Classifications

• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
  • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
  • B21J—FORGING; HAMMERING; PRESSING METAL; RIVETING; FORGE FURNACES
  • B21J1/00—Preparing metal stock or similar ancillary operations prior, during or post forging, e.g. heating or cooling
  • B21J1/04—Shaping in the rough solely by forging or pressing
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
  • C21D1/18—Hardening; Quenching with or without subsequent tempering
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
  • C21D1/18—Hardening; Quenching with or without subsequent tempering
  • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
  • C21D1/26—Methods of annealing
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D6/00—Heat treatment of ferrous alloys
  • C21D6/001—Heat treatment of ferrous alloys containing Ni
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D7/00—Modifying the physical properties of iron or steel by deformation
  • C21D7/13—Modifying the physical properties of iron or steel by deformation by hot working
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0226—Hot rolling
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
  • C21D8/0263—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/002—Bainite
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/005—Ferrite
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/008—Martensite
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/009—Pearlite

Definitions

• the present disclosure relates to a steel sheet and a manufacturing method therefor, and more particularly, to an extremely thick steel material for flange shaving excellent fatigue properties and low-temperature impact toughness, and a method of manufacturing the same.
• a wind turbine has gained attention as an eco-friendly means of generating electricity, and may include components such as a tower flange, a bearing, and a main shaft.
• tower flanges may be joint components required for connecting towers, usually 3 to 4 flanges may be used for one tower, and may be installed in the sea or regions of extreme conditions, such that high durability may be required.
• a wind tower has been increasingly used, and accordingly, a steel material used therefor has also continuously required to have high fatigue properties and low-temperature impact toughness. As a thickness of a material increases, the total deformation amount decreases, such that a microstructure may become large, and the material may tend to deteriorate due to defects in the material such as inclusions or segregation.
• the deformation amount of the central portion of the material may not be large, and accordingly, when a unsolidified shrinkage cavity occurring during continuous casting or casting is not sufficiently compressed during a forging process, the shrinkage cavity may remain in the form of retained voids in the central portion of the flange.
• retained voids may act as crack initiation points when stress in the thickness axial direction is applied, which may eventually cause damage to the entire facility in the form of lamellar tearing. Accordingly, a process of sufficiently compressing central voids such that retained voids are not present may be necessary.
• Cited document 1 related thereto is a technique for applying a strong reduction in a roughing process for a thick sheet.
• cited document 1 uses a technique for determining a limit reduction ratio for each thickness at which plate bite occurs from a reduction ratio per pass determined to be close to design allowance (load and torque) of a rolling mill, a technique for distributing the reduction ratio by adjusting the index of the thickness ratio per pass to secure the target thickness of the roughing mill, and a technique for modifying the reduction ratio such that plate bite does not occur based on the limit reduction ratio for each thickness, thereby providing a manufacturing method in which an average reduction ratio of approximately 27.5% may be applied in the rough-rolling final 3 passes with respect to a thickness of 80 mm.
• the average reduction ratio of the entire thickness of the product is measured, and in the case of extremely thick materials having a maximum thickness of 200 mm or more, there may be a technical difficulty in applying high strain to the central portion in which retained voids are present.
• Cited document 2 discloses that, by hot-forging a slab including, by mass%, C: 0.08 to 0.20%, Si: 0.40% or less, Mn: 0.5 to 5.0%, P: 0.010% or less, S: 0.0050% or less, Cr: 3.0% or less, Ni: 0.1 to 5.0%, Al: 0.010 to 0.080%, N: 0.0070% or less, O: 0.0025% or less with a balance of Fe and inevitable impurities and satisfying equations (1) and (2) with cumulative reduction of 25% or more, heating at the Ac3 point or higher and 1200°C or lower, hot-rolling with an cumulative reduction of 40% or more, rapid cooling from a temperature of Ar3 point or higher to 350°C or lower or a low temperature of Ar3 point or lower, and performing a tempering heat treatment process at a temperature of 450 to 700°C, a thick material with high toughness and strength
• the above manufacturing method may cause surface defects due to local strain concentration when the cumulative reduction is excessively high, and in particular, when a surface or subsurface defect is present in the cast steel before forging, the defect may propagate during the forging process, which may further deteriorate the surface quality of the product after rolling. Also, when the forging reduction per pass is insufficient, it may be difficult to sufficiently compress the voids remaining in the central portion even when the cumulative reduction is high, and because the effective deformation amount in the central portion may be small as compared to the surface deformation, the rolling process may also not suitable for controlling the voids and structure of the central portion of an extremely thick material.
• Cited document 3 discloses that, by heating a material provided with a predetermined alloy composition to 1200 to 1350°C, hot-forging with an cumulative reduction of 25% or more, heating at the Ac3 point or higher and 1200°C or lower, hot-rolling with an cumulative reduction of 40% or more, reheating at the Ac3 point or higher and 1050°C or lower, rapidly cooling from the temperature of Ac3 point or higher to 350°C or lower or the lower temperature of Ar3 point or lower, and tempering at a temperature of 450 to 700°C, a thick steel sheet with high strength having a thickness of 100 mm or more and a yield strength of 620 MPa or more may be manufactured.
• the carbon equivalent (Ceq) and hardenability index (DI) may be high, such that the steel sheet may be vulnerable to surface cracks during casting, and also fatigue propagation resistance of the surface layer may degrade due to the local low-temperature transformation structure such as martensite or bainite in the surface layer, generated during the cooling process, and accordingly, it may be difficult to ensure proper fatigue quality.
• An embodiment of the present disclosure is to provide a steel sheet and a method of manufacturing the same.
• An embodiment of the present disclosure is to provide an extremely thick steel material for a flange, having excellent fatigue properties and low-temperature impact toughness, and a method of manufacturing the same.
• a steel sheet including, by weight%, carbon (C): 0.05 to 0.20%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.00 to 2.00%, aluminum (Al): 0.005 to 0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005 to 0.070%, vanadium (V): 0.001 to 0.300%, titanium (Ti): 0.001 to 0.050%, chromium (Cr): 0.01 to 0.30%, molybdenum (Mo): 0.01 to 0.12%, copper (Cu): 0.01 to 0.60%, nickel: 0.05 to 4.00%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities,
• the steel sheet may further include, by weight %, zirconium (Zr): 0.001 to 0.150%.
• the central portion may include one or more of ferrite and pearlite as a remainder structure.
• one or more of fine NbC, NbCN, VC, CVN precipitates having a diameter of 5 to 50 nm may be 15 or more per 1 ⁇ m 2 .
• a hardness value of the surface layer portion may be 250HB or less.
• the steel sheet may have a tensile strength of 590 to 820 MPa and a fatigue limit ratio (tensile strength/fatigue strength) of 0.30 or more.
• the steel sheet may have a thickness of 133 to 233 mm.
• a method of manufacturing a steel sheet including primary-reheating a steel slab including, by weight%, carbon (C): 0.05 to 0.20%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.00 to 2.00%, aluminum (Al): 0.005 to 0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005 to 0.070%, vanadium (V): 0.001 to 0.300%, titanium (Ti): 0.001 to 0.050%, chromium (Cr): 0.01 to 0.30%, molybdenum (Mo): 0.01 to 0.12%, copper (Cu): 0.01 to 0.60%, nickel: 0.05 to 4.00%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities;
• the steel slab may further include zirconium (Zr): 0.001 to 0.150% by weight%.
• the primary-reheating may be performed in a temperature range of 1100 to 1300°C, and
• the secondary-reheating may be performed in a temperature range of 1000 to 1200°C.
• a thickness of the steel slab may be 650 to 750 mm
• a steel sheet and a method of manufacturing the same may be provided.
• an extremely thick steel material for a flange having excellent fatigue properties and low-temperature impact toughness, and a method of manufacturing the same may be provided.
• an extremely thick steel material for a flange having excellent fatigue properties and low-temperature impact toughness, and a method of manufacturing the same, which may be used in a wind power generation tower and system, or the like.
• FIG. 1 is a CCT diagram indicating a cooling history of disclosure inventive example 1 and comparative example 7 according to an embodiment of the present disclosure.
• % indicating the content of each element may be based on weight.
• a steel sheet may include, by weight%, carbon (C): 0.05 to 0.20%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.00 to 2.00%, aluminum (A1): 0.005 to 0.100%, phosphorus (P): 0.0100% or less, sulfur (S): 0.0150% or less, niobium (Nb): 0.005 to 0.070%, vanadium (V): 0.001 to 0.300%, titanium (Ti): 0.001 to 0.050%, chromium (Cr): 0.01 to 0.30%, molybdenum (Mo): 0.01 to 0.12%, copper (Cu): 0.01 to 0.60%, nickel: 0.05 to 4.00%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities.
• Carbon (C) may be the most important element for ensuring basic strength, and thus, it may be necessary to be included in steel within an appropriate range. To obtain the effect of the addition, 0.05% or more of carbon (C) may be added. According to an embodiment of the present disclosure, 0.10% or more of carbon (C) may be added. When the content exceeds a predetermined level, during QT heat treatment, hardenability may be excessively increased, such that strength and hardness of the base material may be excessively exceeded, which may cause surface cracks during forging and may degrade low-temperature impact toughness properties of the final product. Accordingly, in the present disclosure, an upper limit of carbon (C) content may be limited to 0.20%. According to an embodiment of the present disclosure, an upper limit may be 0.18%.
• Silicon (Si) may be a substitutional element, and may improve strength of a steel material through solid solution and may have a strong deoxidation effect, such that silicon (Si) may be an essential element for manufacturing clean steel. Accordingly, silicon (Si) may be added by 0.05% or more. According to an embodiment of the present disclosure, silicon (Si) may be added by 0.20% or more. When silicon (Si) is added in a large amount, a MA (martensite-austenite) phase may be generated, and ferrite matrix strength may excessively increase, such that surface quality of an ultra-thick product may be deteriorated, and thus, an upper limit of the content may be limited to 0.50%. According to an embodiment of the present disclosure, an upper limit thereof may be 0.40%.
• Manganese (Mn) may be a useful element for improving strength by solid solution and improving hardenability such that a low-temperature transformation phase may be generated. Accordingly, 1.00% or more of manganese (Mn) may be added to ensure a tensile strength of 590MPa or more. According to an embodiment of the present disclosure, the manganese (Mn) content may be 1.10% or more. As the manganese (Mn) content increases, manganese (Mn) may form MnS, which is an elongated non-metallic inclusion, together with S, and MnS may act as an impact initiation point, which may rapidly degrade the low-temperature impact toughness of the product. Accordingly, the manganese (Mn) content may be limited to 2.00% or less. According to an embodiment of the present disclosure, the content may be 1.50% or less.
• Aluminum (Al) may be one of powerful deoxidizers in a steelmaking process along with Si, and may be added at 0.005% or more to obtain this effect.
• a lower limit of the aluminum (Al) content may be 0.010%.
• the aluminum (Al) content is excessive, a fraction of Al 2 O 3 among oxidizing inclusions generated as a result of deoxidation may increase excessively, and the size thereof may become coarse, and it may become difficult to remove the inclusions during refining, which may be a factor degrading low-temperature impact toughness. Accordingly, the aluminum (Al) content may be limited to 0.100% or less. According to an embodiment of the present disclosure, the aluminum (Al) content may be 0.070% or less.
• Phosphorus (P) may cause brittleness in the grain system or may form coarse inclusions, thereby causing brittleness. Thus, to improve brittle crack propagation resistance, phosphorus (P) may be limited to 0.0100% or less. However, considering the level inevitably included during steel manufacturing, 0% may be excluded.
• S may cause brittleness in the grain system or may cause brittleness by forming coarse inclusions. Thus, to improve brittle crack propagation resistance, sulfur (S) may be limited to 0.0015% or less. However, considering the level inevitably included during steel manufacturing, 0% may be excluded.
• Niobium (Nb) may be precipitated in the form of NbC or NbCN and may enhance strength of a base material. Also, when reheated at a high temperature, solid-solute Nb may be precipitated finely in a deformation-induced form at the recrystallization temperature or lower during forging, thereby inhibiting the growth of austenite, such that there may be the effect of refining the structure. Accordingly, niobium (Nb) may be added at 0.005% or more. According to an embodiment of the present disclosure, niobium (Nb) may be 0.010% or more.
• niobium (Nb) When niobium (Nb) is excessively added, undissolved niobium (Nb) may be generated in the form of TiNb(C,N), which may hinder low-temperature impact toughness, and thus, an upper limit of the content may be limited to 0.070%. According to an embodiment of the present disclosure, the niobium (Nb) content may be 0.065% or less.
• V Vanadium (V): 0.001 to 0.300%
• Vanadium (V) may be almost completely re-dissolved during reheating, such that strengthening effect by precipitation or solid solution during subsequent rolling may be minimal, but in the case of an extremely thick forging material, the air cooling speed may be slow, such strength may be improved by precipitating fine carbonitrides during the cooling process or tempering heat treatment process.
• vanadium (V) may be added at 0.001% or more.
• a lower limit of vanadium (V) content may be 0.010%.
• vanadium (V) content may be 0.300% or less.
• vanadium (V) content may be 0.250% or less.
• Titanium (Ti) may be a component significantly improving low-temperature toughness by precipitating as TiN during reheating and inhibiting the growth of prior austenite grains at high temperature, and thus, titanium (Ti) may be added at 0.001% or more to obtain the effect of the addition.
• titanium (Ti) When titanium (Ti) is excessively added, low-temperature toughness may be reduced due to clogging of a continuous casting nozzle or crystallization of a central portion, and titanium (Ti) may be combined with N, such that coarse TiN precipitates may be formed in the thickness central portion, which may degrade the elongation of the product, and accordingly, uniform elongation may be degraded during the forging process, causing surface cracks.
• the titanium (Ti) content may be 0.050% or less.
• the titanium (Ti) content may be 0.030% or less, and may be 0.018% or less.
• Chromium (Cr) may have the effect of preventing a decrease in strength by slowing down the rate of spheroidization of cementite, and may be a component to improve hardenability during the cooling process. For this effect, 0.01% or more of chromium (Cr) may be added.
• the chromium (Cr) content is excessive, the size and fraction of Cr-rich coarse carbides such as M 23 C 6 may increase, which reduces impact toughness of the product may degrade, and solubility of Nb in the product and the fraction of fine precipitates such as NbC may decrease, such that strength degradation of the product may become a problem.
• an upper limit of the content may be limited to 0.30%. According to an embodiment of the present disclosure, an upper limit of the chromium (Cr) content may be 0.25%.
• Molybdenum (Mo) may increase grain boundary strength, may increase hardenability, and may be in solid-solution in precipitates, thereby improving strength, such that molybdenum (Mo) may effectively contribute to increasing strength and ductility of a product. Also, molybdenum (Mo) may have the effect of preventing toughness degradation due to grain boundary segregation of impurities such as P. For this effect, 0.01% or more of molybdenum (Mo) may be added. However, since molybdenum (Mo) is an expensive element, excessive addition may significantly increase manufacturing costs, such that an upper limit of the content may be limited to 0.12%.
• Copper (Cu) may greatly improve strength of a matrix phase by solid-solution strengthening in ferrite. For this effect, 0.01% or more of copper (Cu) may be included. A more preferable copper (Cu) content may be 0.03% or more. However, when the copper (Cu) content is excessive, star cracks may be likely to occur on the surface of the steel sheet, and the manufacturing costs may increase significantly as copper (Cu) is an expensive element, and thus, in the present disclosure, an upper limit of the content may be limited to 0.60%. According to an embodiment of the present disclosure, an upper limit thereof may be 0.50%.
• Nickel (Ni) may increase stacking faults at low temperature and may facilitate cross slip of dislocations, such that nickel (Ni) may contribute to improving impact toughness, and nickel (Ni) may contribute to improving strength by improving hardenability and solid-solution strengthening. For this effect, 0.05% or more of nickel (Ni) may be added. According to an embodiment of the present disclosure, the nickel (Ni) content may be 0.10% or more. When nickel (Ni) is excessively added, the manufacturing cost may also increase due to the high cost, and thus, an upper limit of the content may be limited to 4.00%. According to an embodiment of the present disclosure, an upper limit of the nickel (Ni) content may be 3.50%.
• Calcium (Ca) may combine with S, which can form MnS inclusions, when calcium (Ca) is added after deoxidation by Al, thereby inhibiting the formation of MnS, and at the same time forming spherical CaS, which has the effect of inhibiting the occurrence of cracks due to hydrogen-induced cracking.
• calcium (Ca) may be added at 0.0005% or more.
• the remaining calcium (Ca) after forming CaS may combine with oxygen and coarse oxidized inclusions may be formed, which may be elongated or destroyed during forging, such that elongation and low-temperature impact toughness properties may be degraded, and thus, an upper limit of calcium (Ca) content may be limited to 0.0040%.
• the steel material of the present disclosure may further include iron (Fe) other than the above-described composition.
• Fe iron
• inevitable impurities may be inevitably added, and thus, impurities may not be excluded.
• a person skilled in the art of a general steel manufacturing process may be aware of the impurities, and thus, the descriptions of the impurities may not be provided in the present disclosure.
• zirconium (Zr): 0.001 to 0.150% may be further included by weight%.
• Zirconium (Zr) may be a strong carbide-forming element, may be present in the form of ZrC, and may improve strength of the matrix phase through the form of precipitation strengthening, such as VC or NbC. For this effect, 0.001% or more of zirconium (Zr) may be added. However, when added excessively, the manufacturing cost may increase due to the high cost, and thus, an upper limit thereof may be limited to 0.150%.
• the steel sheet according to an embodiment of the present disclosure may have a Ceq value of 0.40 to 0.65, defined in relational expression 1 as below.
• Ceq [C] + [Mn]/6 + ([Ni]+[Cu])/15 + ([Cr]+[Mo]+[V])/5
• Ceq When the Ceq value defined in relational expression 1 is less than 0.40, there may be a problem in ensuring the appropriate strength of 590 to 820 MPa, which is required in the present disclosure, due to insufficient hardenability. When the value exceeds 0.65, the strength value may excessively increase and the -50°C low-temperature impact toughness may deteriorate. Accordingly, Ceq may be preferably 0.40 to 0.65.
• the percentages indicating the fraction of microstructure may be based on area.
• the surface layer portion which is the region from the surface to 1/8 point in the thickness direction, may include, by area%, 80% or more of polygonal ferrite and a remainder of bainite, and the central portion, which is the region from 1/8 point to 1/2 point in the thickness direction, may include 80% or more of one or more of bainite and martensite.
• the surface layer portion which is a region from the surface to the 1/8 point in the thickness direction, may be limited to include 80% or more of polygonal ferrite.
• the polygonal ferrite fraction is less than 80%, crack propagation resistance of the surface layer portion may decrease and the fatigue limit ratio of the steel material may not be ensured to the target level.
• the central portion which is the region from the 1/8 point to the 1/2 point, may include one or more of bainite and martensite by 80% or more to ensure the desired strength.
• one or more of ferrite and pearlite may be included.
• an average size of packets having a high-angle grain boundary of 15° or more as of bainite or martensite may be 25 ⁇ m or less.
• the average size of packets may be limited to 25 ⁇ m or less to ensure the desired impact toughness.
• a difference between a porosity of the region from the surface to a 3/8 point in the thickness direction and a porosity of the region from the 3/8 point to a 5/8 point in the thickness direction may be 0.100mm 3 /g or less.
• the value obtained by subtracting the porosity of the region from the surface to the 3/8 point in the thickness direction from the porosity of the region from the 3/8 point to the 5/8 point in the thickness direction of the steel sheet may be 0.100mm 3 /g or less.
• the porosity may be obtained by measuring a density (g/mm 3 ) and taking a reciprocal (mm 3 /g).
• the difference When the difference exceeds 0.100mm 3 /g, retained voids may act as a crack initiation point during impact testing, such that it may be difficult to ensure -50°C low-temperature impact toughness properties targeted in the present disclosure.
• the difference may be 0.050mm 3 /g or less.
• the difference may be 0.030mm 3 /g or less. It may not be necessary to particularly limit a lower limit of the difference, but in an embodiment, a lower limit may be 0mm 3 /g.
• one or more of fine NbC, NbCN, VC, CVN precipitates having a diameter of 5 to 50 nm may be present at a rate of 15 or more per 1 ⁇ m2.
• the precipitation strengthening effect may be weakened and the pinning effect and precipitation strengthening effect may be lost, such that it may not be easy to ensure the level of strength required in the present disclosure.
• the steel sheet according to an embodiment of the present disclosure may have a thickness of 133 to 233 mm, a hardness value of the surface layer of 250 HB or less, a tensile strength of 590 to 820 MPa, a Charpy impact absorption energy value of 50 J or more at -50°C, and a fatigue limit ratio (tensile strength/fatigue strength) of 0.30 or more.
• the steel sheet according to an embodiment of the present disclosure may be manufactured into a flange having a maximum thickness of 200 to 500 mm, an inner diameter of 4000 to 7000 mm, and an outer diameter of 5000 to 8000 mm.
• the steel sheet according to an embodiment of the present disclosure may be manufactured by primary-reheating, forging, secondary-reheating, hot-rolling, primary-cooling, secondary-cooling and tempering a steel slab satisfying the alloy composition of the present disclosure.
• a steel slab having a thickness of 650 to 750 mm satisfying the alloy composition of the present disclosure may be primary-reheated in a temperature range of 1100 to 1300°C.
• the primary-reheating may be performed in a temperature range of 1100°C or higher for Ti or Nb composite carbonitrides or TiNb(C,N) coarse crystallites formed during casting to be re-dissolved, to homogenize the structure by heating austenite to the recrystallization temperature or higher before forging, and to reduce surface layer cracks which may occur during the forging process by ensuring the forging end temperature to be sufficiently high.
• an upper limit of the steel slab heating temperature may be limited to 1300°C.
• a thickness of the steel slab during primary-reheating may be 650 to 750 mm. According to an embodiment of the present disclosure, the thickness of the steel slab may be 700 to 750 mm.
• the slab may be manufactured using one of a continuous casting process, a semi-continuous casting process and an ingot casting process.
• the primary-reheated steel slab may be forged with a cumulative reduction ratio of 35 to 65% and a strain rate of 1.0 to 4.0/s.
• the forging may be of processing the heated steel slab into the desired shape of a final intermediate product.
• High-strain, low-speed forging may be essential to sufficiently compress voids, and thus, the forging may be performed under the conditions of a cumulative reduction ratio of 35 to 65% and a strain rate of 1.0/s to 4.0/s.
• the strain rate may indicate the strain rate per unit time, and the unit may be %/s.
• the cumulative reduction ratio of forging When the cumulative reduction ratio of forging is less than 35%, the voids retained in the steel slab may not be sufficiently compressed, such that retained voids may be formed, which may degrade impact resistance of the product. According to an embodiment of the present disclosure, the cumulative reduction ratio of forging may be 40% or more. When the cumulative reduction ratio exceeds 65% and the number of forging passes increases, the surface temperature may continue to decrease, which may cause surface cracks at low temperature.
• the cumulative reduction ratio at the non-recrystallization temperature or lower exceeds 20%
• the uniform elongation value of the surface layer of the steel material at about 850°C to 1150°C, where forging is performed exceeds 20 %. Consequently, due to work hardening caused by overlapping dislocations, the uniform elongation of the surface is significantly reduced, and surface cracks may occur during the forging process. Accordingly, the cumulative reduction ratio at the recrystallization temperature or lower may be limited to 20% or less. According to an embodiment of the present disclosure, the ratio may be 15% or less.
• the thickness of the steel slab after forging may be 350 to 450 mm.
• the forged steel slab may be secondary-reheated in a temperature range of 1000 to 1200°C.
• reheating may be performed in the temperature range of 1000°C or higher.
• reheating may be performed in the temperature range of 1000°C or higher.
• the secondary-reheated steel slab may be hot-rolled at a finishing rolling temperature of Tnr-50 to Tnr+50°C and an accumulated reduction ratio of 20.0% or more.
• the finishing rolling temperature when the finishing rolling temperature is less than Tnr-50°C, the deformation resistance may increase excessively as the temperature decreases, such that it may be difficult to sufficiently refine austenite grains in the central portion in the thickness direction of the product, and accordingly, low-temperature impact toughness properties of the final product may be deteriorated.
• the finishing rolling temperature exceeds Tnr+50°C, austenite grains may become excessively coarse, such that strength and impact toughness may be deteriorated.
• the prior austenite grain size may not be sufficiently refined.
• the thickness of the steel sheet after hot-rolling may be 133 to 233 mm.
• Tnr 887 + 464[C] + 890[Ti] + 363[Al] - 357[Si] + 6445 [Nb] - 644([Nb] 1/2 ) + 732 [V] - 230([V] 1/2 )
• the hot-rolled steel sheet may be heated in a temperature range of 820 to 900°C and may be held for 10 to 40 minutes, and may be primary-cooled to 700°C based on the surface temperature of the steel sheet at an average cooling rate of 0.1 to 5.0°C/s.
• the primary-cooling may be performed to form the desired polygonal ferrite from the surface of the steel material to a region corresponding to a 1/8 point in the thickness direction.
• the temperature and cooling rate may be based on the surface temperature of the steel sheet.
• the waiting time in the air may be excessively long, such that the manufacturing time may increase, which may degrade productivity.
• the cooling rate exceeds 5.0°C/s, polygonal ferrite may not be formed, and bainite or martensite, which are low-temperature transformation structures, may be formed, such that the surface uniform elongation of the final product may be degraded and fatigue quality may deteriorate.
• the primary-cooled steel sheet may be secondary-cooled to room temperature at an average cooling rate of 10.0°C/s or higher based on the surface temperature of the steel sheet.
• the secondary-cooling may be performed such that the microstructure of a region other than the surface layer includes bainite or a dual-phase structure of bainite and ferrite through strong cooling.
• the average cooling rate may be based on the surface temperature of the steel sheet.
• the average cooling rate is less than 10.0°C/s, it may be difficult to obtain the low-temperature transformation structure described above.
• an upper limit of the average cooling rate may not be particularly limited in the secondary cooling, but the upper limit may be limited to 200.0°C/s or less.
• the secondary cooling may be performed by performing quenching while slowing down the passing speed of the steel material, increasing the flow rate of the sprayed water, or the like.
• the secondary-cooled steel sheet may go through a tempering heat treatment of heating to a temperature range of 550 to 700°C and holding for 5 to 60 minutes.
• the tempering heat treatment may reduce dislocation density of bainite or bainite and ferrite mixed structure, which is the low-temperature transformation structure, and may improve strength and toughness by diffusing carbon in a short range.
• tempering heat treatment temperature is less than 550°C, carbon diffusion may not be sufficient, such that strength may be excessively increased, which may degrade toughness.
• temperature exceeds 700°C, fresh martensite may be formed due to reverse transformation at Ac1 or higher temperature, which may severely deteriorate impact toughness and cold workability.
• tempering heat treatment time is less than 5 minutes, there may be no enough time for carbon to sufficiently diffuse during the tempering process, such that strength may be excessively increased, which may degrade toughness and cold workability, and may exceed the appropriate strength range required in the present disclosure.
• tempering heat treatment time exceeds 60 minutes, cementite may become spheroidized due to excessive heating, which may rapidly reduce strength.
• a slab having a thickness of 700 mm and the alloy composition in Table 1 was manufactured. Forging, hot-rolling, primary, secondary-cooling, and tempering heat treatment were performed under the process conditions in Table 2, and a flange having a thickness of 320 mm was finally manufactured.
• the primary-reheating temperature 1220 to 1240°C was applied, and as for the secondary-reheating, 1140 to 1160°C was commonly applied, and as for the reheating time and tempering time for primary-cooling, 30 minutes was commonly applied.
• the hot-rolling temperature was represented based on the Tnr temperature.
• microstructure and physical properties of each manufactured sample were measured and are listed in Table 3 below.
• the microstructure was measured using an image autoanalyzer by collecting samples from the surface region (region from the surface to 1/8 point in the thickness direction) and the remaining region (region from 1/8 point to 1/2 point in the thickness direction). Also, the packet size was automatically measured and represented based on the high-angle grain boundary plane with the orientation relationship determined to 15° or more in the OIM (orientation image microscopy) map of the microstructure analyzed using EBSD (electron back scatter diffraction).
• OIM orientation image microscopy
• the porosity difference was represented as the value obtained by subtracting the porosity of the region from the surface to 3/8 point in the thickness direction from the porosity of the region from 3/8 point to 5/8 point in the thickness direction in the present disclosure.
• the porosity was obtained by measuring a density (g/mm 3 ) and taking a reciprocal (mm 3 /g).
• the number of precipitates of 5 to 50 nm observed in the cross-section of the steel material was measured using TEM. NbC precipitates were confirmed through diffraction patterns of NbC and VC and EDX mapping, and the number of precipitates positioned at 1 ⁇ m 2 was counted.
• the hardness value was measured at three positions on the surface using a Brinell hardness tester and an average value thereof was listed.
• Tensile strength was evaluated through a room temperature tensile test, and as for impact toughness for each sample, an average of the absorbed energy values measured three times at the corresponding temperature through the Charpy V-Notch Test.
• Fatigue strength was measured using a rotational bending fatigue test specimen, after electrolytic polishing was performed to remove the influence of the machined layer on the surface, using the method described in JIS Z 2274. As the fatigue limit ratio, the ratio between fatigue strength to tensile strength was calculated and listed.
• FIG. 1 is a CCT diagram indicating a cooling history of disclosure inventive example 1 and comparative example 7 according to an embodiment of the present disclosure.
• inventive example 1 through multistage cooling, the surface layer portion was controlled to have a polygonal ferrite structure and the central portion to have a bainite structure, thereby ensuring both fatigue properties and strength targeted in the present disclosure.
• comparative example 7 cooling was performed only with strong cooling, and low-temperature transformation structures were formed in both the surface and central portions, such that fatigue properties may be deteriorated.
• the cumulative reduction ratio during forging exceeded the range of the present disclosure.
• the cumulative reduction ratio during forging was excessive, such that work hardening occurred, and low-temperature impact toughness deteriorated.
• the cumulative reduction ratio during forging was insufficient, such that retained voids at the center of the steel sheet were not properly controlled, and accordingly, fatigue strength and impact toughness deteriorated.
• the finishing rolling temperature was not reached during hot-rolling. Accordingly, the packet size in the central portion was coarse, and the surface hardness was excessively high, such that the fatigue crack resistance was degraded.
• the cumulative reduction ratio was insufficient during hot-rolling. Accordingly, the packet size at the target level was not ensured, and impact toughness was degraded.
• the cooling rate was insufficient during secondary-cooling, and ferrite was excessively formed on the central portion, such that appropriate strength was not ensured, such that the strength targeted in the present disclosure was not ensured.
• the tempering temperature was excessively high, such that the formation of precipitates was not smooth, and accordingly, the strength targeted in the present disclosure was not ensured.
• the carbon content exceeded the range suggested in the present disclosure. Accordingly, in comparative example 10, the carbon content was insufficient, such that the strength targeted in the present disclosure was not ensured. As for comparative example 11, the carbon content exceeded the range suggested in the present disclosure, such that an appropriate level of strength was not ensured, and impact toughness was deteriorated.
• the Mn content was excessive, such that bainite was excessively formed in the surface layer, and as a result, strength and toughness were not properly ensured.
• the Nb content did not satisfy the content suggested in the present disclosure. As a result, the formation of precipitates was not easy, and the targeted strength was not ensured.

Landscapes

• Chemical & Material Sciences (AREA)
• Engineering & Computer Science (AREA)
• Mechanical Engineering (AREA)
• Materials Engineering (AREA)
• Metallurgy (AREA)
• Organic Chemistry (AREA)
• Crystallography & Structural Chemistry (AREA)
• Physics & Mathematics (AREA)
• Thermal Sciences (AREA)
• Heat Treatment Of Steel (AREA)
• Heat Treatment Of Articles (AREA)

Applications Claiming Priority (2)

Publications (2)

Family

ID=91589109

Family Applications (1)

Country Status (5)

Citations (2)

Family Cites Families (6)

• 2022
  • 2022-12-22 KR KR1020220181522A patent/KR20240100531A/ko active Pending
• 2023
  • 2023-10-26 JP JP2025536372A patent/JP2025541901A/ja active Pending
  • 2023-10-26 WO PCT/KR2023/016717 patent/WO2024136087A1/ko not_active Ceased
  • 2023-10-26 EP EP23907373.7A patent/EP4640906A4/de active Pending
  • 2023-10-26 AU AU2023412500A patent/AU2023412500A1/en active Pending

Patent Citations (2)

Non-Patent Citations (1)

Also Published As

Similar Documents

Legal Events