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

Abstract

An object of the present invention is to provide an extra-heavy steel material for flanges having excellent strength and low-temperature impact toughness, and a method for producing the same. The extra-thick steel material for flanges of the present invention, which has excellent strength and low-temperature impact toughness, contains, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0. 0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, and the remainder being Fe and other unavoidable impurities. The prior austenite grain size is 35 μm or less, and it has a fine structure containing 90 area % or more of one or more of bainite and martensite, and residual ferrite or pearlite.

Description

[Technical field]

[0001] The present invention relates to a steel material that can be used for wind power towers and systems, and a manufacturing method thereof, and more particularly to an extra-thick steel material for flanges having excellent strength and low-temperature impact toughness, and a manufacturing method thereof. [Background technology]

[0002] Wind turbines are in the limelight as an environmentally friendly means of generating electricity, and include components such as tower flanges, bearings, and main shafts. Among these, tower flanges are joint parts required for connecting towers, and usually 5 to 7 flanges are used for one tower. They are installed at sea or in extreme regions, so high durability is required. In particular, in response to the demand for large-capacity energy production and high efficiency, the scale of wind towers is also increasing, and accordingly, the steel materials used are continuously required to have high strength, high toughness, and thickness. As the thickness of the material increases, the total deformation amount decreases, and the microstructure becomes larger, and the material tends to deteriorate due to defects in the material such as inclusions and segregation. Therefore, in order to improve the external soundness of steel, there is a trend 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 to the utmost.

[0003] In particular, in the case of extremely thick materials with a thickness of more than 200 mmt, the amount of deformation at the center of the material is not large, so if unsolidified shrinkage holes that occur during continuous casting or casting are not sufficiently compressed in the forging process, they will remain in the form of residual voids in the center of the flange.

[0004] Such residual voids can act as starting points for cracks when the structure is subjected to axial stress in thickness, 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 before piercing (hole forging) and ring forging (product forming) which have a small amount of deformation.

[0005] The related Patent Document 1 is a technology for applying a strong reduction in a rough rolling process of a thick plate. Specifically, it utilizes a technology for determining a limit reduction rate for each thickness at which plate bite occurs from a strong reduction rate for each pass set to be close to the design allowable value (load and torque) of the rolling mill, a technology for distributing the reduction rate by adjusting the index of the thickness ratio for each pass in order to secure the target thickness of the rough rolling mill, and a technology for correcting the reduction rate based on the limit reduction rate for each thickness so that plate bite does not occur, and provides a manufacturing method that can apply an average reduction rate of about 27.5% in the final three passes of rough rolling based on 80 mmt as a reference. However, in the case of the above rolling method, the average reduction rate of the entire product thickness is measured, and in the case of an extremely thick material with a maximum thickness of 200 mmt or more, it is technically difficult to apply a high deformation to the center where residual voids exist.

[0006] Another method for manufacturing an extremely thick product is to utilize a forging machine, which has a higher effective deformation per pass than a rolling mill. In Patent Document 2, a slab containing, by mass%, C: 0.08-0.20%, Si: 0.40% or less, Mn: 0.5-5.0%, P: 0.010% or less, S: 0.0050% or less, Cr: 3.0% or less, Ni: 0.1-5.0%, Al: 0.010-0.080%, N: 0.0070% or less, and O: 0.0025% or less, which satisfy the relationships of Formulas 1 and 2, and the remainder being Fe and unavoidable impurities, is hot forged with a cumulative reduction of 25% or more. The document clearly states that a thick, high-toughness, high-strength material can be manufactured by heating the material to between Ac3 point and 1200°C, hot rolling with a cumulative reduction of 40% or more, quenching it at a temperature of Ar3 point or higher to a low temperature of 350°C or lower or below Ar3 point, and tempering it at a temperature of 450 to 700°C. The material has a plate thickness of 100 mmt or more, a yield strength of 620 MPa or more, and an absorbed energy of 70 J or more when evaluated for low-temperature impact toughness at -40°C.

[0007] However, in the above manufacturing method, if the cumulative reduction is too high, surface defects may occur due to localized deformation concentration, and in particular, if surface or subsurface defects exist in the slab before forging, the defects may propagate during the forging process, causing further deterioration of the surface quality in the product after rolling. Also, if the forging reduction per pass is insufficient, it is difficult to sufficiently close the voids remaining in the center even if the cumulative reduction is high, and the rolling process is also not suitable for controlling the voids and structure in the center of an extremely thick material because the effective deformation of the center is smaller than the deformation of the surface.

[0008] On the other hand, Patent Document 3 discloses that a thick, high-strength steel plate of 100 mmt or more and having a yield strength of 620 MPa or more can be manufactured through the steps of heating a material provided with a predetermined alloy composition to 1200 to 1350°C, hot forging with a cumulative reduction of 25% or more, heating to a temperature between the Ac3 point and 1200°C and hot rolling with a cumulative reduction of 40% or more, reheating to a temperature between the Ac3 point and 1050°C, quenching at a temperature of the Ac3 point or more to a temperature of 350°C or less or the Ar3 point or less, whichever is lower, and tempering at a temperature of 450°C to 700°C.

[0009] However, in the case of the above-mentioned ultra-high strength steel plate, the carbon equivalent (Ceq) and hardenability index (DI) are high, making it vulnerable to surface cracks during casting, and in the case of flange steel materials manufactured by normalizing heat treatment, the process conditions cannot be easily applied. In addition, when the carbon equivalent (Ceq) and hardenability index (DI) are high, cracks are likely to occur on the surface of the cast slab due to the generation of surface hard structure during the secondary cooling process of steelmaking, and the cracks may propagate during the forging process, degrading the surface quality of the final product.

[0010] 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 and excellent surface quality of the flange steel. [Prior art documents] [Patent documents]

[0011] [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]

[0012] Therefore, an object of the present invention is to provide an extra-heavy steel material for flanges having excellent strength and low-temperature impact toughness, and a manufacturing method thereof.

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

[0014] The extra-thick steel material for flanges of the present invention is In weight percent, it contains C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, and the remainder is Fe and other unavoidable impurities. The prior austenite grain size is 35 μm or less, and the microstructure contains at least one of bainite and martensite at 90 area % or more and residual ferrite or pearlite, and the low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more, Deformation-induced NbC precipitates of 5 to 50 nm were formed at 1 μm 2 and 5 or less coarse precipitates of 100 nm or more per grain; and The porosity of the steel core, which is in the region of 3 / 8t to 5 / 8t from the surface in the thickness direction, is 0.05 mm 3 / g or less.

[0015] The above steel further contains Zr: 0.001 to 0.15%.

[0016] The steel material has a thickness of 200 to 500 mm.

[0017] The steel material has a tensile strength of 590 to 820 MPa, a yield strength of 440 MPa or more, and an absorbed energy value of 50 J or more in a -50°C Charpy impact test.

[0018] The maximum surface crack depth of the above steel material is 0.1 mm or less (including 0).

[0019] In addition, the method for producing an extra-thick steel material for flanges of the present invention includes the steps of: preparing a slab containing, by weight percent, 0.05-0.2% C, 0.05-0.5% Si, 1.0-2.0% Mn, 0.005-0.1% Al, 0.01% or less P, 0.015% or less S, 0.005-0.07% Nb, 0.001-0.3% V, 0.001-0.05% Ti, 0.01-0.3% Cr, 0.01-0.12% Mo, 0.01-0.6% Cu, 0.05-4.0% Ni, 0.0005-0.004% Ca, with the remainder being Fe and other unavoidable impurities, and heating the slab to a temperature range of 1100-1300°C; The heated slab is first upset at a forging ratio of 1.3 to 2.4, and then bloom forged at a forging ratio of 1.5 to 2.0; Reheating the bloom forged material to a temperature range of 1100 to 1300 ° C.; The reheated bloom material is secondarily upset at a forging ratio of 1.3 to 2.3, and then round forged at a forging ratio of 1.65 to 2.25; The round forged material is subjected to a third upsetting at a temperature below the recrystallization temperature defined by the following Relation 1, with a forging ratio of 2.0 to 2.8 so that the cumulative reduction is 10% or more; After the tertiary upset material is drilled, it is reheated to a temperature range of 1100 to 1300 ° C., and then ring forged with a forging ratio of 1.0 to 1.6; The method is characterized by including a step of heating the ring-forged material to a temperature range of 820 to 930°C as a temperature measurement standard at the center thereof and holding the temperature for 5 to 600 minutes, followed by air-cooling to room temperature, and then heating the material to 550 to 700°C and holding the temperature therefor.

[0020] [Equation 1] T nr (℃) = 887 + 464 × C + 890 × Ti + 363 × Al-357 × Si + (6445 × Nb-644 × Nb 1 / 2 )+(732×V-230×V 1 / 2 ) Here, C, Ti, Al, Si, Nb, and V are in weight percent.

[0021] The slab is manufactured using one of a continuous casting process, a semi-continuous casting process, and an ingot casting process.

[0022] If the size of the forging surface punched out during the above primary upsetting is initially 700mm x 1800mm, then it will be 1000~1200mm x 1800~2000mm.

[0023] In the case of the above bloom forging, if the size of the forged surface at the completion of forging is initially 1000-1200mm x 1800-2000mm, then it is 1450-1850mm x 2100-2500mm.

[0024] After the above-mentioned secondary upsetting and round forging are completed, the product size is 1450~1850Φ×1300~1700mm.

[0025] When the above tertiary upsetting is completed, the product size is 2300~2800Φ×400~800mm.

[0026] The maximum thickness of the flange made of the above steel material can be 200 to 500 mm, the inner diameter is 4000 to 7000 mm, and the outer diameter is 5000 to 8000 mm. Effect of the Invention

[0027] The present invention configured as described above optimizes the forging process, thereby crimping the void in the center of the steel material and improving the internal integrity of the final product, and can effectively provide an extra-thick steel material that can be used for flanges, which has excellent not only strength but also low-temperature impact toughness. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0028] The present invention relates to an extra-thick steel material for flanges having excellent strength and low-temperature impact toughness and a manufacturing method of the product, and the following describes preferred embodiments of the present invention. The embodiments of the present invention can be modified in various forms, and the scope of the present invention should not be interpreted as being limited to the embodiments described below. The embodiments are provided to allow those skilled in the art to understand the present invention in more detail.

[0029] The extra-thick steel material for flanges having excellent strength and low-temperature impact toughness according to the present invention will be described in more detail below.

[0030] The extra-heavy steel material for flanges of the present invention, which has excellent strength and low-temperature impact toughness, contains, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: The alloy contains 0.05-4.0% Ca, 0.0005-0.004%, and the remainder is Fe and other unavoidable impurities. The prior austenite grain size is 35 μm or less, and the alloy has a microstructure containing at least one of bainite and martensite at 90 area % or more and residual ferrite or pearlite. The low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more, and deformation-induced NbC precipitates of 5-50 nm are formed within a range of 1 μm. 2 The steel has 10 or more precipitates per square centimeter, and 5 or less coarse precipitates of 100 nm or more per square centimeter, and the porosity of the steel in the center, which is the region from the surface to the thickness direction of 3 / 8t to 5 / 8t, is 0.05 mm 3 / g or less.

[0031] The alloy composition of the present invention will be described in more detail below. Unless otherwise specified, % and ppm described in the alloy composition are based on weight.

[0032] ·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, and 0.05% or more of carbon (C) can be added to obtain such an effect. Preferably, 0.10% or more of carbon (C) can be added. On the other hand, if the carbon (C) content exceeds a certain level, the hardening ability increases excessively during QT heat treatment, which may exceed the strength and hardness of the base material excessively, and this may cause surface cracks during forging, resulting in a decrease in low-temperature impact toughness properties in the final product. Therefore, in the present invention, the carbon (C) content can be limited to 0.20%, and the more preferable upper limit of the carbon (C) content can be 0.18%.

[0033] Silicon (Si): 0.05-0.50% Silicon (Si) is a substitutional element that improves the strength of steel by solid solution strengthening and has a strong deoxidizing effect, so it is an essential element for manufacturing clean steel. Therefore, silicon (Si) can be added in an amount of 0.05% or more, and more preferably 0.20% or more. On the other hand, if a large amount of silicon (Si) is added, it may generate a MA (Martensite-Austenite) phase, excessively increase the strength of the matrix, and cause deterioration of the surface quality of extremely thick products, so the upper limit of the content can be limited to 0.50%. A more preferable upper limit of the silicon (Si) content can be 0.40%.

[0034] Manganese (Mn): 1.0-2.0% Manganese (Mn) is a useful element that improves strength through 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 590 MPa or more, it is preferable to add 1.0% or more of manganese (Mn). More preferably, the manganese (Mn) content is 1.1% or more. Meanwhile, manganese (Mn) forms MnS, which is a non-metallic inclusion elongated with sulfur (S), and reduces toughness and can act as an impact initiation point, which can be a factor in rapidly reducing the low-temperature impact toughness of a product. Therefore, it is preferable to control the manganese (Mn) content to 2.0% or less, and more preferably, the manganese (Mn) content is 1.5% or less.

[0035] Aluminum (Al): 0.005~0.1% Aluminum (Al) is one of the powerful deoxidizing agents in the steelmaking process along with silicon (Si), and in order to obtain this effect, it is preferable to add 0.005% or more. A more preferable lower limit of the aluminum (Al) content may be 0.01%. On the other hand, if the aluminum (Al) content 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, which makes it difficult to remove the inclusions during refining, and this may be a factor in reducing low-temperature impact toughness. Therefore, it is preferable to control the aluminum (Al) content to 0.1% or less. A more preferable aluminum (Al) content may be 0.07% or less.

[0036] Phosphorus (P): 0.010% or less (including 0%), Sulfur (S): 0.0015% or less (including 0%) Phosphorus (P) and sulfur (S) are elements that induce brittleness in grain boundaries or form coarse inclusions that induce brittleness, so in order to improve resistance to brittle crack propagation, it is preferable to limit phosphorus (P) to 0.010% or less and sulfur (S) to 0.0015% or less.

[0037] Niobium (Nb): 0.005-0.07% Niobium (Nb) is an element that precipitates in the form of NbC or NbCN to improve the strength of the base material. In addition, niobium (Nb) that is dissolved during high-temperature reheating precipitates very finely in the form of deformation induction below the recrystallization temperature during forging, suppressing the growth of austenite, and thus has the effect of refining the structure. Therefore, niobium (Nb) is preferably added at 0.005% or more, and a more preferable niobium (Nb) content may be 0.01% or more. On the other hand, when niobium (Nb) is added in excess, undissolved niobium (Nb) is generated in the form of TiNb (C, N), which becomes a factor that inhibits low-temperature impact toughness, so the upper limit of the niobium (Nb) content is preferably limited to 0.07%. A more preferable niobium (Nb) content may be 0.065% or less.

[0038] Vanadium (V): 0.001-0.3% Vanadium (V) almost completely redissolves during reheating, so the strengthening effect due to precipitation or dissolution during subsequent rolling is insufficient. However, in the case of very thick forged materials, the air cooling rate is very slow, so it precipitates as very fine carbonitrides during the cooling process or tempering heat treatment process, which has the effect of improving strength. In order to fully obtain such an effect, it is necessary to add 0.001% or more of vanadium (V). A more preferable lower limit of the vanadium (V) content may be 0.01%. On the other hand, if the content is excessive, the surface hardness of the slab increases excessively due to its high hardening ability, which may act as a cause of surface cracks during flange processing, and the manufacturing cost increases sharply, which is not commercially useful. Therefore, the vanadium (V) content may be limited to 0.3% or less. A more preferable vanadium (V) content may be 0.25% or less.

[0039] Titanium (Ti): 0.001-0.05% Titanium (Ti) is a component that precipitates as TiN during reheating, inhibits the growth of prior austenite grains at high temperatures, and greatly improves low-temperature toughness. In order to obtain such an effect, it is preferable to add 0.001% or more of titanium (Ti). On the other hand, if titanium (Ti) is added excessively, there is a possibility that the low-temperature toughness will decrease due to clogging of the continuous casting nozzle and crystallization in the center. In addition, titanium (Ti) combines with nitrogen (N) to form coarse TiN precipitates in the thickness center, which reduces the elongation of the product, and therefore may reduce the uniform elongation during the forging process and cause surface cracks. Therefore, the titanium (Ti) content may be 0.05% or less. A preferable titanium (Ti) content may be 0.03% or less, and a more preferable titanium (Ti) content may be 0.018% or less.

[0040] Chromium (Cr): 0.01-0.30% Chromium (Cr) has the effect of preventing a decrease in strength by slowing down the spheroidization rate of cementite, and is an ingredient that improves hardenability during the cooling process. For this effect, 0.01% or more of chromium (Cr) can be added. On the other hand, if the chromium (Cr) content is excessive, M 23 The size and fraction of Cr-rich coarse carbides such as C6 increases, reducing the impact toughness of the product, and the solid solubility of niobium (Nb) in the product and the fraction of fine precipitates such as NbC decrease, which may cause a problem of reduced strength of the product. Therefore, in the present invention, the upper limit of the chromium (Cr) content may be limited to 0.30%. The preferred upper limit of the chromium (Cr) content may be 0.25%.

[0041] Molybdenum (Mo): 0.01-0.12% Molybdenum (Mo) is an element that increases grain boundary strength, increases hardening ability, and dissolves in precipitates to improve strength, and is an element that effectively contributes to increasing the strength and ductility of products. Molybdenum (Mo) also has the effect of preventing a decrease in toughness caused by grain boundary segregation of impurity elements such as phosphorus (P). For this effect, 0.01% or more of molybdenum (Mo) can be added. However, since molybdenum (Mo) is an expensive element and excessive addition may significantly increase manufacturing costs, the upper limit of the molybdenum (Mo) content can be limited to 0.12%.

[0042] Copper (Cu): 0.01-0.60% Copper (Cu) is an element that can greatly improve the strength of the matrix phase by solid solution strengthening in ferrite. To achieve this effect, 0.01% or more of copper (Cu) may be contained. A more preferable copper (Cu) content may be 0.03% or more. However, if the copper (Cu) content is excessive, there is a high possibility that star cracks will be induced on the surface of the steel sheet, and since copper (Cu) is an expensive element, there is a possibility that problems will arise in that the manufacturing cost will increase significantly. Therefore, in the present invention, the upper limit of the copper (Cu) content may be limited to 0.60%. A preferable upper limit of the copper (Cu) content may be 0.5%.

[0043] Nickel (Ni): 0.05-4.00% Nickel (Ni) is an element that effectively contributes to improving strength by increasing stacking faults at low temperatures, facilitating cross slip of dislocations to improve impact toughness, improving hardenability, and improving the degree of solid solution hardening. To achieve this effect, 0.05% or more of nickel (Ni) can be added. A preferred nickel (Ni) content is 0.10% or more. Meanwhile, if nickel (Ni) is added excessively, the manufacturing cost may increase due to the high cost, so the upper limit of the nickel (Ni) content may be limited to 4.00%. A preferred upper limit of the nickel (Ni) content may be 3.5%.

[0044] Calcium (Ca): 0.0005-0.0040% Adding calcium (Ca) after deoxidation with aluminum (Al) has the effect of suppressing the formation of MnS by combining with sulfur (S) that forms MnS inclusions, and suppressing the occurrence of cracks due to hydrogen-induced cracking by forming spherical CaS. In order to sufficiently form the sulfur (S) contained as an impurity into CaS, it is preferable to add 0.0005% or more of calcium (Ca). However, if the amount added is excessive, the calcium (Ca) remaining after forming CaS combines with oxygen (O) to form coarse oxidized inclusions, which are stretched and broken during forging, and may be a factor in reducing the elongation and low-temperature impact toughness properties. Therefore, the upper limit of the calcium (Ca) content can be limited to 0.0040%.

[0045] Zirconium (Zr): 0.001-0.15% The steel material of the present invention may selectively contain Zr in the range of 0.001 to 0.15. Zirconium (Zr) is a strong carbide forming element and may exist in the form of ZrC, and may improve the strength of the matrix phase in the form of precipitation strengthening like VC and NbC. For this effect, 0.001% or more of zirconium (Zr) may be added. However, if it is added in excess, it may increase the manufacturing cost due to the high cost, so the upper limit of the Zr content may be 0.15%.

[0046] The extra-thick flange steel material and its product, which are excellent in strength and low-temperature impact toughness, according to the present invention, may contain the remaining Fe and other inevitable impurities in addition to the above-mentioned components. However, since unintended impurities may be inevitably mixed in from raw materials or the surrounding environment during normal manufacturing processes, it is not possible to completely exclude them. Since these impurities are known to anyone with ordinary knowledge in this technical field, the present specification does not specifically mention all of the contents thereof. Furthermore, the addition of further effective components in addition to the above-mentioned components is not completely excluded.

[0047] On the other hand, the extra heavy steel material of the present invention has a fine structure in which the prior austenite grain size is 35 μm or less and at least one of bainite and martensite accounts for 90 area % or more, with the remainder being ferrite or pearlite.

[0048] In the present invention, the prior austenite grain size in the microstructure of the final product is 35 μm or less. If the prior austenite grain size exceeds 35 μm, the length of the crack path during impact fracture becomes shorter, the DBTT (Ductile Brittle Transition Temperature) increases, and low-temperature impact toughness deteriorates. Therefore, the prior austenite grain size is preferably 35 μm or less.

[0049] The steel material of the present invention has a microstructure containing at least one of bainite and martensite at 90% by area or more, and the remainder ferrite or pearlite. If the phase fraction of low-temperature transformation phases such as bainite and martensite is less than 95% by area, the strength of the base phase decreases, and the steel material cannot meet the requirements of the present invention for tensile strength of 590 to 820 MPa and yield strength of 440 MPa or more.

[0050] In the present invention, the low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more. When the matrix base structure is bainite or martensite, cracks during impact testing propagate along packet boundaries based on high-angle grain boundaries, so if the packet size is large, DBTT increases and impact toughness may deteriorate. Therefore, in order to ensure an absorbed energy value of 50 J or more in a -50°C Charpy impact test required in the present invention, it is appropriate that the packet size is 15 μm or less.

[0051] In addition, the steel material of the present invention has deformation-induced NbC precipitates of 5 to 50 nm in size in the matrix structure. 2If the number of deformation-induced NbC precipitates of 5 to 50 nm is less than 10, the precipitation strengthening effect is weakened, and if the number of coarse precipitates of 100 nm or more is more than 5, the pinning effect and the precipitation strengthening effect are similarly lost, so that it is not easy to ensure the tensile strength of 590 to 820 MPa and the yield strength of 440 MPa or more required in the present invention.

[0052] Furthermore, the steel material of the present invention has a porosity of 0.05 mm in the center of the steel material, which is the region from the surface to the thickness direction of 3 / 8t to 5 / 8t. 3 / g or less.

[0053] Moreover, the extra heavy steel material of the present invention can have a thickness of 200 to 500 mm.

[0054] Moreover, the extra heavy-gauge steel material of the present invention can have a tensile strength of 590 to 820 MPa, a yield strength of 440 MPa or more, and an absorbed energy value in a Charpy impact test at -50°C of 50 J or more.

[0055] The maximum surface crack depth of the steel material may be 0.1 mm or less (including 0).

[0056] Next, a method for producing an extra-thick steel material for flanges having excellent strength and low-temperature impact toughness, which is another aspect of the present invention, will be described in detail.

[0057] The method for manufacturing an extra-thick steel material of the present invention includes the steps of preparing a slab having the above-mentioned compositional components, heating the slab to a temperature range of 1100 to 1300°C, performing a first upset of the heated slab with a forging ratio of 1.3 to 2.4, and then performing bloom forging at a forging ratio of 1.5 to 2.0, reheating the bloom forged material to a temperature range of 1100 to 1300°C, performing a second upset of the reheated bloom material with a forging ratio of 1.3 to 2.3, and then performing round forging at a forging ratio of 1.65 to 2.25, and then performing round forging of the round forged material. The method includes a step of tertiary upsetting at a temperature equal to or lower than the recrystallization temperature defined by the following relational expression 1 with a forging ratio of 2.0 to 2.8 so that the cumulative reduction is 10% or more; a step of reheating the tertiary upset material to a temperature range of 1100 to 1300°C after drilling holes, and then ring forging at a forging ratio of 1.0 to 1.6; a step of heating the ring forged material to a temperature range of 820 to 930°C as a temperature measurement standard for the center part and holding it for 5 to 600 minutes, followed by air cooling to room temperature, and then heating it to 550 to 700°C and holding it there.

[0058] Slab heating In the present invention, first, a slab having the above-mentioned compositional components is prepared, and then heated to a temperature range of 1100 to 1300°C.

[0059] The slab needs to be heated at a temperature above a certain range in order to redissolve titanium (Ti) and niobium (Nb) complex carbonitrides or TiNb (C, N) coarse crystals formed during casting. It is also preferable to heat the slab to a temperature above the recrystallization temperature before the first upset forging and hold it there to homogenize the structure, ensure a sufficiently high forging end temperature, and minimize surface cracks that may occur during the forging process. Therefore, it is preferable to heat the slab in the present invention at a temperature of 1100°C or higher.

[0060] On the other hand, if the slab heating temperature is too high, high-temperature oxide scale may be excessively generated, and the manufacturing cost may increase excessively due to high-temperature heating and holding. Therefore, it is preferable that the primary heating of the slab in the present invention is performed in the range of 1300°C or less.

[0061] Primary upset and bloom forging Next, in the present invention, the heated slab is subjected to a primary upset at a forging ratio of 1.3 to 2.4, and then to bloom forging at a forging ratio of 1.5 to 2.0.

[0062] Upsetting is a method of performing strong plastic deformation perpendicular to the longitudinal axis, and the forging ratio during the first upsetting is preferably 1.3 to 2.4, and more preferably 1.5 to 2.0. Here, the forging ratio refers to the ratio of the cross-sectional area changed by forging. In such a first upsetting, if the size of the forged surface to be punched out is initially 700 mm x 1800 mm, it may be 1000 to 1200 mm x 1800 to 2000 mm.

[0063] If the forging ratio during the first upset is less than 1.3, it is difficult to sufficiently compress the porosity remaining in the center of the slab. Therefore, the porosity required in the final product of the present invention is set at an appropriate level of 0.05 mm 3 / g or less, it is difficult to ensure low-temperature impact toughness in the center. On the other hand, if the forging ratio in the first upset exceeds 2.4, buckling occurs during the forging process, making it difficult to achieve the surface quality and appropriate shape control required for flange products. Therefore, the forging ratio in the first upset is preferably 1.3 to 2.4.

[0064] In the present invention, the primarily upset material is subjected to bloom forging at a forging ratio of 1.5 to 2.0.

[0065] Bloom forging is a method of processing a first upset material into a bloom shape by further pressing the material, and the area is expanded while processing both the upper and lower surfaces in a certain direction of width or length. In the case of bloom forging, if the size of the forged surface at the time of forging completion is initially 1000-1200mm x 1800-2000mm, it can be 1450-1850mm x 2100-2500mm. In the case of bloom forging, the forging ratio is preferably 1.5-2.0. The reason is that if the forging ratio is less than 1.5, it is difficult to ensure the appropriate porosity quality required in the present invention, as in the case of upset forging, and if it exceeds 2.0, surface cracks may occur.

[0066] The forging direction can be either the length direction or the width direction, but in the case of the length direction, the cast structure is more dense, so the elongation rate of the surface layer structure is high and the workability is excellent. Therefore, from the viewpoint of surface cracks, the bloom forging in the length direction may be more suitable than that in the width direction.

[0067] reheating In the present invention, the bloom forged material is reheated to a temperature range of 1100 to 1300°C.

[0068] When the bloom forging is completed, the bloom surface temperature is 950°C or less, and if processing is continued, surface cracks or material breakage may occur. Therefore, after bloom forging, the material can be heated again to a temperature range of 1100 to 1300°C. As described above, it is preferable to heat the material to 1100°C or more for reasons such as re-dissolving the crystallized material, homogenizing the structure, and preventing surface cracks, and it is preferable to control the temperature to 1300°C or less due to problems such as excessive scaling and coarsening of crystal grains.

[0069] Secondary Upset - Round Forging Next, in the present invention, the reheated bloom material is subjected to secondary upsetting with a forging ratio of 1.3 to 2.3, and then round forging with a forging ratio of 1.65 to 2.25.

[0070] That is, the bloom that has been heated is subjected to a secondary upsetting at a forging ratio of 1.3 to 2.3, and then round forging is performed at a forging ratio of 1.65 to 2.25 to process the bloom into a circular flange edge. When the secondary upsetting and round forging are completed, the product size can be 1450 to 1850Φ×1300 to 1700mm.

[0071] If the forging ratio during the secondary upsetting and round forging is less than the level required by the present invention, the porosity of the center of the final product is reduced to 0.05 mm 3 / g or less, it is difficult to ensure the low-temperature impact toughness of the center of the steel material. On the other hand, if the forging ratio exceeds the standard of the present invention, it is not possible to process it into the desired flange product shape due to problems such as buckling and surface cracks and shape defects.

[0072] Tertiary upset (formation of deformation-induced precipitates) Subsequently, in the present invention, the round forged material is tertiary upset at a temperature equal to or lower than the recrystallization temperature defined by the following Relational Formula 2 with a forging ratio of 2.0 to 2.8 so that the cumulative reduction amount is 10% or more.

[0073] The cylindrical material can be processed to an appropriate flange thickness by a tertiary upsetting before piercing. After the tertiary upsetting, the product size can be 2300~2800Φ×400~800mm.

[0074] The forging ratio of the tertiary upset may be 2.0 to 2.8. If the forging ratio is insufficient or excessive, problems such as the above-mentioned inability to control residual voids and surface cracks / shape control may occur.

[0075] In this tertiary upsetting process, the important thing is the cumulative reduction at a temperature below the recrystallization temperature (Rst) of the steel material, and forging and rolling are performed so that the cumulative reduction is 10% or more. At this time, the recrystallization temperature can be calculated by the following Relation 2.

[0076] [Equation 2] T nr (℃) = 887 + 464 × C + 890 × Ti + 363 × Al-357 × Si + (6445 × Nb-644 × Nb 1 / 2 )+(732×V-230×V 1 / 2 ) Here, C, Ti, Al, Si, Nb, and V are in weight percent.

[0077] If the cumulative reduction is less than 10% at a temperature below the recrystallization temperature, the formation of deformation-induced ultrafine NbC or NbCN precipitates of 5 to 50 nm is not easy, and the precipitates are 1 μm 2 There is a possibility that the number of precipitates per unit area is less than 10, or the number of coarse precipitates with a size of 100 nm or more exceeds 5. If the amount of fine precipitates is reduced or the size is increased, the precipitation strengthening effect is small, and the pinning effect is reduced during the subsequent reheating for quenching, making it difficult to ensure that the average prior austenite grain size in the center of the product is 35 μm or less. Therefore, it is preferable to control the cumulative reduction at a temperature below the recrystallization temperature to 10% or more, more preferably 15% or more, and most preferably 20% or more.

[0078] Hole drilling and ring forging Next, in the present invention, the tertiary upset material is subjected to hole machining, reheated to a temperature range of 1100 to 1300° C., and then ring forged with a forging ratio of 1.0 to 1.6.

[0079] After the above-mentioned tertiary upsetting is completed, a hole can be drilled in the center of the material using a 500 to 1000Φ punch.

[0080] The material with the holes machined is reheated again to the temperature range of 1100 to 1300°C, and then processed into the final flange ring shape. The maximum thickness of the flange made of the steel material can be 200 to 500 mm, the inner diameter can be 4000 to 7000 mm, and the outer diameter can be 5000 to 8000 mm. Since ring forging is a process in which the final shape and dimensional control are more important than gap pressing, severe plastic processing is not applied. Therefore, the forging ratio can be 1.0 to 1.6, and more preferably 1.2 to 1.4.

[0081] Meanwhile, in the present invention, the deformation rate in all the forging processes presented above can be 1 / s to 4 / s. At a deformation rate of less than 1 / s, the finish forging temperature drops, and there is a possibility that surface cracks may occur. In contrast, when a high deformation rate of more than 4 / s is applied in the unrecrystallized region, there is a possibility that surface cracks may be induced due to a drop in elongation caused by excessive local work hardening.

[0082] Quenching and tempering heat treatment Finally, in the present invention, the ring-forged material is heated to a temperature range of 820 to 930°C based on the temperature measurement standard of its center (t / 2), held for 5 to 600 minutes, air-cooled to room temperature, and then heated to 550 to 700°C and held there, thereby performing quenching and tempering heat treatment.

[0083] If the heating temperature is less than 820°C or the holding time is less than 5 minutes during quenching, the carbides generated during cooling after forging and the impurity elements segregated at the grain boundaries may not be smoothly redissolved, and the low-temperature toughness of the steel after heat treatment may be significantly reduced. On the other hand, if the heating temperature exceeds 930°C or the holding time exceeds 600 minutes, the prior austenite grain size may exceed 35 μm required in the present invention, or the precipitate phases such as Nb(C,N) and V(C,N) may become coarse, resulting in a deterioration in strength and low-temperature impact toughness.

[0084] On the other hand, the cooling rate can be 0.5°C / s to 30°C / s based on the center (t / 2) of the product. If the cooling rate is less than 0.5°C / s, the fraction of bainite or martensite, which is the low-temperature transformation structure required in the present invention, cannot be ensured to be 90% or more, making it difficult to ensure appropriate strength, and if it exceeds 30°C / s, the strength increases excessively and low-temperature impact toughness may deteriorate. Therefore, the cooling rate during quenching is preferably 0.5°C / s to 30°C / s.

[0085] Meanwhile, tempering can be performed for 5 to 600 minutes in a temperature range of 550 to 700°C. If the tempering temperature is below 550°C, carbon does not diffuse properly after quenching, resulting in excessive strength, which may result in deterioration of low-temperature impact toughness at -50°C. On the other hand, if the tempering temperature exceeds 700°C, fresh martensite formed during air cooling by heating in the two-phase region may also deteriorate low-temperature impact toughness. Therefore, the tempering temperature is preferably 550 to 700°C.

[0086] If the tempering holding time is less than 5 min, the dislocation density after quenching is not appropriately reduced, and carbon diffusion does not occur sufficiently, so that the strength is excessively high, and the low-temperature impact toughness is reduced, similar to the reason why the temperature is low. Also, if the holding time exceeds 600 min, carbon becomes excessively spheroidized and coarsened, and the impact toughness is deteriorated. Therefore, the appropriate holding time for the tempering heat treatment is preferably 5 to 600 min. EXAMPLES

[0087] The present invention will be described in 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.

[0088] [Table 1] *In Table 1, the content units of the component elements are in weight percent, but the units of P, S, and Ca are in ppm. The remaining components are Fe and unavoidable impurities.

[0089] A slab having a thickness of 700 mm was manufactured having the alloy composition shown in Table 1 above. Using this slab, a flange having a final thickness of 320 mm was manufactured through the process conditions shown in Tables 2-3 below, including slab preparation, forging (reheating and first upset, bloom forging, reheating, second upset-round forging, third upset, reheating and ring forging), and quenching and tempering. At this time, after the bloom forging was completed, the reheating temperature for the second upset was 1230°C ± 10°C, and the forging ratio of the round forging after the second upset was 2.0, which was also applied. In addition, all process conditions that satisfy the scope of the present invention were applied to the processes other than those shown in Tables 2-3 below.

[0090] Thereafter, the physical properties of each of the test pieces produced above were measured and are shown in Table 4 below.

[0091] Here, the prior austenite grain size and the fraction of low-temperature transformation phases (bainite and / or martensite) were measured by taking a specimen from the central microstructure specimen after QT heat treatment and using an automatic image analyzer, and the packet size of bainite was automatically analyzed by setting the boundary condition to 15° through EBSD (Electron Back Scattered Diffraction). Meanwhile, in this example, the residual structure excluding the low-temperature transformation phase in both the inventive examples and the comparative examples is ferrite and / or pearlite.

[0092] The yield / tensile strength was evaluated by a room temperature tensile test according to ASTM A370, and a 0.2% offset was applied to the yield strength. The impact toughness of each test piece was calculated by using the average of the absorbed energy measured three times at the same temperature through the Charpy V-Notch test.

[0093] The number of deformation-induced NbC precipitates of 5 to 50 nm observed in the cross section of the steel material was measured using a TEM. NbC precipitates were confirmed by the NbC diffraction pattern and EDX mapping, and the precipitates were measured at 1 μm. 2 The number of NbC precipitates located at the positions were counted.

[0094] The porosity of the product center is calculated by the density (g / mm 3 ) and reciprocal (mm 3 It was measured by taking the mass (g).

[0095] In addition, after visually observing the surface of each test piece, the surface was ground at the point where a surface crack had formed, and the length ground until the crack disappeared was measured as the length of the surface crack.

[0096] [Table 2]

[0097] [Table 3]

[0098] [Table 4] *In Table 3, a* is the 1μm diameter of deformation-induced NbC precipitates of 5 to 50 nm observed in the cross section of the steel material. 2 b* indicates the number of coarse precipitates of 100 nm or more.

[0099] As can be seen from Tables 1 to 3 above, in the case of Examples 1 to 5, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, not only do they have excellent strength and excellent low-temperature impact toughness at -50°C, but they also have good surface quality in the flange product state.

[0100] In contrast, Comparative Examples 1 to 12 are cases in which the alloy composition proposed by the present invention is satisfied but the manufacturing conditions are not satisfied, and since the properties proposed by the present invention in the flange product state, such as the prior austenite grain size, the fraction and packet size of the low-temperature transformation phase, and the porosity, are not satisfied, it can be seen that the strength and low-temperature impact toughness are at a low level. Furthermore, even if the material quality is good, if the forging ratio conditions are not satisfied at each stage of forging, poor surface quality properties in the product state can be confirmed by the occurrence of surface cracks or through cracks.

[0101] On the other hand, Comparative Examples 13 to 16 satisfy the manufacturing conditions proposed by the present invention, but do not satisfy the alloy composition, and are found to have low quality levels, such as excessive strength (insufficient impact toughness) or insufficient strength.

[0102] As described above, the detailed description of the present invention has been given of the preferred embodiment of the present invention, but it goes without saying that a person having ordinary skill in the art to which the present invention pertains can make various modifications without departing from the scope of the present invention. Therefore, the scope of the present invention should not be limited to the described embodiment, but should be determined by the claims below as well as equivalents thereto.

Claims

1. The alloy contains, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, and the remainder being Fe and other unavoidable impurities, The prior austenite grain size is 35 μm or less, and the microstructure contains at least one of bainite and martensite at 90 area % or more and residual ferrite or pearlite, and the low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more, Deformation-induced NbC precipitates of 5 to 50 nm are reduced to 1 μm. ２ and 5 or less coarse precipitates of 100 nm or more per grain; and The porosity of the steel core, which is the area from the surface to the thickness direction, is 0.05 mm ３ / g or less.

2. The extra-thick steel material for flanges according to claim 1, characterized in that the steel material further contains, by weight percent, Zr: 0.001 to 0.15%.

3. 2. The extra-thick steel material for flanges according to claim 1, wherein the steel material has a thickness of 200 to 500 mm.

4. The extra-thick steel material for flanges according to claim 1, characterized in that the steel material has a tensile strength of 590 to 820 MPa, a yield strength of 440 MPa or more, and an absorbed energy value of 50 J or more in a -50 ° C. Charpy impact test.

5. 2. The extremely thick steel material for flanges according to claim 1, characterized in that the maximum surface crack depth of the steel material is 0.1 mm or less (including 0).

6. preparing a slab containing, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, with the remainder being Fe and other unavoidable impurities; and heating the slab to a temperature range of 1100-1300°C. performing a first upsetting of the heated slab at a forging ratio of 1.3 to 2.4, and then performing bloom forging at a forging ratio of 1.5 to 2.0; reheating the bloom forged material to a temperature range of 1100 to 1300°C; Secondarily upsetting the reheated bloom material at a forging ratio of 1.3 to 2.3, and then round forging at a forging ratio of 1.65 to 2.25; The round forged material is subjected to a third upsetting at a temperature equal to or lower than the recrystallization temperature defined by the following Relation 1, with a forging ratio of 2.0 to 2.8 so that the cumulative reduction is 10% or more; After the tertiary upset material is drilled, it is reheated to a temperature range of 1100 to 1300 ° C., and then ring forged with a forging ratio of 1.0 to 1.6; The method for manufacturing an extra-thick steel material for flanges includes a step of heating the ring-forged material to a temperature range of 820 to 930°C as a temperature measurement standard at the center of the material, holding the temperature range for 5 to 600 minutes, air-cooling the material to room temperature, and then heating the material to 550 to 700°C and holding the temperature range for 5 to 600 minutes. [Relationship 1] T ｎｒ (℃)=887+464×C+890×Ti+363×Al-357×Si+(6445×Nb-644×Nb １／２ )+(732×V-230×V １／２ ) (Here, C, Ti, Al, Si, Nb, and V are in weight percent.)

7. 7. The method of claim 6, wherein the slab is manufactured using one of a continuous casting process, a semi-continuous casting process, and an ingot casting process.

8. The method for manufacturing an extremely thick steel material for flanges according to claim 6, characterized in that the size of the forged surface punched during the first upsetting is 1000 to 1200 mm x 1800 to 2000 mm when the size is initially 700 mm x 1800 mm.

9. 7. The method for manufacturing an extremely thick steel material for flanges according to claim 6, characterized in that in the case of bloom forging, the size of the forged surface at the completion of forging is 1450-1850 mm x 2100-2500 mm when the initial size is 1000-1200 mm x 1800-2000 mm.

10. The manufacturing method of the extra thick steel material for flanges according to claim 6, characterized in that when the secondary upsetting and round forging are completed, the product size is 1450 to 1850Φ x 1300 to 1700 mm.

11. The method for manufacturing an extremely thick steel material for flanges according to claim 6, characterized in that when the tertiary upsetting is completed, the product size is 2300 to 2800Φ×400 to 800 mm.

12. The method for manufacturing an extremely thick steel material for flanges according to claim 6, characterized in that the maximum thickness of the flange made of the steel material can be 200 to 500 mm, the inner diameter is 4000 to 7000 mm, and the outer diameter is 5000 to 8000 mm.

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