([Nb]1 / 2)+732[V]-230([V]1 / 2) (In the formula, [C], [Ti], [Al], [Si], [Nb], and [V] are the weight percentages of each element.)

[0086] Primary cooling The hot-rolled steel sheet is heated to a temperature range of 820 to 900°C and held for 10 to 40 minutes, and then primarily cooled to 700°C at an average cooling rate of 0.1 to 5.0°C / s based on the surface temperature of the steel sheet.

[0087] In the present invention, the primary cooling can be performed to form the desired polygonal ferrite from the surface of the steel material to a region corresponding to a 1 / 8 point in the thickness direction. The temperature and cooling rate during the primary cooling are based on the surface temperature of the steel sheet.

[0088] If the reheating temperature is below 820°C or the holding time is less than 10 minutes, the carbides formed during cooling after rolling and the impurity elements segregated at the grain boundaries do not redissolve smoothly, which can significantly reduce the through-thickness elongation (ZRA) and low-temperature toughness of the steel after heat treatment.On the other hand, if the temperature exceeds 900°C or the holding time exceeds 40 minutes, the impact toughness quality can be reduced due to the coarsening of austenite and precipitate phases such as Nb(C,N) and V(C,N).

[0089] If the average cooling rate during the primary cooling is less than 0.1°C / s, the waiting time in air is too long, which can increase manufacturing time and reduce productivity.On the other hand, if the cooling rate exceeds 5.0°C / s, polygonal ferrite is not formed and bainite or martensite, a low-temperature transformation structure, is formed, which can reduce the uniform elongation rate of the surface layer of the final product and deteriorate fatigue quality.

[0090] Secondary cooling The primarily cooled steel sheet can be secondarily cooled to room temperature at an average cooling rate of 10.0°C / s or more based on the temperature of the steel sheet surface.

[0091] The secondary cooling can be carried out in order to make the microstructure of the region other than the surface layer contain bainite or a two-phase structure of bainite and ferrite through strong cooling.

[0092] During the secondary cooling, the average cooling rate may be based on the surface temperature of the steel sheet. If the average cooling rate is less than 10.0°C / s, it may be difficult to obtain the above-mentioned low-temperature transformation structure. In the present invention, the upper limit of the average cooling rate during the secondary cooling is not particularly limited, but it may be limited to 200.0°C / s or less. Meanwhile, the secondary cooling involves quenching, which may be performed by slowing the steel sheet threading speed and increasing the flow rate of sprayed water.

[0093] Tempering heat treatment The secondarily cooled steel sheet can be subjected to a tempering heat treatment by heating it to a temperature range of 550 to 700°C and holding it for 5 to 60 minutes.

[0094] The tempering heat treatment reduces the dislocation density of the bainite or mixed structure of bainite and ferrite, which is a low-temperature transformation structure, and diffuses carbon over a short range, thereby improving strength and toughness.

[0095] If the tempering temperature is less than 550°C, the carbon diffusion is insufficient, resulting in excessive strength and reduced toughness. On the other hand, if the temperature exceeds 700°C, fresh martensite is formed due to reverse transformation at temperatures above Ac1, resulting in significant degradation of impact toughness and cold workability.

[0096] If the tempering heat treatment time is less than 5 minutes, the time for sufficient diffusion of carbon during the tempering process is insufficient, resulting in excessively excessive strength, which may result in reduced toughness and cold workability, and may fall outside the appropriate strength range required in the present invention. On the other hand, if the tempering heat treatment time is more than 60 minutes, excessive heating may cause cementite to spheroidize, resulting in a rapid decrease in strength. [Example]

[0097] The present invention will be described in more detail with reference to the following examples. However, it should be noted that the following examples are intended to illustrate and explain the present invention in more detail, and are not intended to limit the scope of the present invention.

[0098] (Example) A 700 mm thick slab was manufactured using the alloy composition shown in Table 1. Using the slab, forging, hot rolling, and tempering after primary and secondary cooling were carried out according to the process conditions shown in Table 2 below, ultimately producing a 320 mm thick flange. The primary reheating temperature was commonly 1220-1240°C, the secondary reheating temperature was commonly 1140-1160°C, and the reheating time for primary cooling and tempering time was commonly 30 minutes. Meanwhile, the hot rolling temperature was indicated based on the Tnr temperature.

[0099] [Table 1] *The units of P, S, Ca and O are ppm. [Equation 1] Ceq=[C]+[Mn] / 6+([Ni]+[Cu]) / 15+([Cr]+[Mo]+[V]) / 5 (In the formula, [C], [Mn], [Ni], [Cu], [Cr], [Mo], and [V] are the weight percentages of each element.)

[0100] [Table 2]

[0101] The microstructure and physical properties of each specimen were measured and are shown in Table 3 below. The microstructure was measured using an automatic image analyzer on specimens taken from the surface region (region from the surface to the 1 / 8th point in the thickness direction) and the remaining region (region from the 1 / 8th point to the 1 / 2nd point in the thickness direction). The packet size was calculated by automatically measuring the grain size based on the high-angle grain boundary with an orientation relationship of 15° or more on an Orientation Image Microscopy (OIM) map of the microstructure analyzed using Electron Backscatter Diffraction (EBSD).

[0102] In the present invention, the difference in porosity is expressed as the value obtained by subtracting the porosity of the region from the surface to the 3 / 8 point in the thickness direction from the porosity of the region from the 3 / 8 point to the 5 / 8 point in the thickness direction. Here, the porosity is expressed as the density (g / mm 3 ) and measure the reciprocal (mm 3 / g) was calculated.

[0103] The number of 5-50 nm precipitates observed in the cross section of the steel was measured using TEM. NbC precipitates were identified through NbC and VC diffraction patterns and EDX mapping, and 1 μm precipitates were identified. 2 The number of precipitates located at the position was counted.

[0104] The hardness values ​​were measured at three locations in the surface layer using a Brinell hardness tester and the average value was shown. Tensile strength was evaluated by a room temperature tensile test, and the impact toughness of each test piece was calculated by using the average absorbed energy values ​​measured three times at the same temperature via a Charpy V-Notch test. Fatigue strength was measured using rotating bending fatigue test pieces, electrolytically polished to remove the effects of the surface worked layer, and then measured according to the method described in JIS Z 2274. The fatigue limit ratio was calculated and shown as the ratio of fatigue strength to tensile strength.

[0105] After visually observing the surface of each test piece, grinding was carried out at the point where a surface crack had formed, and the grinding length until the crack disappeared was measured as the surface crack length.

[0106] [Table 3] *F: Polygonal ferrite, B: Bainite, M: Martensite

[0107] As shown in Table 3, in the case of the invention examples that satisfy the alloy composition and manufacturing conditions of the present invention, the microstructure characteristics proposed in the present invention were satisfied and the physical properties aimed at in the present invention were also secured.

[0108] Fig. 1 is a CCT diagram showing the cooling histories of Example 1 and Comparative Example 7. In Example 1, multi-stage cooling was used to control the surface layer to a polygonal ferrite structure and the center to a bainite structure, thereby ensuring all of the fatigue properties and strength targeted by the present invention. On the other hand, in Comparative Example 7, cooling was performed using only deep cooling, and low-temperature transformed structures were formed in both the surface layer and center, resulting in deterioration of fatigue properties.

[0109] On the other hand, Comparative Examples 1 and 2 are examples in which the cumulative reduction during forging was outside the range of the present invention. In Comparative Example 1, the cumulative reduction during forging was excessive, resulting in work hardening and a deterioration in low-temperature impact toughness. In Comparative Example 2, the cumulative reduction during forging was insufficient, resulting in inadequate control of residual voids in the center of the steel sheet, resulting in a deterioration in fatigue strength and impact toughness.

[0110] Comparative Example 3 is an example in which the rolling reduction below the recrystallization temperature exceeded the range of the present invention, resulting in the occurrence of surface defects.

[0111] Comparative Example 4 is an example in which the deformation rate during forging exceeded the range proposed by the present invention, resulting in excessive work hardening and making it impossible to ensure appropriate physical properties.

[0112] Comparative Example 5 is an example in which the finish rolling temperature was not reached during hot rolling, resulting in a coarse packet size in the center and excessively high hardness in the surface layer, resulting in reduced fatigue crack resistance.

[0113] Comparative Example 6 is an example in which the cumulative reduction rate was not reached during hot rolling, which meant that the desired packet size could not be achieved, resulting in a decrease in impact toughness.

[0114] Comparative Example 7 is an example in which the cooling rate during primary cooling exceeded the range proposed by the present invention, resulting in the formation of martensite, a hard structure, in the surface layer, making the surface layer too hard and resulting in fatigue crack resistance that did not meet the range of the present invention.

[0115] In Comparative Example 8, the cooling rate was not reached during the secondary cooling, and excessive ferrite was formed in the center, making it impossible to ensure appropriate strength, and the strength targeted in the present invention could not be achieved.

[0116] In Comparative Example 9, the tempering temperature was excessively high, and the precipitates were not smoothly formed, so that the strength aimed at in the present invention could not be ensured.

[0117] Comparative Examples 10 and 11 are examples in which the carbon content exceeded the range proposed by the present invention. As a result, the carbon content of Comparative Example 10 was insufficient, and the strength targeted by the present invention could not be ensured. In Comparative Example 11, the carbon content exceeded the range proposed by the present invention, and the appropriate level of strength could not be ensured, resulting in poor impact toughness.

[0118] Comparative Example 12 was a case where the Mn content was excessive, and bainite was formed in an excessive amount in the surface layer portion, resulting in failure to ensure adequate strength and toughness.

[0119] Comparative Example 13 is an example in which the Nb content did not reach the content proposed in the present invention, and as a result, precipitates were not easily formed, and the desired strength could not be ensured.

[0120] Comparative Example 14 is an example in which the V content exceeded the range of the present invention, resulting in a deterioration in low-temperature impact toughness.

[0121] Although the present invention has been described in detail with reference to the examples above, other embodiments are possible, and the spirit and scope of the claims set forth below should not be limited to the examples.

Claims

1. In weight percent, carbon (C): 0.05-0.20%, silicon (Si): 0.05-0.50%, manganese (Mn): 1.00-2.00%, aluminum (Al): 0.005-0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005-0.070%, vanadium (V):

0. 0.001 to 0.300%, titanium (Ti): 0.001 to 0.050%, chromium (Cr): 0.01 to 0.30%, molybdenum (Mo): 0.01 to 0.12%, copper (Cu): 0.01 to 0.60%, nickel (Ni): 0.05 to 4.00%, calcium (Ca): 0.0005 to 0.0040%, the balance being Fe and inevitable impurities, The microstructure includes a surface layer portion, which is a region from the surface to the 1 / 8 point in the thickness direction, containing 80% or more polygonal ferrite and the remainder bainite, in terms of area percent, and a central portion, which is a region from the 1 / 8 point to the 1 / 2 point in the thickness direction, containing 80% or more of one or more of bainite and martensite, The bainite or martensite in the central portion has a packet average size of 25 μm or less having a high-angle grain boundary of 15° or more, The difference between the porosity of the region from the surface to the 3 / 8 point in the thickness direction and the porosity of the region from the 3 / 8 point to the 5 / 8 point in the thickness direction is 0.100 mm ３ / g or less, A steel plate having a Charpy impact absorption energy value of 50 J or more at -50°C.

2. The steel plate according to claim 1, further comprising, by weight percent, zirconium (Zr): 0.001 to 0.150%.

3. The steel sheet according to claim 1, characterized in that the steel sheet has a Ceq value defined by the following relational expression 1 of 0.40 to 0.

65. [Relationship 1] Ceq=[C]+[Mn] / 6+([Ni]+[Cu]) / 15+([Cr]+[Mo]+[V]) / 5 (In the formula, [C], [Mn], [Ni], [Cu], [Cr], [Mo], and [V] are the weight percentages of each element.)

4. The steel plate according to claim 1 , wherein the central portion contains at least one of ferrite and pearlite as a remaining structure.

5. The steel sheet contains at least one of fine NbC, NbCN, VC, and CVN precipitates having a diameter of 5 to 50 nm, with a size of 1 μm or more. ２ 2. The steel plate according to claim 1, wherein the number of defects is 15 or more per one.

6. The steel sheet according to claim 1, wherein the surface layer portion has a hardness value of 250 HB or less.

7. 2. The steel plate according to claim 1, wherein the steel plate has a tensile strength of 590 to 820 MPa and a fatigue limit ratio (tensile strength / fatigue strength) of 0.30 or more.

8. The steel plate according to claim 1, wherein the steel plate has a thickness of 133 to 233 mm.

9. In weight percent, carbon (C): 0.05 to 0.20%, silicon (Si): 0.05 to 0.5%, manganese (Mn): 1.00 to 2.00%, aluminum (Al): 0.005 to 0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005 to 0.070%, vanadium (V): 0.001 to 0.30 a step of primarily reheating a steel slab consisting of 0.0% Fe, 0.001 to 0.050% titanium (Ti), 0.01 to 0.30% chromium (Cr), 0.01 to 0.12% molybdenum (Mo), 0.01 to 0.60% copper (Cu), 0.05 to 4.00% nickel (Ni), 0.0005 to 0.0040% calcium (Ca), and the balance being Fe and unavoidable impurities; forging the primarily reheated steel slab at a cumulative reduction rate of 35 to 65% and a deformation rate of 1.0 to 4.0 / s; Secondly reheating the forged steel slab; hot rolling the second reheated steel slab at a finish rolling temperature of Tnr-50 to Tnr+50°C and a cumulative reduction of 20.0% or more; heating the hot-rolled steel sheet to a temperature range of 820 to 900°C, holding the temperature for 10 to 40 minutes, and then primarily cooling the steel sheet to 700°C at an average cooling rate of 0.1 to 5.0°C / s based on the surface temperature of the steel sheet; Secondarily cooling the primarily cooled steel sheet to room temperature at an average cooling rate of 10°C / s or more based on the temperature of the steel sheet surface; a tempering heat treatment step of heating the second-cooled steel sheet to a temperature range of 550 to 700°C and holding the temperature for 5 to 60 minutes; A method for producing a steel plate, characterized in that during forging, a cumulative reduction rate at a temperature equal to or lower than the recrystallization temperature is 20% or less.

10. The method for manufacturing a steel plate according to claim 9, wherein the steel slab further contains, by weight percent, zirconium (Zr): 0.001 to 0.150%.

11. The method for manufacturing a steel plate according to claim 9, wherein the steel slab has a Ceq value defined by the following relational expression 1 of 0.40 to 0.

65. [Relationship 1] Ceq=[C]+[Mn] / 6+([Ni]+[Cu]) / 15+([Cr]+[Mo]+[V]) / 5 (In the formula, [C], [Mn], [Ni], [Cu], [Cr], [Mo], and [V] are the weight percentages of each element.)

12. The first reheating step is performed at a temperature range of 1100 to 1300°C, The method of claim 9, wherein the second reheating is performed at a temperature in the range of 1000 to 1200°C.

13. During the first reheating, the thickness of the steel slab is 650 to 750 mm, After said forging, the thickness of the steel slab is 350-450 mm; The method for producing a steel sheet according to claim 9, wherein after the hot rolling, the thickness of the steel sheet is 133 to 233 mm.

Images