Extra-thick steel material with excellent low-temperature impact toughness and its manufacturing method

Abstract

The present invention provides an extra-thick steel material that has excellent low-temperature impact toughness after long-term PWHT even when the steel plate is thick, and a method for manufacturing the same. [Solution] In weight percent, C: 0.10-0.25%, Si: 0.05-0.50%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.010% or less, S: 0.0015% or less, Nb: 0.001-0.03%, V: 0.001-0.03%, Ti: 0.001-0.03%, Cr: 0.01-0.20%, Mo: 0.01-0.1 The core microstructure in the range of t / 4 to t / 2 is composed of a composite structure of 35 to 40% ferrite and the remainder bainite, with a bainite packet size of 10 μm or less and a central porosity of 0.1 mm 3 / g or less, the depth of surface cracks is 0.5 mm or less, and the cross-sectional hardness of the center portion is 200HB or less.

Description

[Technical field]

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

[0002] Recently, as crude oil refining and storage facilities have become larger and larger capacity storage has been required, there has been a continuous increase in the demand for thicker steel materials used in these facilities. In particular, the temperature required to ensure low-temperature impact toughness has gradually decreased due to increased use in cold environments.

[0003] In manufacturing large structures, there is a trend to control defects in steel materials such as non-metallic inclusions, segregation, and internal voids to the utmost in order to improve the external soundness of the steel materials. In addition, there is a demand to lower the carbon equivalent (Ceq) to ensure the structural stability of not only the base material but also the heat-affected zone after welding.

[0004] In particular, for very thick materials with a thickness of over 100 mm, the rolling reduction ratio is not high compared to thin materials, so unsolidified shrinkage holes that occur during continuous casting or casting are not fully compressed during the rough rolling process and remain in the form of residual voids in the center of the product. These residual voids act as starting points for cracks in structures when impacted, and ultimately cause the entire equipment to break due to a decrease in low-temperature impact toughness. Therefore, a process is required to fully compress the central void so that no residual voids remain before rolling.

[0005] In this regard, Patent Document 1 corresponds to a strong reduction technology in a rough rolling process of thick plates, and utilizes a technology for determining a limit reduction ratio for each thickness at which plate biting occurs from a strong reduction ratio for each pass set to approach the design allowable value (load and torque) of the rolling mill, a technology for distributing the reduction ratio by adjusting the index of the thickness ratio for each pass to ensure the target thickness of the rough rolling mill, and a technology for correcting the reduction ratio based on the limit reduction ratio for each thickness so that plate biting does not occur, thereby providing a manufacturing method that can apply an average reduction ratio of 27.5% in the final three passes of rough rolling based on an 80 mm standard. However, in the case of the above rolling method, the average reduction ratio of the entire product thickness is measured, and there is a disadvantage that it is difficult to apply a high deformation to the center of an extremely thick material with a maximum thickness of 250 mm where residual voids exist.

[0006] Meanwhile, the thicker the steel, the higher the post-weld heat treatment (PWHT) temperature or time. PWHT is a method to prevent deformation of structures and ensure shape and dimensional stability by removing residual stress in the welded parts. Usually, PWHT is performed on the entire structure, but even if it is performed locally, the base material other than the welded parts is also exposed to the heat source, which can cause deterioration of the physical properties of the base material. As a result, in the case of very thick materials, the quality of the base material can deteriorate after high temperature and long PWHT heat treatment, which can shorten the equipment life of the manufactured pressure vessel. During such PWHT, in the case of high-strength pressure vessel steels consisting of hard phases such as bainite, martensite, and island martensite (MA), the base material not only loses strength through a series of processes such as carbon re-diffusion, dislocation recovery, grain growth (interface migration of bainite or martensite), and carbide growth and precipitation, but also shows a tendency to increase the ductile-brittle transition temperature (DBTT).

[0007] As a means to prevent the deterioration of physical properties due to long-term PWHT at high temperatures, the first method is to increase the amount of alloying elements that can increase hardening ability even if Ceq is high, and to increase the fraction of tempered low-temperature phases after heat treatment to reduce the amount of strength loss. The second method is to increase the content of elements with solid solution strengthening effects, such as Mo, Cu, Si, and C, in order to increase the strength of the ferrite matrix phase without changes in structure and dislocation density after heat treatment, while realizing a two-phase structure consisting of ferrite and bainite in the microstructure of QT (Quenching-Tempering) steel, or a three-phase structure that includes a certain amount of martensite in addition to the above structures.

[0008] However, both of the above-mentioned two methods have the disadvantage that the toughness of the heat affected zone (HAZ) of the weld is likely to decrease due to an increase in Ceq, and the manufacturing cost increases due to the addition of solid solution strengthening elements.

[0009] Another method is a precipitation strengthening method using rare earth elements, which is effective under specific component ranges and application temperature conditions. Related Patent Document 2 discloses that the PWHT guaranteed time can be up to 16 hours by heating and hot rolling a slab consisting of, by weight percent, 0.05-0.20% C, 0.02-0.5% Si, 0.2-2.0% Mn, 0.005-0.10% Al, the balance being Fe and unavoidable impurities, and further including one or more of Cu, Ni, Cr, Mo, V, Nb, Ti, B, Ca, and rare earth elements as necessary, and then air-cooling it to room temperature, heating it at the Ac1-Ac3 transformation point, and then slowly cooling it.

[0010] However, the guaranteed PWHT time obtained by the above technique is extremely insufficient when the steel material is thickened and the weld conditions are severe, and there is a problem in that it is impossible to apply PWHT for a longer period of time. [Prior art documents] [Patent documents]

[0011] [Patent Document 1] Korean Patent Publication No. 10-2012-0075246 [Patent Document 2] Japanese Patent Application Publication No. 9-256037 Summary of the Invention [Problem to be solved by the invention]

[0012] According to one aspect of the present invention, there is provided an extra-thick steel material having excellent low-temperature impact toughness after a long-term PWHT even when the steel plate is thick, and a manufacturing method thereof.

[0013] The subject of the present invention is not limited to the above-mentioned contents. A person skilled in the art will have no difficulty in understanding further subjects of the present invention from the overall content of this specification. [Means for solving the problem]

[0014] One aspect of the present invention is a composition comprising, by weight percent, 0.10 to 0.25% carbon (C), 0.05 to 0.50% silicon (Si), 1.0 to 2.0% manganese (Mn), 0.005 to 0.1% aluminum (Al), 0.010% or less phosphorus (P), 0.0015% or less sulfur (S), 0.001 to 0.03% niobium (Nb), 0.001 to 0.03% vanadium (V), 0.001 to 0.03% titanium (Ti), 0.01 to 0.20% Cr, 0.01 to 0.15% molybdenum (Mo), 0.01 to 0.50% copper (Cu), 0.05 to 0.50% nickel (Ni), 0.0005 to 0.0040% calcium (Ca), with the balance being Fe and unavoidable impurities; The microstructure of the central part in the range of t / 4 to t / 2 (where t means the thickness of the steel plate) is composed of 35 to 40% ferrite and the remainder bainite complex structure in terms of area percent, the bainite packet size is 10 μm or less, and the porosity of the central part is 0.1 mm 3 / g or less, The depth of the surface crack is 0.5 mm or less, We can provide steel products with a cross-sectional hardness of the center of 200HB or less.

[0015] The average size of the prior austenite grains of the steel material may be 20 μm or less.

[0016] The thickness of the steel material may be 133 to 250 mm.

[0017] The above steel material can have a tensile strength of 450 to 650 MPa after PWHT, and a low-temperature impact toughness of 80 J or more at the center at -60°C.

[0018] Another aspect of the present invention is an alloy having, by weight percent, 0.10 to 0.25% carbon (C), 0.05 to 0.50% silicon (Si), 1.0 to 2.0% manganese (Mn), 0.005 to 0.1% aluminum (Al), 0.010% or less phosphorus (P), 0.0015% or less sulfur (S), 0.001 to 0.03% niobium (Nb), 0.001 to 0.03% vanadium (V), 0.001 to 0.03% titanium (Ti), 0.001 to 0.03% chromium (Cr), a step of obtaining a primary intermediate material by primarily heating a steel slab having a thickness of 650 to 750 mm, the steel slab being composed of 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, the balance being Fe and unavoidable impurities, in a temperature range of 1100 to 1300°C, and then performing a primary forging process with a cumulative reduction of 3 to 15% and a deformation speed of 1 to 4 / s; a step of obtaining a second intermediate material by subjecting the first intermediate material to a second heating at a temperature range of 1000 to 1500°C and then subjecting the first intermediate material to a second forging process at a cumulative reduction of 3 to 30% and a deformation speed of 1 to 4 / s; The third heating step heats the secondary intermediate material to a temperature range of 1000~1200℃; hot rolling the tertiarily heated secondary intermediate material at a finish hot rolling temperature of 900 to 1100°C to obtain a hot rolled material; cooling the hot-rolled material; quenching the cooled hot-rolled material by heating it to a temperature range of 820 to 900° C. and maintaining it for 10 to 40 minutes, and then cooling it at a cooling rate of 5° C. / s or more; and A method for manufacturing a steel material may be provided, which includes a tempering step of maintaining the quenched steel material at 600 to 680° C. for 10 to 40 minutes.

[0019] The cooling step may comprise cooling the hot-rolled material to a temperature range of Bs+20 to Ar1+20° C. at a cooling rate of 3° C. / s or more.

[0020] The method may further include cooling the hot-rolled material to a cooling end temperature and then air-cooling the material to room temperature.

[0021] The thickness of the primary intermediate material may be 450 to 550 mm.

[0022] The thickness of the secondary intermediate material may be 300 to 340 mm.

[0023] The hot-rolled material may have a thickness of 133 to 250 mm. Effect of the Invention

[0024] According to one aspect of the present invention, it is possible to provide an extra-thick steel material having excellent low-temperature impact toughness after a long-term PWHT even when the steel plate is thick, and a manufacturing method thereof.

[0025] According to another aspect of the present invention, there is provided a steel material that can be used in petrochemical manufacturing facilities, storage tanks, etc., and a manufacturing method thereof. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0026] Hereinafter, preferred embodiments of the present invention will be described. The embodiments of the present invention can be modified in various forms, and the scope of the present invention should not be construed as being limited to the embodiments described below. The embodiments are provided to further explain the present invention in detail to those skilled in the art to which the present invention pertains.

[0027] The present invention will be described in detail below.

[0028] The steel composition of the present invention will be described in detail below.

[0029] In the present invention, unless otherwise specified, % and ppm representing the content of each element are based on weight.

[0030] A steel material according to one aspect of the present invention may consist, by weight, of 0.10 to 0.25% carbon (C), 0.05 to 0.50% silicon (Si), 1.0 to 2.0% manganese (Mn), 0.005 to 0.1% aluminum (Al), 0.010% or less phosphorus (P), 0.0015% or less sulfur (S), 0.001 to 0.03% niobium (Nb), 0.001 to 0.03% vanadium (V), 0.001 to 0.03% titanium (Ti), 0.01 to 0.20% chromium (Cr), 0.01 to 0.15% molybdenum (Mo), 0.01 to 0.50% copper (Cu), 0.05 to 0.50% nickel (Ni), 0.0005 to 0.0040% calcium (Ca), with the balance being Fe and unavoidable impurities.

[0031] Carbon (C): 0.10~0.25% Carbon (C) is the most important element for ensuring the strength of steel, so it must be contained in steel within an appropriate range, and 0.10% or more must be added to obtain this added effect. On the other hand, if the content exceeds a certain level, the martensite fraction increases during quenching, which can cause the base material strength and hardness to increase excessively, which can cause surface cracks during forging and reduce the low-temperature impact toughness of the final product, so the upper limit is limited to 0.25%.

[0032] Therefore, the carbon (C) content may be 0.10 to 0.25%, and a more preferable upper limit may be 0.20%.

[0033] Silicon (Si): 0.05 to 0.50% Silicon (Si) is a substitutional element that improves the strength of steel by solid solution strengthening and has a strong deoxidizing effect, making it an essential element for manufacturing clean steel. In order to obtain the above effects, 0.05% or more must be added, and more preferably 0.20% or more can be added. On the other hand, if the content exceeds 0.5%, it may generate MA phase, excessively increase the strength of the ferrite matrix, and cause deterioration of the surface quality of extremely thick products.

[0034] Therefore, the silicon (Si) content may be 0.05 to 0.50%, more preferably the upper limit may be 0.40%, and more preferably the lower limit may be 0.20%.

[0035] Manganese (Mn): 1.0-2.0% Manganese (Mn) is a useful element that improves strength by solid solution strengthening and improves hardening ability so that a low-temperature transformation phase is generated. Therefore, in order to ensure a tensile strength of 450 MPa or more, it is preferable to add 1.0% or more of manganese (Mn). A more preferable lower limit can be 1.1%. On the other hand, if the content of manganese (Mn) is excessive, it may form MnS, which is a non-metallic inclusion elongated together with S, and may reduce toughness. When tensioned in the thickness direction, this may act as a factor in reducing the elongation rate and may be a factor in rapidly reducing the low-temperature impact toughness of the center part, so the upper limit is limited to 2.0%, and more preferably 1.5%.

[0036] Therefore, the manganese (Mn) content may be 1.0 to 2.0%, more preferably the upper limit may be 1.5%, and more preferably the lower limit may be 1.1%.

[0037] Aluminum (Al): 0.005 to 0.1% Aluminum (Al), together with Si, is one of the powerful deoxidizing agents in the steelmaking process, and in order to obtain the above effect, it is preferable to add 0.005% or more, and a more preferable lower limit may be 0.01%. On the other hand, if the content of aluminum (Al) is excessive, the proportion of Al2O3 in the oxidized inclusions generated as a result of deoxidation increases excessively, causing the size of the inclusions to become coarse, making it difficult to remove the inclusions during refining, and may be a factor in reducing the impact toughness properties, so the upper limit is set to 0.1%, and a more preferable upper limit may be 0.07%.

[0038] Therefore, the content of aluminum (Al) may be 0.005 to 0.1%, more preferably the upper limit may be 0.07%, and more preferably the lower limit may be 0.01%.

[0039] Phosphorus (P): 0.010% or less Phosphorus (P) is an element that forms coarse inclusions at grain boundaries, causing brittleness, and in order to improve resistance to brittle crack propagation, the upper limit of P is limited to 0.010% or less.

[0040] Therefore, the phosphorus (P) content can be 0.010% or less.

[0041] Sulfur (S): 0.0015% or less Sulfur (S) is an element that forms coarse inclusions at grain boundaries, causing brittleness, and the upper limit is limited to 0.0015% in order to improve resistance to brittle crack propagation.

[0042] Therefore, the sulfur (S) content can be 0.0015% or less.

[0043] Niobium (Nb): 0.001-0.03% Niobium (Nb) is an element that improves the strength of the base material by precipitating in the form of NbC or NbCN, and Nb dissolved during reheating at high temperature is very finely precipitated in the form of NbC during rolling, suppressing the recrystallization of austenite and refining the structure. In order to obtain the above-mentioned effect, it is preferable to add 0.001% or more of niobium (Nb), and a more preferable lower limit may be 0.005%. On the other hand, if the content is excessively added, undissolved niobium (Nb) is generated in the form of TiNb(C,N), which becomes a factor that inhibits the impact toughness characteristics, so the upper limit may be limited to 0.03%, and more preferably 0.02%.

[0044] Therefore, the niobium (Nb) content may be 0.001 to 0.03%. A more preferable upper limit may be 0.02%, and a more preferable lower limit may be 0.005%.

[0045] Vanadium (V): 0.001 to 0.03% Vanadium (V) is almost completely redissolved during reheating, so the strengthening effect due to precipitation or dissolution during the subsequent rolling is small, but it has the effect of improving strength by precipitating as very fine carbonitrides during the subsequent heat treatment process such as PWHT. In order to fully ensure the above-mentioned effects, it is necessary to add 0.001% or more of vanadium (V). More preferably, it can be contained in 0.01% or more. On the other hand, if the content is excessive, it can excessively increase the strength and hardness of the base material and welded parts, which can act as a cause of surface cracks during pressure vessel processing, and the manufacturing cost increases rapidly, which is commercially disadvantageous, so the upper limit can be set to 0.03%, more preferably 0.02%.

[0046] Therefore, the vanadium (V) content may be 0.001 to 0.03%, with a more preferable upper limit being 0.02% and a more preferable lower limit being 0.01%.

[0047] Titanium (Ti): 0.001 to 0.03% Titanium (Ti) is an element that precipitates as TiN during reheating, suppresses grain growth in the base material and the welded heat affected zone, and greatly improves low-temperature toughness, and is preferably added at 0.001% or more to obtain the above effect. On the other hand, if titanium (Ti) is excessive, it may cause clogging of the continuous casting nozzle or reduce low-temperature impact toughness due to central crystallization, and may combine with N to form coarse TiN precipitates in the thickness center, reducing the elongation rate of the product, thereby reducing the lamellar tearing resistance of the final material, so the upper limit is limited to 0.03%, more preferably 0.025%, and more preferably 0.018%.

[0048] Therefore, the titanium (Ti) content can be 0.001 to 0.03%, and the more preferable upper limit can be 0.025%, and more preferably 0.018%.

[0049] Chromium (Cr): 0.01-0.20% Chromium (Cr) has the effect of increasing the hardenability and forming a low-temperature transformation structure, thereby increasing the yield and tensile strength, and slowing down the decomposition rate of cementite during tempering after quenching or heat treatment after welding, thereby preventing a decrease in strength. In order to obtain the above-mentioned effect, the lower limit of the content can be limited to 0.01%. On the other hand, if the chromium (Cr) content is excessive, the size and fraction of Cr-rich coarse carbides such as M23C6 increase, the impact toughness of the product decreases, and the solid solubility of Nb in the product and the fraction of fine precipitates such as NbC decrease, which may decrease the strength of the product. Therefore, the upper limit can be set to 0.20%, more preferably 0.15%.

[0050] Therefore, the chromium (Cr) content may be 0.01 to 0.20%, and a more preferable upper limit may be 0.15%.

[0051] Molybdenum (Mo): 0.01-0.15% Molybdenum (Mo) is an element that increases grain boundary strength and has a large effect of solid solution strengthening in ferrite, and is an element that effectively contributes to increasing the strength and ductility of products. Furthermore, molybdenum (Mo) has the effect of preventing a decrease in toughness due to grain boundary segregation of impurities such as P. In order to obtain the above-mentioned effect, it is preferable to add 0.01% or more. On the other hand, molybdenum (Mo) is an expensive element, and if added in excess, there is a possibility that the manufacturing cost will increase significantly, so the upper limit can be limited to 0.15%.

[0052] Therefore, the molybdenum (Mo) content may be 0.01 to 0.15%, more preferably 0.05% as the lower limit and 0.12% as the upper limit.

[0053] Copper (Cu): 0.01 to 0.50% Copper (Cu) is an advantageous element in the present invention since it not only can greatly improve the strength of the matrix phase by solid solution strengthening in ferrite, but also has the effect of suppressing corrosion in a wet hydrogen sulfide atmosphere. In order to obtain such an effect, it can be added at 0.01% or more, and more preferably at 0.03% or more. On the other hand, if the content of copper (Cu) is excessive, it may cause star cracks on the surface of the steel sheet, and as an expensive element, there is a problem that the manufacturing cost increases significantly, so the upper limit can be limited to 0.50%, and preferably 0.30%.

[0054] Therefore, the copper (Cu) content may be 0.01 to 0.50%, more preferably 0.30% as the upper limit, and 0.03% as the lower limit.

[0055] Nickel (Ni): 0.05-0.50% Nickel (Ni) is an important element for increasing stacking faults at low temperatures, facilitating potential cross slip, improving impact toughness, and improving hardening ability to improve strength. In order to obtain the above effects, it is preferable to add 0.05% or more, more preferably 0.10% or more. On the other hand, if the content is excessive, the manufacturing cost may increase due to the high cost, so the upper limit may be limited to 0.50%, more preferably 0.30%.

[0056] Therefore, the nickel (Ni) content may be 0.05 to 0.50%, more preferably 0.30% as the upper limit, and 0.10% as the lower limit.

[0057] Calcium (Ca): 0.0005-0.0040% Adding calcium (Ca) after deoxidation with Al has the effect of suppressing the generation of MnS by combining with S, and suppressing the generation of cracks due to hydrogen-induced cracks by forming spherical CaS. In order to form sufficient CaS from the S contained as an impurity, it is preferable to add 0.0005% or more. On the other hand, if the content is excessive, the Ca remaining after forming CaS combines with O to form coarse oxidized inclusions, which are stretched and broken during rolling, resulting in a problem of reduced low-temperature impact toughness properties, so the upper limit can be limited to 0.0040%.

[0058] Therefore, the calcium (Ca) content may be 0.0005 to 0.0040%. A more preferable lower limit may be 0.0015%, and a more preferable upper limit may be 0.003%.

[0059] The steel material of the present invention may be composed of the above-mentioned composition, and the remaining iron (Fe) and inevitable impurities. The inevitable impurities cannot be excluded because they may be unintentionally mixed in during normal manufacturing processes. Since such impurities are known to any engineer in the field of normal steel manufacturing, the entire contents of the impurities will not be mentioned in this specification.

[0060] In the following, the steel microstructure of the present invention will be described in detail.

[0061] In the present invention, unless otherwise specified, the percentage indicating the fraction of the microstructure is based on the area.

[0062] The microstructure of the central part of the steel material satisfying the alloy composition according to one aspect of the present invention, in the range of t / 4 to t / 2 (where t means the thickness of the steel sheet), is composed of 35 to 40% ferrite and the remainder bainite, in terms of area percent, and the packet size of the bainite may be 10 μm or less. 3 / g or less.

[0063] When structures other than 35-40% ferrite and the remainder bainite are formed, it is difficult to ensure the low-temperature impact toughness properties targeted in the present invention. In particular, when ferrite is less than 35%, the strength is excessively excessive and the low-temperature impact toughness of the center at -60°C cannot be adequately ensured, and when ferrite exceeds 40%, the strength decreases and the tensile strength value required in the present invention cannot be ensured.

[0064] The bainite packet size, when measured by EBSD, can determine the size of the crystal grains centered on the 15° high hardness grain boundary, and can be limited to 10 μm or less in consideration of low-temperature impact toughness at -60°C, and more preferably 8 μm or less. However, when considering the level at which crystal grains can be refined by rolling, the lower limit can be limited to 5 μm.

[0065] In order to secure the low-temperature impact toughness characteristic targeted in this invention, the porosity in the center of the steel must be 0.1 mm 3 / g or less, and 3 If it exceeds 1 / g, it may act as a crack initiation point and cause the product to break upon impact.

[0066] The average size of prior austenite grains in the steel according to one aspect of the present invention may be 20 μm or less.

[0067] Immediately after hot rolling, the grain size in the center of the steel is controlled to ensure an appropriate impact toughness absorption energy value at -60°C. If the average size of the prior austenite grains exceeds 20 μm, coarse ferrite is formed, and it is difficult to control the remaining bainite packet size.

[0068] The steel production method of the present invention will now be described in detail.

[0069] The steel according to one aspect of the present invention may be manufactured by subjecting a steel slab having the above-mentioned alloy composition to primary heating and primary forging, secondary heating and secondary forging, tertiary heating and hot rolling, and cooling.

[0070] First heating and first forging A steel slab satisfying the above-mentioned alloy composition can be heated in the temperature range of 1100 to 1300°C, and then subjected to primary forging with a cumulative reduction of 3 to 15% and a deformation rate of 1 to 4 / s to produce a primary intermediate material.

[0071] The Ti and Nb composite carbonitrides or TiNb(C,N) coarse crystals formed during casting are redissolved, and the austenite is heated to the recrystallization temperature or higher before the first forging and maintained to homogenize the structure. The forging end temperature is kept high enough to minimize surface cracks that may occur during the forging process, so the heating temperature can be in the range of 1100°C or higher. On the other hand, if the heating temperature is too high, problems may occur due to oxidation scale at high temperatures, and the manufacturing cost may increase excessively due to the increase in cost associated with heating and maintenance, so the upper limit can be limited to 1300°C. The thickness of the slab in the present invention may be 650 to 750 mm, and preferably 700 mm.

[0072] In the first forging, the slab is forged to a thickness of 450-550 mm at a temperature range of 1100-1300°C, which is the primary heating temperature, and then processed to the desired width of the primary intermediate material. Since high deformation, low speed forging is essential to fully compress the voids, the forging speed can be limited to 1-4 / s.

[0073] If the cumulative reduction is less than 3%, the voids remaining in the slab cannot be sufficiently compressed, resulting in the generation of residual voids, which may result in a decrease in the lamellar tearing resistance of the product. The cumulative reduction of the first forging may be 5% or more, and more preferably 7% or more. However, if the cumulative reduction at or below the non-recrystallization temperature at which dislocation density is not recovered or offset by recrystallization exceeds 15%, the uniform elongation rate of the surface is significantly reduced due to work hardening of overlapping dislocations, and surface cracks may occur during the forging process. The cumulative reduction of the first forging may be 13% or less, and more preferably 11% or less.

[0074] Secondary heating and secondary forging The primary intermediate material is secondarily heated in the temperature range of 1000 to 1200° C., and then secondarily forged at a cumulative reduction of 3 to 30% and a deformation speed of 1 to 4 / s to produce a secondary intermediate material.

[0075] This is the stage where the primary intermediate material is heated at a temperature range of 1000 to 1200°C and forged to the desired thickness and length of the secondary intermediate material. As with the primary forging, the central porosity of the secondary intermediate material is reduced to 0.1 mm 3 In order to ensure that the thickness is equal to or less than 1 / g, high deformation, low speed forging is also required in the secondary forging. The thickness of the secondary intermediate material in the present invention may be 300 to 340 mm.

[0076] If the cumulative reduction in the second forging is less than 3%, the microvoids remaining after the first forging cannot be completely closed, and when deformation is applied to the ends of the oval closed voids, the notch effect may cause the physical properties to deteriorate more than when the voids are circular, so it is necessary to close the voids sufficiently with a deformation of 3% or more. However, if the cumulative reduction exceeds 30%, surface cracks may occur due to surface work hardening.

[0077] The deformation speed of the secondary forging can be 1 to 4 / s, similar to the primary forging. At a speed of less than 1 / s, there is a possibility that surface cracks may occur due to the temperature drop during finish forging, and a high deformation speed of more than 4 / s in the unrecrystallized region may also cause a decrease in elongation rate and surface cracks.

[0078] Tertiary heating The secondary intermediate material can be heated in the temperature range of 1000 to 1200°C.

[0079] The Ti and Nb complex carbonitrides or TiNb(C,N) coarse crystals formed during casting are redissolved, and the austenite is heated to and maintained at the recrystallization temperature before hot rolling to homogenize the structure, and the tertiary heating can be performed at a temperature of 1000℃ or higher to ensure that the rolling end temperature is high enough to minimize the crushing of inclusions during the rolling process. However, if the heating temperature is too high, problems may occur due to oxidation scale at high temperatures, and the manufacturing cost may increase excessively due to the increase in costs associated with heating and maintenance, so the upper limit of the temperature can be limited to 1200℃.

[0080] Hot rolling The tertiarily heated secondary intermediate material may be hot-rolled at a finish hot rolling temperature of 900 to 1100° C. to produce a hot-rolled material, and the thickness of the hot-rolled material may be 133 to 233 mm.

[0081] If the finish hot rolling temperature is less than 900°C, the deformation resistance increases excessively with decreasing temperature, making it difficult to sufficiently refine the austenite grains in the center of the product thickness direction, which may result in poor low-temperature impact toughness in the center of the final product. On the other hand, if the temperature exceeds 1100°C, the austenite grains may become excessively coarse, resulting in poor strength and impact toughness.

[0082] cooling The hot-rolled material produced as described above can be cooled to a temperature range of Bs+20 to Ar1+20° C. at a cooling rate of 3° C. / s or more.

[0083] After the hot rolling is completed, an accelerated cooling process at a cooling rate of 3°C / s or more is required to obtain a fine ferrite and pearlite composite structure transformed at a low temperature. If the cooling rate is less than 3°C / s, ferrite transformation begins during the cooling process, making it difficult to ensure the fine ferrite structure of the hot rolled material required in the present invention. In addition, if the cooling end temperature exceeds Ar1+20°C, ferrite grows after nucleation at high temperatures, making it difficult to refine it, and if the temperature is less than Bs+20°C, the hot rolled material structure is transformed into bainite or martensite, and additional grain refinement may not be achieved due to the Austenite Memory Effect during the heating process during quenching. The cooling conditions from cooling to the cooling end temperature to room temperature are not particularly limited, but air cooling can be applied in the present invention.

[0084] Quenching and tempering The hot-rolled material can be heated to a temperature range of 820 to 900°C and maintained at that temperature for 10 to 40 minutes, then quenched by cooling at a cooling rate of 5°C / s or more, and then tempered by maintaining at 600 to 680°C for 10 to 40 minutes.

[0085] If the temperature during quenching is less than 820℃ or the holding time is less than 10 minutes, the redissolution of carbides generated during cooling after rolling and impurity elements segregated at grain boundaries cannot be smoothly performed, and the low-temperature impact toughness of the center of the steel after heat treatment may be significantly reduced. On the other hand, if the temperature exceeds 900℃ or the holding time exceeds 40 minutes, the lamellar tearing resistance may be reduced due to the coarsening of austenite and precipitate phases such as Nb(C,N) and V(C,N).

[0086] If the tempering temperature is less than 600°C, the impingement carbon is not properly precipitated, and the strength increases excessively, making it difficult to secure the low-temperature impact toughness properties targeted in the present invention. If the tempering temperature exceeds 680°C, the dislocation density of the matrix decreases, and the cementite becomes excessively spheroidized and coarsened, making it difficult to secure appropriate strength.

[0087] Post-weld heat treatment (PWHT) After welding the steel material quenched and tempered in the present invention, a post-weld heat treatment can be performed. The conditions of the post-weld heat treatment are not particularly limited, and the heat treatment can be performed under normal conditions.

[0088] The steel material of the present invention manufactured as described above can have a thickness of 133 to 250 mm, a cross-sectional hardness of the center portion of the steel material of 200HB or less, a tensile strength of 450 to 620MPa after PWHT heat treatment of the steel material, and a low-temperature impact toughness of 80J or more at -60°C at the center portion of the steel material, and the steel material can have excellent low-temperature impact toughness properties without cracks occurring on the surface of the steel material.

[0089] The present invention will be described in more detail below with reference to 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. EXAMPLES

[0090] A 700 mm thick cast slab having the alloy composition shown in Table 1 was produced. Primary forging, secondary forging, hot rolling, cooling, and QT heat treatment were performed according to the process conditions shown in Table 2. A primary heating temperature of 1200°C, a secondary heating temperature of 1100°C, and a tertiary heating temperature of 1050°C were commonly applied, and a quenching and tempering time of 30 minutes was commonly applied. The thickness of the primary intermediate material was 550 mm, and the thickness of the secondary intermediate material was 400 mm. The cooling end temperature after hot rolling and the cooling rate during quenching, which are not disclosed in Table 2, were applied under conditions that satisfied the range of the present invention.

[0091] [Table 1] *Rank is ppm

[0092] [Table 2]

[0093] The microstructure and mechanical properties of the steel material manufactured above were measured. The microstructure fraction was measured using a scanning electron microscope, and the microstructure test piece was etched with Lepera and an optical image was taken, and the microstructure fraction was measured using an automatic image analyzer. At this time, the microstructure and porosity of the center part in the range of t / 4 to t / 2 (where t means the thickness of the steel plate) were measured. The uniform elongation of the slab surface layer was the elongation value measured at the maximum tensile stress part after making a tensile test piece on the slab surface layer in the first forging temperature range and conducting a tensile test. The bainite packet size was determined by determining the grain size centered on the grain boundary surface with a high inclination angle of 15° using EBSD, and the cross-sectional surface hardness was measured by measuring the reference cross-sectional hardness of the center part of the test piece using a Brinell hardness tester.

[0094] In addition, the mechanical properties shown in Table 4 below are tensile strength after PWHT and low-temperature impact toughness at -60°C. After visually observing the surface of the steel material, grinding was performed at the point where a surface crack was formed, and the grinding depth until the crack disappeared was measured as the surface crack depth.

[0095] [Table 3] F: Ferrite, B: Bainite, FM: Fresh martensite

[0096] [Table 4]

[0097] As shown in Table 3, it can be confirmed that the examples of the invention that satisfy the alloy composition and manufacturing method proposed in the present invention satisfy all of the mechanical properties targeted by the present invention.

[0098] On the other hand, in Comparative Examples 1 and 2, the cumulative reduction amount and deformation rate in the first forging exceeded the range of the present invention, and the slab surface uniform elongation rate in the forging temperature range did not satisfy the range of the present invention, and cracks occurred on the surface of the steel material.

[0099] In Comparative Example 3, the deformation rate during the secondary forging did not reach the range of the present invention, and there was an excessive amount of void in the center of the steel material, so that the low-temperature impact toughness did not satisfy the range proposed by the present invention.

[0100] In Comparative Example 4, the finish hot rolling temperature exceeded the range of the present invention, so that the average size of the prior austenite crystal grains became excessive, and the bainite packet size after quenching and tempering became coarse, resulting in deterioration of the low-temperature impact toughness value.

[0101] In Comparative Examples 5 and 6, the heating temperatures during quenching and tempering did not reach the range of the present invention. In Comparative Example 5, fresh martensite was formed, resulting in excessive hardness, and in Comparative Example 6, the hardness of bainite was excessive, resulting in an excessive increase in cross-sectional hardness at the center.

[0102] In the case of Comparative Example 7, the C content exceeded the range of the present invention, and bainite was formed excessively, which resulted in an excessive increase in tensile strength, a decrease in low-temperature impact toughness, and the occurrence of cracks.

[0103] In the case of Comparative Example 8, the Mn content did not satisfy the range of the present invention, and ferrite was excessively formed, so that sufficient tensile strength could not be ensured.

[0104] Although the present invention has been described in detail with reference to the above exemplary embodiments, other embodiments may be possible, and therefore the spirit and scope of the appended claims should not be limited to the exemplary embodiments.

Claims

1. In weight percent, it is composed of carbon (C): 0.10 to 0.25%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.0 to 2.0%, aluminum (Al): 0.005 to 0.1%, phosphorus (P): 0.010% or less, sulfur (S): 0.0015% or less, niobium (Nb): 0.001 to 0.03%, vanadium (V): 0.001 to 0.03%, titanium (Ti): 0.001 to 0.03%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and the balance being Fe and unavoidable impurities, The microstructure in the center part in the range of t / 4 to t / 2 (where t means the thickness of the steel plate) is composed of 35 to 40% ferrite and the remainder bainite composite structure in terms of area percent, the average bainite packet size is 10 μm or less, and the porosity in the center part is 0.1 mm ３ / g or less, The depth of the surface crack is 0.5 mm or less, The cross-sectional hardness of the center portion is 200 HB or less, A steel material having a thickness of 133 to 250 mm.

2. 2. The steel material according to claim 1, wherein the average size of the prior austenite grains of the steel material is 20 μm or less.

3. 2. The steel material according to claim 1, characterized in that the steel material has a tensile strength of 450 to 650 MPa after PWHT and a low-temperature impact toughness of 80 J or more at the center at -60°C.

4. In weight percent, carbon (C): 0.10-0.25%, silicon (Si): 0.05-0.50%, manganese (Mn): 1.0-2.0%, aluminum (Al): 0.005-0.1%, phosphorus (P): 0.010% or less, sulfur (S): 0.0015% or less, niobium (Nb): 0.001-0.03%, vanadium (V): 0.001-0.03%, titanium (Ti): 0.001-0.03%, chromium (Cr): 0.01-0 20%, Mo: 0.01-0.15%, Copper (Cu): 0.01-0.50%, Nickel (Ni): 0.05-0.50%, Calcium (Ca): 0.0005-0.0040%, the balance being Fe and unavoidable impurities. A steel slab having a thickness of 650-750 mm is primarily heated in a temperature range of 1100-1300°C, and then subjected to a primary forging process with a cumulative reduction of 3-15% and a deformation speed of 1-4 / s to obtain a primary intermediate material; The first intermediate material is secondarily heated at a temperature range of 1000 to 1500° C., and then secondarily forged at a cumulative reduction of 3 to 30% and a deformation speed of 1 to 4 / s to obtain a second intermediate material; a third heating step of heating the second intermediate material at a temperature range of 1000 to 1200° C.; hot rolling the tertiarily heated secondary intermediate material at a finish hot rolling temperature of 900 to 1100° C. to obtain a hot rolled material; Cooling the hot-rolled material to a temperature range of Bs+20 to Ar1+20° C. at a cooling rate of 3° C. / s or more, and then air-cooling the hot-rolled material to room temperature; quenching the cooled hot-rolled material by heating it to a temperature range of 820 to 900 ° C. and maintaining it for 10 to 40 minutes, and then cooling it at a cooling rate of 5 ° C. / s or more; and maintaining the quenched steel at 600 to 680° C. for 10 to 40 minutes for tempering, and the steel has a thickness of 133 to 250 mm.

5. The method for manufacturing a steel material according to claim 4, characterized in that the thickness of the primary intermediate material is 450 to 550 mm.

6. The method for manufacturing a steel material according to claim 4, characterized in that the thickness of the secondary intermediate material is 300 to 340 mm.

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