Steel sheet and manufacturing method therefor

Abstract

The present invention relates to a steel sheet and a manufacturing method therefor, and, more specifically, to a steel sheet for an extremely thick pressure vessel, having excellent cold formability and lamellar tearing resistance quality, which can be used in a petrochemical manufacturing facility, a storage tank, or the like, and a manufacturing method therefor.

Description

TECHNICAL FIELD

[0001] The present disclosure relates to a steel sheet and a manufacturing method therefor, and more particularly, to a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamella tearing quality, and a method of manufacturing the same.BACKGROUND ART

[0002] Recently, due to enlargement of crude oil refining and storage facilities and demand for high-capacity storage, demand for increasing a thickness of steel used in the facility has been continuously increased. In manufacturing a large structure, to improve the internal and external soundness of steel, a concentration of impurities such as non-metallic inclusions or segregation are being reduced, or cracks and voids on the surface and inside the material are controlled to the extreme.

[0003] In particular, in the case of an extremely thick material having a thickness exceeding 100 mm, a rolling reduction ratio may not be high as compared to a thin material, such that unsolidified shrinkage occurring during continuous casting or casting may not be sufficiently compressed during a rough rolling process and may remain in the form of retained voids in a central portion of a product. The retained voids may act as an initiation point of cracks when the structure receives stress in a thickness axial direction, and may eventually cause damage to the entire facility in the form of lamellar tearing. Accordingly, a process to sufficiently compress central voids such that no retained voids are present may be necessary in the step before rolling.

[0004] Cited document 1, related thereto, may disclose a steel reduction technique in a thick steel rough rolling process, and may use a technique of determining a limit reduction ratio for different thickness ratios at which sheet sticking occurs from the steel reduction ratio for each pass determined to be close to a design allowable value (load and torque) of the rolling mill, a technique of distributing a reduction ratio by adjusting an index of thickness ratio for each pass to ensure a target thickness of a rough rolling machine, a technique of modifying a reduction ratio to prevent sheet sticking based on the limit reduction ratio for different thicknesses, and a manufacturing method for allowing an average reduction ratio of approximately 27.5% in the final 3 passes of rough rolling based on a thickness of 80 mm may be provided. However, in the case of the above rolling method, the average reduction ratio of the entire product thickness may be measured, such that it may be difficult to apply high strain to a central portion of an extremely thick material having a maximum thickness of 233 mm in which retained voids are present.

[0005] One of the other methods for manufacturing an extremely thick material may be a method of using a forging machine having a higher effective strain per pass than a rolling mill. Cited document 2 is a method of vertically erecting the continuous casting slab extracted from the furnace and providing a full width forging reduction of 400 mm or more, eliminating pores in the edge portion in the width direction and a central portion by performing a forging pass with a reduction of less than 2 passes, which is within the buckling limit reduction ratio, and increasing a central portion strain rate. The retained void in the central portion, which was a problem in cited document 1, may be effectively compressed, lamellar tearing quality of the product may be improved when applied to forging.

[0006] However, there may be a disadvantage in that surface defects may occur due to local strain concentration during a width forging process. In particular, when a surface or subsurface defect is present in the cast state before forging, the defect may propagate during the forging process and the surface quality may become inferior in the product state after rolling.

[0007] Cited document 3 discloses that a thick-walled, high-strength steel sheet having a yield strength of 620 MPa or more and a thickness of 100 mmt or more may be manufactured through a process of heating a material provided with a predetermined alloy composition at 1200 to 1350° C., performing hot forging with a cumulative reduction of 25% or more, heating at Ac3 point or more and 1200° C. higher, performing hot rolling with a cumulative reduction of 40% or more, reheating to Ac3 point or more and 1050° C. or higher, rapidly cooling from the temperature of Ac3 point or more to a lower temperature of 350° C. or less or Ar3 point or less, and tempering at a temperature of 450° C. to 700° C. However, in the case of the ultra-high strength steel sheet described above, the carbon equivalent (Ceq) and hardenability index may be high, such that the sheet may be vulnerable to surface cracks during casting, and rapid cooling in the austenite single phase region may locally form rapid cooling structures of bainite or martensite on the surface layer of the sheet, and the sheet may not be used for a storage tank and a pressure vessel having excellent cold formability.PRIOR ARTCited Document

[0008] (Cited document 1) Korean Laid-Open Patent Publication No. 10-2012-0075246 (publicized on Jul. 6, 2012)

[0009] (Cited document 2) Korean Laid-Open Patent Publication No. 10-2012-0074039 (publicized on Jul. 5, 2012)

[0010] (Cited document 3) Korean Laid-Open Patent Publication No. 10-2017-7019203 (publicized on Jan. 15, 2016)DETAILED DESCRIPTION OF PRESENT DISCLOSURETechnical Problems to Solve

[0011] An embodiment of the present disclosure is to provide a steel sheet and a method of manufacturing the same.

[0012] An embodiment of the present disclosure is to provide a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamella tearing quality, and a method of manufacturing the same.

[0013] The subject matter of the present disclosure is not limited to the above-described contents. Those skilled in the art will have no difficulty in understanding additional purposes of the present disclosure from the overall contents of this specification.Solution to Problem

[0014] According to an embodiment of the present disclosure, provided is a steel sheet including, by weight %, carbon (C): 0.10 to 0.25%, 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): 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V): 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities,

[0015] wherein, in a microstructure, a surface layer portion, which is a region from a surface to a ⅛ point in a thickness direction, includes 80% or more of polygonal ferrite by area %, and a central portion, which is the other region from a ⅛ point to a ½ point in the thickness direction, includes 50% or more of bainite and a remainder of ferrite by area %,

[0016] wherein a difference between a porosity of a region from a surface to a ⅜ point in the thickness direction and a porosity of a region from a ⅜ point to a ⅝ point in the thickness direction is 0.100 mm3 / g or less, and

[0017] wherein a hardness value of the surface layer portion is 200 to 220 HB.

[0018] The surface layer portion may include one or more of pearlite and bainite as a remainder structure.

[0019] In the steel sheet, one or more of fine NbC, NbCN, VC, CVN precipitates having a diameter of 5 to 50 nm may be present in a density of 10 or more per 1 μm2.

[0020] The steel sheet may have a tensile strength of 510 to 690 MPa, an elongation (ZRA) in the thickness direction of 35% or more, and a Charpy impact absorption energy value at −50° C. of 50 J or more.

[0021] The steel sheet may have a maximum surface crack depth of 1 μm or less in a 180° bending test at room temperature.

[0022] The steel sheet may have a thickness of 133 to 233 mm.

[0023] According to an embodiment of the present disclosure, provided is a method of manufacturing a steel sheet including primary-reheating a steel slab including, by weight %, carbon (C): 0.10 to 0.25%, 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) 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V) 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities;

[0024] forging the primary-reheated steel slab at a cumulative reduction ratio of 35 to 65% and a strain rate of 1.0 to 4.0 / s;

[0025] secondary-reheating the forged steel slab;

[0026] hot-rolling the secondary-reheated steel slab at a finishing-rolling temperature of 900 to 1100° C.;

[0027] heating the hot-rolled steel sheet in a temperature range of 820 to 900° C., holding for 10 to 40 minutes, and primary-cooling the steel sheet to 700° C. at an average cooling rate of 0.1 to 5.0° C. / s based on a steel sheet surface temperature;

[0028] secondary-cooling the primary-cooled steel sheet to room temperature at an average cooling rate of 10.0° C. / s or higher based on the steel sheet surface temperature; and

[0029] performing a tempering heat treatment of heating the secondary-cooled steel sheet in a temperature range of 550 to 700° C. and holding for 5 to 60 minutes,

[0030] wherein, in the forging, a cumulative reduction ratio at a recrystallization temperature or lower is 20% or less.

[0031] The primary-reheating may be performed in a temperature range of 1100 to 1300° C., and

[0032] The secondary-reheating may be performed in a temperature range of 1000 to 1200° C.

[0033] In the primary-reheating, a thickness of the steel slab may be 650 to 750 mm,

[0034] after the forging, a thickness of the steel slab may be 350 to 450 mm, and

[0035] after the hot-rolling, a thickness of the steel sheet may be 133 to 233 mm.Advantageous Effects of Invention

[0036] According to an aspect of the present disclosure, a steel sheet and a method of manufacturing the same may be provided.

[0037] According to an aspect of the present disclosure, a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamella tearing quality, and a method of manufacturing the same may be provided.

[0038] According to an aspect of the present disclosure, a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamellar tearing quality which may be used in a petrochemical manufacturing facility, a storage tank, or the like.BRIEF DESCRIPTION OF DRAWINGS

[0039] FIG. 1 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of inventive example 1 according to an embodiment of the present disclosure.

[0040] FIG. 2 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of comparative example 6, which falls outside the scope of one embodiment of the present disclosure.BEST MODE FOR INVENTION

[0041] Hereinafter, preferred embodiments of the present invention may be described. The implementation examples of the present disclosure may be modified in various forms, and the scope of the present disclosure should not be construed as being limited to the implementation examples described below. The implementation examples are provided to describe the present disclosure in greater detail to those skilled in the art to which the present invention pertains.

[0042] In the description below, the present disclosure will be described in detail.

[0043] In the description below, a composition of steel of the present disclosure will be described in detail.

[0044] Unless otherwise indicated, the percentages indicating the content of each element may be based on weight.

[0045] A steel sheet according to an embodiment of the present disclosure may include, by weight %, carbon (C): 0.10 to 0.25%, 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): 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V): 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities.

[0046] Carbon (C): 0.10 to 0.25%

[0047] Carbon (C) may be the most important element for ensuring basic strength, and thus, it may be necessary to be contained in steel within an appropriate range. To obtain the effect of the addition, 0.10% or more of carbon (C) may be added. According to an embodiment of the present disclosure, 0.12% or more of carbon (C) may be added. When the content exceeds a predetermined level, a fraction of low-temperature transformation structure may increase during QT heat treatment, and strength and hardness of the base material may be excessive, and accordingly, uniform elongation may be degraded due to the formation of a low-temperature structure rather than an elongated ferrite as a surface structure. Accordingly, in the present disclosure, an upper limit of carbon (C) content may be limited to 0.25%. According to an embodiment of the present disclosure, an upper limit thereof may be 0.20%.

[0048] Silicon (Si): 0.05 to 0.50%

[0049] Silicon (Si) may be a substitutional element, and may improve strength of a steel material through solid solution strengthening 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 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%.

[0050] Manganese (Mn): 1.00 to 2.00%

[0051] Manganese (Mn) may be a useful element for improving strength by solid solution strengthening and improving hardenability by generating a low-temperature transformation phase. Accordingly, 1.00% or more of manganese (Mn) may be added to ensure a tensile strength of 450 MPa 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 with S, which is an elongated non-metallic inclusion, which may degrade toughness, and when tensioned in the thickness direction, MnS may act as a factor degrading an elongation, and thus, MnS may be a factor rapidly degrading lamellar tearing quality. 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.

[0052] Aluminum (Al): 0.005 to 0.100%

[0053] 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. According to an embodiment of the present disclosure, a lower limit of the aluminum (Al) content may be 0.01%. When the aluminum (Al) content is excessive, a fraction of Al2O3 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 lamellar tearing properties. 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.

[0054] Phosphorus (P): 0.0100% or less

[0055] Phosphorus (P) may cause brittleness in the grain system or may form coarse inclusions, thereby causing brittleness. Thus, to improve resistance to brittle crack propagation, phosphorus (P) may be limited to 0.0100% or less. However, considering the level inevitably contained during steel manufacturing, 0% may be excluded.

[0056] Sulfur (S): 0.0015% or less

[0057] Sulfur (S) may cause brittleness in the grain system or forming coarse inclusions, thereby causing brittleness. Thus, to improve resistance to brittle crack propagation, sulfur (S) may be limited to 0.0015% or less. However, considering the level inevitably contained during steel manufacturing, 0% may be excluded.

[0058] Niobium (Nb): 0.001 to 0.030%

[0059] Niobium (Nb) may be precipitated in the form of NbC or NbCN and may enhance strength of a base material. Also, when reheating at a high temperature, dissolved Nb may be relatively finely precipitated in the form of NbC during rolling, which may suppress recrystallization of austenite and may have the effect of refining the structure. Accordingly, niobium (Nb) may be added at 0.001% or more. According to an embodiment of the present disclosure, niobium (Nb) may be 0.005% or more. When niobium (Nb) is excessively added, undissolved niobium (Nb) may be generated in the form of TiNb(C,N), which may hinder lamellar tearing properties, and thus, an upper limit of the content may be limited to 0.030%. According to an embodiment of the present disclosure, the niobium (Nb) content may be 0.020% or less.

[0060] Vanadium (V): 0.001 to 0.030%

[0061] 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 vanadium (V) may have the effect of improving strength by precipitating fine carbonitrides during a subsequent heat treatment process such as PWHT. To sufficiently obtain this effect, vanadium (V) may be added at 0.001% or more. According to an embodiment of the present disclosure, a lower limit of vanadium (V) content may be 0.010%. When the content is excessive, vanadium (V) may excessively increase strength and hardness of a base material and a welded zone, which may act as a factor for surface cracks during pressure vessel processing, and manufacturing costs may rapidly increase, which may not be commercially advantageous. Accordingly, the vanadium (V) content may be 0.030% or less. According to an embodiment of the present disclosure, vanadium (V) content may be 0.020% or less.

[0062] Titanium (Ti): 0.001 to 0.030%

[0063] Titanium (Ti) may be a component significantly improving low-temperature toughness by precipitating as TiN during reheating and suppressing grain growth in the base material and the weld heat-affected zone, and thus, titanium (Ti) may be added at 0.001% or more to obtain the effect of the addition. When titanium (Ti) is excessively added, it may cause clogging of the continuous casting nozzle or segregation of a central portion, leading to a reduction in low-temperature toughness. In addition, 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, lamellar tearing properties of the final product may be degraded. Accordingly, the titanium (Ti) content may be 0.030% or less. According to an embodiment of the present disclosure, the titanium (Ti) content may be 0.025% or less, and may be 0.018% or less.

[0064] Chromium (Cr): 0.01 to 0.20%

[0065] Chromium (Cr) may increase hardenability and may form a low-temperature transformation structure, thereby increasing yield and tensile strength, and chromium (Cr) may slow down the decomposition rate of cementite during tempering after rapid cooling or heat treatment after welding, thereby preventing the decrease in strength. For this effect, 0.01% or more of chromium (Cr) may be added. When the chromium (Cr) content is excessive, the size and fraction of Cr-rich coarse carbides such as M23C6 may increase, thereby impact toughness of the product may degrade. In addition, 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. Thus, an upper limit of the content may be limited to 0.20%. According to an embodiment of the present disclosure, an upper limit of the chromium (Cr) content may be 0.15%.

[0066] Molybdenum (Mo): 0.01 to 0.15%

[0067] Molybdenum (Mo) may increase grain boundary strength and may have a relatively large solid solution strengthening effect in ferrite, 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 0.15%.

[0068] Copper (Cu): 0.01 to 0.50%

[0069] Copper (Cu) may greatly improve strength of a matrix phase by solid solution strengthening in ferrite, and may have the effect of inhibiting corrosion in a wet hydrogen sulfide atmosphere, and thus, copper (Cu) may be a favorable element in the present disclosure. 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 0.50%. According to an embodiment of the present disclosure, an upper limit thereof may be 0.30%.

[0070] Nickel (Ni): 0.05 to 0.50%

[0071] Nickel (Ni) may increase stacking faults at low temperature and may facilitate cross slip of dislocations, thereby improving impact toughness and hardenability and improving strength. 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 0.50%. According to an embodiment of the present disclosure, an upper limit of the nickel (Ni) content may be 0.30%.

[0072] Calcium (Ca): 0.0005 to 0.0040%

[0073] When calcium (Ca) added after deoxidation by Al, calcium (Ca) may combine with S, which forms MnS inclusions, such that the formation of MnS may be suppressed. In addition, calcium (Ca) may form spherical CaS thereby cracks due to hydrogen-induced cracking may be suppressed. To sufficiently form S contained as impurities into CaS, calcium (Ca) may be added at 0.0005% or more. However, when the amount added is excessive, 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 rolling, such that lamellar tearing properties may be degraded, and thus, an upper limit of calcium (Ca) content may be limited to 0.0040%.

[0074] The steel material of the present disclosure may further include iron (Fe) other than the above-described composition. In a general manufacturing process, 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.

[0075] Hereinafter, the steel microstructure of the present disclosure will be described in detail.

[0076] Unless otherwise indicated in the present disclosure, “%” indicating a fraction of the microstructure may be based on an area.

[0077] According to an embodiment of the present disclosure, as for a microstructure of the steel sheet, by area %, a surface layer portion, which is a region from a surface to a ⅛ point in a thickness direction, may include 80% or more of polygonal ferrite as area %, and a central portion, which is a region from a ⅛ point to a ½ point in the thickness direction, may include 50% or more of bainite and a remainder of ferrite.

[0078] In the present disclosure, to improve cold workability through the formation of a soft microstructure of the surface layer, the surface layer portion, which is a region from the surface to the ⅛ point in the thickness direction, may be limited to include 80% or more of polygonal ferrite. In an embodiment of the present disclosure, polygonal ferrite may be 90% or more. In addition to polygonal ferrite, the remainder structure may include one or more of pearlite and bainite.

[0079] The central portion, which is the region from the ⅛ point to the ½ point, which is the other region, may include 50% or more of bainite to ensure strength suggested in the present disclosure, and may include ferrite as the remainder structure.

[0080] According to an embodiment of the present disclosure, a difference between a porosity of the region from the surface to a ⅜ point in the thickness direction and a porosity of the region from the ⅜ point to a ⅝ point in the thickness direction may be 0.100 mm3 / g or less.

[0081] That is, the value obtained by subtracting the porosity of the region from the surface to the ⅜ point in the thickness direction from the porosity of the region from the ⅜ point to the ⅝ point in the thickness direction of the steel sheet may be 0.100 mm3 / g or less. Here, the porosity may be obtained by measuring a density (g / mm3) and taking a reciprocal (mm3 / g).

[0082] In the present disclosure, the difference in porosities between the surface layer portion and the central portion may be limited for lamellar tearing properties. When the difference in porosities exceeds 0.100 mm3 / g, the retained void may act as a crack initiation point degrading the elongation during the thickness direction tensile test, such that the elongation (ZRA) in the thickness direction may be deteriorated. According to an embodiment of the present disclosure, the difference may be 0.050 mm3 / g or less. According to an embodiment of the present disclosure, the difference may be 0.030 mm3 / g or less. It may not be necessary to specifically limit a lower limit of the difference, but in an embodiment, a lower limit may be 0 mm3 / g.

[0083] According to an embodiment of the present disclosure, one or more of fine NbC, NbCN, VC, and CVN precipitates having a diameter of 5 to 50 nm may be present in a density of 10 or more per 1 μm2.

[0084] In the present disclosure, the precipitates may be limited to ensure strength through precipitation strengthening. When the number of precipitates is less than 10 per 1 μm2, the strengthening effect may be insufficient, such that it may be difficult to ensure appropriate strength suggested in the present disclosure. When the diameter of the precipitate is excessively small, less than 5 nm, it may be difficult to ensure strength similarly to the effect of having a small number of precipitates, and when the diameter exceeds 50 nm, the pinning effect may also be reduced, such that the strength improvement effect may be reduced. Also, when the size of precipitate exceeds 50 nm, coarse precipitates may act as brittle fracture initiation points during an impact test, and thus, the diameter of the precipitate may be 5 to 50 nm appropriately.

[0085] According to an embodiment of the present disclosure, a steel sheet may have a thickness of 133 to 233 mm, a hardness value of the surface layer portion of 200 to 220 HB, a tensile strength of 510 to 690 MPa, an elongation in the thickness direction (ZRA) of 35% or more, a Charpy impact absorption energy value at −50° C. of 50 J or more, and a maximum surface crack depth of 1 μm or less in a 180° bending test at room temperature.

[0086] Hereinafter, the method of manufacturing steel of the present disclosure will be described in detail.

[0087] 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.Primary-Reheating

[0088] 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.

[0089] The primary-reheating may be performed in a temperature range of 1100° C. or higher to re-dissolve Ti or Nb composite carbonitride or TiNb(C,N) coarse crystallites formed during casting, to homogenize the structure by heating and holding 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. When the steel slab is reheated at an excessively high temperature, problems may occur due to the oxidation scale at high temperature, and the manufacturing costs may increase excessively due to the increased cost due to heating and holding, and thus, an upper limit of the steel slab heating temperature may be limited to 1300° C.

[0090] According to an embodiment of the present disclosure, 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.Forging

[0091] 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.

[0092] The forging may be of processing the heated steel slab into the shape of a final desired 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.

[0093] In the present disclosure, the strain rate may indicate the strain rate per unit time, and the unit may indicate % / s.

[0094] When the cumulative reduction ratio of forging is less than 35%, the void retained in the steel slab may not be sufficiently compressed, such that retained voids may be formed, which may degrade lamellar tearing properties 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%, resulting in an increased number of forging passes, the surface temperature may continue to decrease, which may cause surface cracks at low temperature.

[0095] When the cumulative reduction ratio exceeds 20% at the non-recrystallization temperature or lower, at which dislocation density is recovered or not offset by recrystallization, the uniform elongation of the steel surface layer during forging at approximately 850° C. to 1150° C. exceeds 20%. As a result, due to work hardening caused by accumulated dislocations, the surface uniform elongation significantly decreases, 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.

[0096] When the strain rate is less than 1.0 / s during forging, the forging productivity may be degraded, and since the temperature of the surface layer portion is excessively reduced during the forging process, surface cracks may occur during the finishing forging process. When a strain rate thereof exceeds 4.0 / s, surface quality may be degraded due to excessive work hardening.

[0097] According to an embodiment of the present disclosure, the thickness of the steel slab after forging may be 350 to 450 mm.Secondary-Reheating

[0098] The forged steel slab may be secondary-reheated in a temperature range of 1000 to 1200° C.

[0099] To re-dissolve the composite carbonitride of Ti or Nb or the coarse-grained TiNb(C,N) formed during casting, to homogenize the structure by heating and holding the austenite to the recrystallization temperature or higher before hot-rolling, and to minimize crushing of inclusions during the rolling process by ensuring the rolling end temperature to be sufficiently high, reheating may be performed in the temperature range of 1000° C. or higher. When the slab is heated at an excessively high temperature, a problem may occur due to the oxidation scale at high temperature, and the manufacturing cost may increase excessively due to the increased cost due to heating and holding, and thus, an upper limit of the secondary heating temperature may be limited to 1200° C.Hot-Rolling

[0100] The secondary-reheated steel slab may be hot-rolled at a finishing-rolling temperature of 900 to 1100° C.

[0101] When the finishing-rolling temperature is lower than 900° C. during hot-rolling, the strain resistance value may increase excessively as the temperature decreases, such that it may be difficult to sufficiently refine austenite grains in the central portion of the product in the thickness direction, and accordingly, lamellar tearing properties of the final product may deteriorate. When the finishing-rolling temperature exceeds 1100° C., the austenite grains may become excessively coarse, such that strength and impact toughness may deteriorate.

[0102] According to an embodiment of the present disclosure, the thickness of the steel sheet after hot-rolling may be 133 to 233 mm.Primary-Cooling

[0103] 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.

[0104] In the present disclosure, 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 ⅛ point in the thickness direction. During the primary-cooling, the temperature and cooling rate may be based on the surface temperature of the steel sheet.

[0105] When the reheating, when the temperature is less than 820° C. or the holding time is less than 10 minutes, carbides generated during cooling after rolling or impurities segregated at the grain boundaries may not be re-dissolved smoothly, such that the thickness direction elongation (ZRA) and low-temperature toughness of the steel material after heat treatment may be significantly degraded. When the temperature exceeds 900° C. or the holding time exceeds 40 minutes, lamellar tearing quality may be degraded due to coarsening of austenite and coarsening of precipitated phases such as Nb(C,N), V(C,N).

[0106] When the average cooling rate is less than 0.1° C. / s during primary-cooling, the waiting time in the air may be excessively long, such that the manufacturing time may increase, which may degrade productivity. When 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 uniform elongation of the surface layer of the final product may be degraded, and thus, cold workability may be rapidly deteriorated.Secondary-Cooling

[0107] 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.

[0108] The secondary-cooling may be performed to include the microstructure of the region other than the surface layer portion into bainite or dual-phase structure of bainite and ferrite through the rapid cooling.

[0109] During the secondary-cooling, the average cooling rate may be based on the surface temperature of the steel sheet. When the average cooling rate is less than 10.0° C. / s, it may be difficult to obtain the low-temperature transformation structure described above. In the present disclosure, 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 quenching, for example, by slowing down the passing speed of the steel material and increasing the flow rate of the sprayed water, or the like.Tempering Heat Treatment

[0110] 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.

[0111] 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.

[0112] When the 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. When the 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.

[0113] When the 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. When the tempering heat treatment time exceeds 60 minutes, cementite may become spheroidized due to excessive heating, which may rapidly reduce strength.MODE FOR INVENTION

[0114] Hereinafter, the present disclosure may be described more specifically through embodiments. However, it should be noted that the embodiments below are merely intended to describe the present disclosure in greater detail based on embodiments, and are not intended to limit the scope of the rights of the present disclosure.Embodiment

[0115] 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 steel sheet having a thickness of 133 mm was finally manufactured. In this case, as for the primary-reheating temperature, 1200° C. was commonly applied, as for the secondary-reheating temperature, 1100° C. was commonly applied, and as for the reheating time for primary-cooling and tempering time, 30 minutes was commonly applied.TABLE 1SteelAlloy composition (weight %)typeCSiMnAlPSNbVTiCrMoCuNiCaA0.160.351.320.03090100.01700.0070.0110.060.090.050.1919B0.140.311.430.02782110.01100.0090.0070.080.080.070.2322C0.130.321.350.03177100.01800.0120.0050.120.100.050.2118D0.160.351.220.0288490.01700.0230.0100.10.090.100.2218E0.150.401.210.0308990.01500.0190.0110.090.060.080.1917F0.280.351.210.0307480.01900.0090.0090.120.080.090.2518G0.030.310.30.03080100.01300.0090.0050.20.080.070.1919H0.160.352.350.0307990.00050.0140.0040.250.090.060.1818I0.170.301.210.02982350.01200.1500.0050.230.100.100.1920J0.060.251.320.31085100.01500.0160.0120.150.070.060.232* Units for P, S, Ca and O may be ppm.TABLE 2ForgingReductionratio atHot-recrystal-rollingTemper-CumulativelizationFinishingPrimary-Secondary-ingreduction temperatureStrainrollingcoolingcoolingTemper-SampleSteelratioor lowerratetemperatureRateRateatureNo.type(%)(%)( / s)(° C.)(° C. / s)(° C. / s)(° C.) 1A45151.79151.324.5620 2B47132.59300.916.6618 3C51113.910501.618.4625 4D61142.79872.426.5636 5E39 92.49683.629.3641 6A84101.89154.128.0632 7B12131.99214.327.7629 8B48392.69253.816.6621 9C47148.09641.519.461810C51163.37101.826.559911D50132.798183.034.458012D59153.59900.77.466413E51143.49801.527.378814E49173.89121.536.063615G39123.910433.247.062916H42112.910542.640.062717I46171.89833.035.064018J50141.69912.733.0633The microstructure and physical properties of each manufactured sample were measured and listed in Table 3 below. The microstructure was measured by collecting samples from the surface layer region (the region from the surface to a ⅛ point in the thickness direction) and the other region (the region from the ⅛ point to the ½ point in the thickness direction) using an image auto-analyzer. In this case, the observed microstructures were polygonal ferrite (P.F), martensite (M), and bainite (B), and area % thereof may be 100%.

[0117] The porosity difference was represented as the value obtained by subtracting the porosity of the region from the surface to the ⅜ point in the thickness direction from the porosity of the region from the ⅜ point to the ⅝ point in the thickness direction in the present disclosure. Here, the porosity was obtained by measuring the density (g / mm3) and taking the reciprocal (mm3 / g)

[0118] The number of precipitates of 5 to 50 nm observed in the cross-section of 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 μm2 was counted.

[0119] Also, the hardness value was measured at three points of the surface layer portion using a Brinell hardness tester and the average value 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 was used.

[0120] Also, the surface of each sample was visually observed, grinding was performed at the point at which surface cracks were formed, and the grinding length until the cracks disappeared was measured as the surface crack length.TABLE 3PropertiesMicrostructureImpactSurfaceCentralPorosityPrecip-tough-layerportiondiffer-itateshard-TensileElonga-nessCrackSampleSteelportion(area % )ence(pernessstrengthtion(−50° C.,DepthNo.type(area %)BF(mm3 / g)μm2)(HB)(MPa)(%)J)(μm)Classification 1AP.F9280.00224202514723150Inventiveexample 1 2BP.F89110.00229215534593020Inventiveexample 2 3CP.F88120.01325207524642980Inventiveexample 3 4DP.F76240.01536205508643000Inventiveexample 4 5EP.F89110.00719203517593150Inventiveexample 5 6AP.F9190.012342875266630719Comparativeexample 1 7BP.F84160.36033207500123200Comparativeexample 2 8BP.F76240.017452765435925716Comparativeexample 3 9CP.F88120.015272745296128918Comparativeexample 410CP.F86140.0071928369321149Comparativeexample 511DM84160.006182745345428815Comparativeexample 612DP.F22780.00723206433573040Comparativeexample 713EP.F81190.015 418246164130Comparativeexample 814FM10000.01741215798 81217Comparativeexample 915GP.F9910.01535208434302890Comparativeexample 1016HP.F84160.009 1210451613010Comparativeexample 1117IB9820.0118921468755258Comparativeexample 1218JP.F30700.01638211431572860Comparativeexample 13* P.F: polygonal ferrite, M: martensite, B: bainite

[0121] As listed in Table 3, in the case of the inventive example satisfying the alloy composition and manufacturing conditions of the present disclosure, the microstructure characteristics suggested in the present disclosure were satisfied, and the properties targeted in the present disclosure was also ensured.

[0122] FIG. 1 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of inventive example 1 according to an embodiment of the present disclosure. In inventive example 1, as indicated in FIG. 1, polygonal ferrite was formed on the surface layer portion, and the ¼ point was formed as bainite, which was advantageous for ensuring cold formability and strength.

[0123] As for comparative examples 1 and 2, the cumulative reduction ratio exceeded the range of the present disclosure during forging. In comparative example 1, the reduction ratio was excessive during forging, such that forging cracks were created, and surface quality was poor during cold bending. Also, the surface layer portion hardness value was excessive. In comparative example 2, the reduction ratio was insufficient during forging, such that the retained voids were not sufficiently compressed at the center of the steel sheet, resulting in degradation of elongation.

[0124] As for comparative example 3, the reduction ratio exceeded the range of the present disclosure at the recrystallization temperature or lower. Accordingly, surface defects occurred.

[0125] As for comparative example 4, the strain rate during forging exceeded the range suggested in the present disclosure. Accordingly, the surface layer portion hardness value was excessive, and cracks also occurred.

[0126] As for comparative example 5, the finishing-rolling temperature was not reached during hot-rolling. Accordingly, strength was excessively high, low-temperature impact toughness was deteriorated, and cold formability was also deteriorated.

[0127] As for comparative example 6, the cooling rate exceeded the range suggested in the present disclosure during primary-cooling. Accordingly, martensite, a hard structure, was formed in the surface layer, which reduced cold formability, and surface defects occurred in the product during the cold bending process.

[0128] FIG. 2 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of comparative example 6, which does not satisfy an embodiment of the present disclosure. As indicated in FIG. 2, it may be confirmed that the surface layer portion was formed as a martensite structure.

[0129] As for comparative example 7, the cooling rate was not satisfied during secondary-cooling, and adequate strength was not ensured in the central portion, such that strength targeted in the present disclosure was not ensured.

[0130] As for comparative example 8, the tempering temperature was excessively high, and the formation of precipitates was not smooth, and accordingly, strength targeted in the present disclosure was not ensured.

[0131] As for comparative examples 9 and 10, the carbon content exceeded the range suggested in the present disclosure. Accordingly, in comparative example 9, martensite, a hard structure, was formed in the surface layer portion, such that it was difficult to ensure the target impact toughness. As for comparative example 10, the Mn content was also deteriorated, such that it was difficult to ensure appropriate strength in the central portion, resulting in degradation of strength.

[0132] As for comparative example 11, the Nb content did not satisfy the content suggested in the present disclosure. Accordingly, the formation of precipitates was not easy, and strength targeted was not ensured.

[0133] As for comparative example 12, the V content exceeded the range suggested in the present disclosure. Accordingly, the bainite structure was formed in the surface layer portion, and accordingly, low-temperature impact toughness was deteriorated.

[0134] As for comparative example 13, the C content did not satisfy the content suggested in the present disclosure. Accordingly, bainite was not properly formed in the central portion, and strength was also degraded.

[0135] The present disclosure has been described in detail through embodiments above, and other forms of embodiments may also be possible. Thus, the technical spirit and scope of the claims described below are not limited to the embodiments.

Examples

embodiment

[0115]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 steel sheet having a thickness of 133 mm was finally manufactured. In this case, as for the primary-reheating temperature, 1200° C. was commonly applied, as for the secondary-reheating temperature, 1100° C. was commonly applied, and as for the reheating time for primary-cooling and tempering time, 30 minutes was commonly applied.

TABLE 1SteelAlloy composition (weight %)typeCSiMnAlPSNbVTiCrMoCuNiCaA0.160.351.320.03090100.01700.0070.0110.060.090.050.1919B0.140.311.430.02782110.01100.0090.0070.080.080.070.2322C0.130.321.350.03177100.01800.0120.0050.120.100.050.2118D0.160.351.220.0288490.01700.0230.0100.10.090.100.2218E0.150.401.210.0308990.01500.0190.0110.090.060.080.1917F0.280.351.210.0307480.01900.0090.0090.120.080.090.2518G0.030.310.30.03080100.01300...

TECHNICAL FIELD

[0001] The present disclosure relates to a steel sheet and a manufacturing method therefor, and more particularly, to a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamella tearing quality, and a method of manufacturing the same.BACKGROUND ART

[0002] Recently, due to enlargement of crude oil refining and storage facilities and demand for high-capacity storage, demand for increasing a thickness of steel used in the facility has been continuously increased. In manufacturing a large structure, to improve the internal and external soundness of steel, a concentration of impurities such as non-metallic inclusions or segregation are being reduced, or cracks and voids on the surface and inside the material are controlled to the extreme.

[0003] In particular, in the case of an extremely thick material having a thickness exceeding 100 mm, a rolling reduction ratio may not be high as compared to a thin material, such that unsolidified shrinkage occurring during continuous casting or casting may not be sufficiently compressed during a rough rolling process and may remain in the form of retained voids in a central portion of a product. The retained voids may act as an initiation point of cracks when the structure receives stress in a thickness axial direction, and may eventually cause damage to the entire facility in the form of lamellar tearing. Accordingly, a process to sufficiently compress central voids such that no retained voids are present may be necessary in the step before rolling.

[0004] Cited document 1, related thereto, may disclose a steel reduction technique in a thick steel rough rolling process, and may use a technique of determining a limit reduction ratio for different thickness ratios at which sheet sticking occurs from the steel reduction ratio for each pass determined to be close to a design allowable value (load and torque) of the rolling mill, a technique of distributing a reduction ratio by adjusting an index of thickness ratio for each pass to ensure a target thickness of a rough rolling machine, a technique of modifying a reduction ratio to prevent sheet sticking based on the limit reduction ratio for different thicknesses, and a manufacturing method for allowing an average reduction ratio of approximately 27.5% in the final 3 passes of rough rolling based on a thickness of 80 mm may be provided. However, in the case of the above rolling method, the average reduction ratio of the entire product thickness may be measured, such that it may be difficult to apply high strain to a central portion of an extremely thick material having a maximum thickness of 233 mm in which retained voids are present.

[0005] One of the other methods for manufacturing an extremely thick material may be a method of using a forging machine having a higher effective strain per pass than a rolling mill. Cited document 2 is a method of vertically erecting the continuous casting slab extracted from the furnace and providing a full width forging reduction of 400 mm or more, eliminating pores in the edge portion in the width direction and a central portion by performing a forging pass with a reduction of less than 2 passes, which is within the buckling limit reduction ratio, and increasing a central portion strain rate. The retained void in the central portion, which was a problem in cited document 1, may be effectively compressed, lamellar tearing quality of the product may be improved when applied to forging.

[0006] However, there may be a disadvantage in that surface defects may occur due to local strain concentration during a width forging process. In particular, when a surface or subsurface defect is present in the cast state before forging, the defect may propagate during the forging process and the surface quality may become inferior in the product state after rolling.

[0007] Cited document 3 discloses that a thick-walled, high-strength steel sheet having a yield strength of 620 MPa or more and a thickness of 100 mmt or more may be manufactured through a process of heating a material provided with a predetermined alloy composition at 1200 to 1350° C., performing hot forging with a cumulative reduction of 25% or more, heating at Ac3 point or more and 1200° C. higher, performing hot rolling with a cumulative reduction of 40% or more, reheating to Ac3 point or more and 1050° C. or higher, rapidly cooling from the temperature of Ac3 point or more to a lower temperature of 350° C. or less or Ar3 point or less, and tempering at a temperature of 450° C. to 700° C. However, in the case of the ultra-high strength steel sheet described above, the carbon equivalent (Ceq) and hardenability index may be high, such that the sheet may be vulnerable to surface cracks during casting, and rapid cooling in the austenite single phase region may locally form rapid cooling structures of bainite or martensite on the surface layer of the sheet, and the sheet may not be used for a storage tank and a pressure vessel having excellent cold formability.PRIOR ARTCited Document

[0008] (Cited document 1) Korean Laid-Open Patent Publication No. 10-2012-0075246 (publicized on Jul. 6, 2012)

[0009] (Cited document 2) Korean Laid-Open Patent Publication No. 10-2012-0074039 (publicized on Jul. 5, 2012)

[0010] (Cited document 3) Korean Laid-Open Patent Publication No. 10-2017-7019203 (publicized on Jan. 15, 2016)DETAILED DESCRIPTION OF PRESENT DISCLOSURETechnical Problems to Solve

[0011] An embodiment of the present disclosure is to provide a steel sheet and a method of manufacturing the same.

[0012] An embodiment of the present disclosure is to provide a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamella tearing quality, and a method of manufacturing the same.

[0013] The subject matter of the present disclosure is not limited to the above-described contents. Those skilled in the art will have no difficulty in understanding additional purposes of the present disclosure from the overall contents of this specification.Solution to Problem

[0014] According to an embodiment of the present disclosure, provided is a steel sheet including, by weight %, carbon (C): 0.10 to 0.25%, 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): 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V): 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities,

[0015] wherein, in a microstructure, a surface layer portion, which is a region from a surface to a ⅛ point in a thickness direction, includes 80% or more of polygonal ferrite by area %, and a central portion, which is the other region from a ⅛ point to a ½ point in the thickness direction, includes 50% or more of bainite and a remainder of ferrite by area %,

[0016] wherein a difference between a porosity of a region from a surface to a ⅜ point in the thickness direction and a porosity of a region from a ⅜ point to a ⅝ point in the thickness direction is 0.100 mm3 / g or less, and

[0017] wherein a hardness value of the surface layer portion is 200 to 220 HB.

[0018] The surface layer portion may include one or more of pearlite and bainite as a remainder structure.

[0019] In the steel sheet, one or more of fine NbC, NbCN, VC, CVN precipitates having a diameter of 5 to 50 nm may be present in a density of 10 or more per 1 μm2.

[0020] The steel sheet may have a tensile strength of 510 to 690 MPa, an elongation (ZRA) in the thickness direction of 35% or more, and a Charpy impact absorption energy value at −50° C. of 50 J or more.

[0021] The steel sheet may have a maximum surface crack depth of 1 μm or less in a 180° bending test at room temperature.

[0022] The steel sheet may have a thickness of 133 to 233 mm.

[0023] According to an embodiment of the present disclosure, provided is a method of manufacturing a steel sheet including primary-reheating a steel slab including, by weight %, carbon (C): 0.10 to 0.25%, 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) 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V) 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities;

[0024] forging the primary-reheated steel slab at a cumulative reduction ratio of 35 to 65% and a strain rate of 1.0 to 4.0 / s;

[0025] secondary-reheating the forged steel slab;

[0026] hot-rolling the secondary-reheated steel slab at a finishing-rolling temperature of 900 to 1100° C.;

[0027] heating the hot-rolled steel sheet in a temperature range of 820 to 900° C., holding for 10 to 40 minutes, and primary-cooling the steel sheet to 700° C. at an average cooling rate of 0.1 to 5.0° C. / s based on a steel sheet surface temperature;

[0028] secondary-cooling the primary-cooled steel sheet to room temperature at an average cooling rate of 10.0° C. / s or higher based on the steel sheet surface temperature; and

[0029] performing a tempering heat treatment of heating the secondary-cooled steel sheet in a temperature range of 550 to 700° C. and holding for 5 to 60 minutes,

[0030] wherein, in the forging, a cumulative reduction ratio at a recrystallization temperature or lower is 20% or less.

[0031] The primary-reheating may be performed in a temperature range of 1100 to 1300° C., and

[0032] The secondary-reheating may be performed in a temperature range of 1000 to 1200° C.

[0033] In the primary-reheating, a thickness of the steel slab may be 650 to 750 mm,

[0034] after the forging, a thickness of the steel slab may be 350 to 450 mm, and

[0035] after the hot-rolling, a thickness of the steel sheet may be 133 to 233 mm.Advantageous Effects of Invention

[0036] According to an aspect of the present disclosure, a steel sheet and a method of manufacturing the same may be provided.

[0037] According to an aspect of the present disclosure, a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamella tearing quality, and a method of manufacturing the same may be provided.

[0038] According to an aspect of the present disclosure, a steel sheet for an ultra-thick pressure vessel having excellent cold formability and lamellar tearing quality which may be used in a petrochemical manufacturing facility, a storage tank, or the like.BRIEF DESCRIPTION OF DRAWINGS

[0039] FIG. 1 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of inventive example 1 according to an embodiment of the present disclosure.

[0040] FIG. 2 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of comparative example 6, which falls outside the scope of one embodiment of the present disclosure.BEST MODE FOR INVENTION

[0041] Hereinafter, preferred embodiments of the present invention may be described. The implementation examples of the present disclosure may be modified in various forms, and the scope of the present disclosure should not be construed as being limited to the implementation examples described below. The implementation examples are provided to describe the present disclosure in greater detail to those skilled in the art to which the present invention pertains.

[0042] In the description below, the present disclosure will be described in detail.

[0043] In the description below, a composition of steel of the present disclosure will be described in detail.

[0044] Unless otherwise indicated, the percentages indicating the content of each element may be based on weight.

[0045] A steel sheet according to an embodiment of the present disclosure may include, by weight %, carbon (C): 0.10 to 0.25%, 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): 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V): 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities.

[0046] Carbon (C): 0.10 to 0.25%

[0047] Carbon (C) may be the most important element for ensuring basic strength, and thus, it may be necessary to be contained in steel within an appropriate range. To obtain the effect of the addition, 0.10% or more of carbon (C) may be added. According to an embodiment of the present disclosure, 0.12% or more of carbon (C) may be added. When the content exceeds a predetermined level, a fraction of low-temperature transformation structure may increase during QT heat treatment, and strength and hardness of the base material may be excessive, and accordingly, uniform elongation may be degraded due to the formation of a low-temperature structure rather than an elongated ferrite as a surface structure. Accordingly, in the present disclosure, an upper limit of carbon (C) content may be limited to 0.25%. According to an embodiment of the present disclosure, an upper limit thereof may be 0.20%.

[0048] Silicon (Si): 0.05 to 0.50%

[0049] Silicon (Si) may be a substitutional element, and may improve strength of a steel material through solid solution strengthening 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 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%.

[0050] Manganese (Mn): 1.00 to 2.00%

[0051] Manganese (Mn) may be a useful element for improving strength by solid solution strengthening and improving hardenability by generating a low-temperature transformation phase. Accordingly, 1.00% or more of manganese (Mn) may be added to ensure a tensile strength of 450 MPa 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 with S, which is an elongated non-metallic inclusion, which may degrade toughness, and when tensioned in the thickness direction, MnS may act as a factor degrading an elongation, and thus, MnS may be a factor rapidly degrading lamellar tearing quality. 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.

[0052] Aluminum (Al): 0.005 to 0.100%

[0053] 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. According to an embodiment of the present disclosure, a lower limit of the aluminum (Al) content may be 0.01%. When the aluminum (Al) content is excessive, a fraction of Al2O3 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 lamellar tearing properties. 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.

[0054] Phosphorus (P): 0.0100% or less

[0055] Phosphorus (P) may cause brittleness in the grain system or may form coarse inclusions, thereby causing brittleness. Thus, to improve resistance to brittle crack propagation, phosphorus (P) may be limited to 0.0100% or less. However, considering the level inevitably contained during steel manufacturing, 0% may be excluded.

[0056] Sulfur (S): 0.0015% or less

[0057] Sulfur (S) may cause brittleness in the grain system or forming coarse inclusions, thereby causing brittleness. Thus, to improve resistance to brittle crack propagation, sulfur (S) may be limited to 0.0015% or less. However, considering the level inevitably contained during steel manufacturing, 0% may be excluded.

[0058] Niobium (Nb): 0.001 to 0.030%

[0059] Niobium (Nb) may be precipitated in the form of NbC or NbCN and may enhance strength of a base material. Also, when reheating at a high temperature, dissolved Nb may be relatively finely precipitated in the form of NbC during rolling, which may suppress recrystallization of austenite and may have the effect of refining the structure. Accordingly, niobium (Nb) may be added at 0.001% or more. According to an embodiment of the present disclosure, niobium (Nb) may be 0.005% or more. When niobium (Nb) is excessively added, undissolved niobium (Nb) may be generated in the form of TiNb(C,N), which may hinder lamellar tearing properties, and thus, an upper limit of the content may be limited to 0.030%. According to an embodiment of the present disclosure, the niobium (Nb) content may be 0.020% or less.

[0060] Vanadium (V): 0.001 to 0.030%

[0061] 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 vanadium (V) may have the effect of improving strength by precipitating fine carbonitrides during a subsequent heat treatment process such as PWHT. To sufficiently obtain this effect, vanadium (V) may be added at 0.001% or more. According to an embodiment of the present disclosure, a lower limit of vanadium (V) content may be 0.010%. When the content is excessive, vanadium (V) may excessively increase strength and hardness of a base material and a welded zone, which may act as a factor for surface cracks during pressure vessel processing, and manufacturing costs may rapidly increase, which may not be commercially advantageous. Accordingly, the vanadium (V) content may be 0.030% or less. According to an embodiment of the present disclosure, vanadium (V) content may be 0.020% or less.

[0062] Titanium (Ti): 0.001 to 0.030%

[0063] Titanium (Ti) may be a component significantly improving low-temperature toughness by precipitating as TiN during reheating and suppressing grain growth in the base material and the weld heat-affected zone, and thus, titanium (Ti) may be added at 0.001% or more to obtain the effect of the addition. When titanium (Ti) is excessively added, it may cause clogging of the continuous casting nozzle or segregation of a central portion, leading to a reduction in low-temperature toughness. In addition, 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, lamellar tearing properties of the final product may be degraded. Accordingly, the titanium (Ti) content may be 0.030% or less. According to an embodiment of the present disclosure, the titanium (Ti) content may be 0.025% or less, and may be 0.018% or less.

[0064] Chromium (Cr): 0.01 to 0.20%

[0065] Chromium (Cr) may increase hardenability and may form a low-temperature transformation structure, thereby increasing yield and tensile strength, and chromium (Cr) may slow down the decomposition rate of cementite during tempering after rapid cooling or heat treatment after welding, thereby preventing the decrease in strength. For this effect, 0.01% or more of chromium (Cr) may be added. When the chromium (Cr) content is excessive, the size and fraction of Cr-rich coarse carbides such as M23C6 may increase, thereby impact toughness of the product may degrade. In addition, 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. Thus, an upper limit of the content may be limited to 0.20%. According to an embodiment of the present disclosure, an upper limit of the chromium (Cr) content may be 0.15%.

[0066] Molybdenum (Mo): 0.01 to 0.15%

[0067] Molybdenum (Mo) may increase grain boundary strength and may have a relatively large solid solution strengthening effect in ferrite, 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 0.15%.

[0068] Copper (Cu): 0.01 to 0.50%

[0069] Copper (Cu) may greatly improve strength of a matrix phase by solid solution strengthening in ferrite, and may have the effect of inhibiting corrosion in a wet hydrogen sulfide atmosphere, and thus, copper (Cu) may be a favorable element in the present disclosure. 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 0.50%. According to an embodiment of the present disclosure, an upper limit thereof may be 0.30%.

[0070] Nickel (Ni): 0.05 to 0.50%

[0071] Nickel (Ni) may increase stacking faults at low temperature and may facilitate cross slip of dislocations, thereby improving impact toughness and hardenability and improving strength. 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 0.50%. According to an embodiment of the present disclosure, an upper limit of the nickel (Ni) content may be 0.30%.

[0072] Calcium (Ca): 0.0005 to 0.0040%

[0073] When calcium (Ca) added after deoxidation by Al, calcium (Ca) may combine with S, which forms MnS inclusions, such that the formation of MnS may be suppressed. In addition, calcium (Ca) may form spherical CaS thereby cracks due to hydrogen-induced cracking may be suppressed. To sufficiently form S contained as impurities into CaS, calcium (Ca) may be added at 0.0005% or more. However, when the amount added is excessive, 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 rolling, such that lamellar tearing properties may be degraded, and thus, an upper limit of calcium (Ca) content may be limited to 0.0040%.

[0074] The steel material of the present disclosure may further include iron (Fe) other than the above-described composition. In a general manufacturing process, 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.

[0075] Hereinafter, the steel microstructure of the present disclosure will be described in detail.

[0076] Unless otherwise indicated in the present disclosure, “%” indicating a fraction of the microstructure may be based on an area.

[0077] According to an embodiment of the present disclosure, as for a microstructure of the steel sheet, by area %, a surface layer portion, which is a region from a surface to a ⅛ point in a thickness direction, may include 80% or more of polygonal ferrite as area %, and a central portion, which is a region from a ⅛ point to a ½ point in the thickness direction, may include 50% or more of bainite and a remainder of ferrite.

[0078] In the present disclosure, to improve cold workability through the formation of a soft microstructure of the surface layer, the surface layer portion, which is a region from the surface to the ⅛ point in the thickness direction, may be limited to include 80% or more of polygonal ferrite. In an embodiment of the present disclosure, polygonal ferrite may be 90% or more. In addition to polygonal ferrite, the remainder structure may include one or more of pearlite and bainite.

[0079] The central portion, which is the region from the ⅛ point to the ½ point, which is the other region, may include 50% or more of bainite to ensure strength suggested in the present disclosure, and may include ferrite as the remainder structure.

[0080] According to an embodiment of the present disclosure, a difference between a porosity of the region from the surface to a ⅜ point in the thickness direction and a porosity of the region from the ⅜ point to a ⅝ point in the thickness direction may be 0.100 mm3 / g or less.

[0081] That is, the value obtained by subtracting the porosity of the region from the surface to the ⅜ point in the thickness direction from the porosity of the region from the ⅜ point to the ⅝ point in the thickness direction of the steel sheet may be 0.100 mm3 / g or less. Here, the porosity may be obtained by measuring a density (g / mm3) and taking a reciprocal (mm3 / g).

[0082] In the present disclosure, the difference in porosities between the surface layer portion and the central portion may be limited for lamellar tearing properties. When the difference in porosities exceeds 0.100 mm3 / g, the retained void may act as a crack initiation point degrading the elongation during the thickness direction tensile test, such that the elongation (ZRA) in the thickness direction may be deteriorated. According to an embodiment of the present disclosure, the difference may be 0.050 mm3 / g or less. According to an embodiment of the present disclosure, the difference may be 0.030 mm3 / g or less. It may not be necessary to specifically limit a lower limit of the difference, but in an embodiment, a lower limit may be 0 mm3 / g.

[0083] According to an embodiment of the present disclosure, one or more of fine NbC, NbCN, VC, and CVN precipitates having a diameter of 5 to 50 nm may be present in a density of 10 or more per 1 μm2.

[0084] In the present disclosure, the precipitates may be limited to ensure strength through precipitation strengthening. When the number of precipitates is less than 10 per 1 μm2, the strengthening effect may be insufficient, such that it may be difficult to ensure appropriate strength suggested in the present disclosure. When the diameter of the precipitate is excessively small, less than 5 nm, it may be difficult to ensure strength similarly to the effect of having a small number of precipitates, and when the diameter exceeds 50 nm, the pinning effect may also be reduced, such that the strength improvement effect may be reduced. Also, when the size of precipitate exceeds 50 nm, coarse precipitates may act as brittle fracture initiation points during an impact test, and thus, the diameter of the precipitate may be 5 to 50 nm appropriately.

[0085] According to an embodiment of the present disclosure, a steel sheet may have a thickness of 133 to 233 mm, a hardness value of the surface layer portion of 200 to 220 HB, a tensile strength of 510 to 690 MPa, an elongation in the thickness direction (ZRA) of 35% or more, a Charpy impact absorption energy value at −50° C. of 50 J or more, and a maximum surface crack depth of 1 μm or less in a 180° bending test at room temperature.

[0086] Hereinafter, the method of manufacturing steel of the present disclosure will be described in detail.

[0087] 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.Primary-Reheating

[0088] 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.

[0089] The primary-reheating may be performed in a temperature range of 1100° C. or higher to re-dissolve Ti or Nb composite carbonitride or TiNb(C,N) coarse crystallites formed during casting, to homogenize the structure by heating and holding 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. When the steel slab is reheated at an excessively high temperature, problems may occur due to the oxidation scale at high temperature, and the manufacturing costs may increase excessively due to the increased cost due to heating and holding, and thus, an upper limit of the steel slab heating temperature may be limited to 1300° C.

[0090] According to an embodiment of the present disclosure, 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.Forging

[0091] 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.

[0092] The forging may be of processing the heated steel slab into the shape of a final desired 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.

[0093] In the present disclosure, the strain rate may indicate the strain rate per unit time, and the unit may indicate % / s.

[0094] When the cumulative reduction ratio of forging is less than 35%, the void retained in the steel slab may not be sufficiently compressed, such that retained voids may be formed, which may degrade lamellar tearing properties 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%, resulting in an increased number of forging passes, the surface temperature may continue to decrease, which may cause surface cracks at low temperature.

[0095] When the cumulative reduction ratio exceeds 20% at the non-recrystallization temperature or lower, at which dislocation density is recovered or not offset by recrystallization, the uniform elongation of the steel surface layer during forging at approximately 850° C. to 1150° C. exceeds 20%. As a result, due to work hardening caused by accumulated dislocations, the surface uniform elongation significantly decreases, 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.

[0096] When the strain rate is less than 1.0 / s during forging, the forging productivity may be degraded, and since the temperature of the surface layer portion is excessively reduced during the forging process, surface cracks may occur during the finishing forging process. When a strain rate thereof exceeds 4.0 / s, surface quality may be degraded due to excessive work hardening.

[0097] According to an embodiment of the present disclosure, the thickness of the steel slab after forging may be 350 to 450 mm.Secondary-Reheating

[0098] The forged steel slab may be secondary-reheated in a temperature range of 1000 to 1200° C.

[0099] To re-dissolve the composite carbonitride of Ti or Nb or the coarse-grained TiNb(C,N) formed during casting, to homogenize the structure by heating and holding the austenite to the recrystallization temperature or higher before hot-rolling, and to minimize crushing of inclusions during the rolling process by ensuring the rolling end temperature to be sufficiently high, reheating may be performed in the temperature range of 1000° C. or higher. When the slab is heated at an excessively high temperature, a problem may occur due to the oxidation scale at high temperature, and the manufacturing cost may increase excessively due to the increased cost due to heating and holding, and thus, an upper limit of the secondary heating temperature may be limited to 1200° C.Hot-Rolling

[0100] The secondary-reheated steel slab may be hot-rolled at a finishing-rolling temperature of 900 to 1100° C.

[0101] When the finishing-rolling temperature is lower than 900° C. during hot-rolling, the strain resistance value may increase excessively as the temperature decreases, such that it may be difficult to sufficiently refine austenite grains in the central portion of the product in the thickness direction, and accordingly, lamellar tearing properties of the final product may deteriorate. When the finishing-rolling temperature exceeds 1100° C., the austenite grains may become excessively coarse, such that strength and impact toughness may deteriorate.

[0102] According to an embodiment of the present disclosure, the thickness of the steel sheet after hot-rolling may be 133 to 233 mm.Primary-Cooling

[0103] 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.

[0104] In the present disclosure, 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 ⅛ point in the thickness direction. During the primary-cooling, the temperature and cooling rate may be based on the surface temperature of the steel sheet.

[0105] When the reheating, when the temperature is less than 820° C. or the holding time is less than 10 minutes, carbides generated during cooling after rolling or impurities segregated at the grain boundaries may not be re-dissolved smoothly, such that the thickness direction elongation (ZRA) and low-temperature toughness of the steel material after heat treatment may be significantly degraded. When the temperature exceeds 900° C. or the holding time exceeds 40 minutes, lamellar tearing quality may be degraded due to coarsening of austenite and coarsening of precipitated phases such as Nb(C,N), V(C,N).

[0106] When the average cooling rate is less than 0.1° C. / s during primary-cooling, the waiting time in the air may be excessively long, such that the manufacturing time may increase, which may degrade productivity. When 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 uniform elongation of the surface layer of the final product may be degraded, and thus, cold workability may be rapidly deteriorated.Secondary-Cooling

[0107] 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.

[0108] The secondary-cooling may be performed to include the microstructure of the region other than the surface layer portion into bainite or dual-phase structure of bainite and ferrite through the rapid cooling.

[0109] During the secondary-cooling, the average cooling rate may be based on the surface temperature of the steel sheet. When the average cooling rate is less than 10.0° C. / s, it may be difficult to obtain the low-temperature transformation structure described above. In the present disclosure, 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 quenching, for example, by slowing down the passing speed of the steel material and increasing the flow rate of the sprayed water, or the like.Tempering Heat Treatment

[0110] 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.

[0111] 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.

[0112] When the 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. When the 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.

[0113] When the 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. When the tempering heat treatment time exceeds 60 minutes, cementite may become spheroidized due to excessive heating, which may rapidly reduce strength.MODE FOR INVENTION

[0114] Hereinafter, the present disclosure may be described more specifically through embodiments. However, it should be noted that the embodiments below are merely intended to describe the present disclosure in greater detail based on embodiments, and are not intended to limit the scope of the rights of the present disclosure.Embodiment

[0115] 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 steel sheet having a thickness of 133 mm was finally manufactured. In this case, as for the primary-reheating temperature, 1200° C. was commonly applied, as for the secondary-reheating temperature, 1100° C. was commonly applied, and as for the reheating time for primary-cooling and tempering time, 30 minutes was commonly applied.TABLE 1SteelAlloy composition (weight %)typeCSiMnAlPSNbVTiCrMoCuNiCaA0.160.351.320.03090100.01700.0070.0110.060.090.050.1919B0.140.311.430.02782110.01100.0090.0070.080.080.070.2322C0.130.321.350.03177100.01800.0120.0050.120.100.050.2118D0.160.351.220.0288490.01700.0230.0100.10.090.100.2218E0.150.401.210.0308990.01500.0190.0110.090.060.080.1917F0.280.351.210.0307480.01900.0090.0090.120.080.090.2518G0.030.310.30.03080100.01300.0090.0050.20.080.070.1919H0.160.352.350.0307990.00050.0140.0040.250.090.060.1818I0.170.301.210.02982350.01200.1500.0050.230.100.100.1920J0.060.251.320.31085100.01500.0160.0120.150.070.060.232* Units for P, S, Ca and O may be ppm.TABLE 2ForgingReductionratio atHot-recrystal-rollingTemper-CumulativelizationFinishingPrimary-Secondary-ingreduction temperatureStrainrollingcoolingcoolingTemper-SampleSteelratioor lowerratetemperatureRateRateatureNo.type(%)(%)( / s)(° C.)(° C. / s)(° C. / s)(° C.) 1A45151.79151.324.5620 2B47132.59300.916.6618 3C51113.910501.618.4625 4D61142.79872.426.5636 5E39 92.49683.629.3641 6A84101.89154.128.0632 7B12131.99214.327.7629 8B48392.69253.816.6621 9C47148.09641.519.461810C51163.37101.826.559911D50132.798183.034.458012D59153.59900.77.466413E51143.49801.527.378814E49173.89121.536.063615G39123.910433.247.062916H42112.910542.640.062717I46171.89833.035.064018J50141.69912.733.0633The microstructure and physical properties of each manufactured sample were measured and listed in Table 3 below. The microstructure was measured by collecting samples from the surface layer region (the region from the surface to a ⅛ point in the thickness direction) and the other region (the region from the ⅛ point to the ½ point in the thickness direction) using an image auto-analyzer. In this case, the observed microstructures were polygonal ferrite (P.F), martensite (M), and bainite (B), and area % thereof may be 100%.

[0117] The porosity difference was represented as the value obtained by subtracting the porosity of the region from the surface to the ⅜ point in the thickness direction from the porosity of the region from the ⅜ point to the ⅝ point in the thickness direction in the present disclosure. Here, the porosity was obtained by measuring the density (g / mm3) and taking the reciprocal (mm3 / g)

[0118] The number of precipitates of 5 to 50 nm observed in the cross-section of 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 μm2 was counted.

[0119] Also, the hardness value was measured at three points of the surface layer portion using a Brinell hardness tester and the average value 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 was used.

[0120] Also, the surface of each sample was visually observed, grinding was performed at the point at which surface cracks were formed, and the grinding length until the cracks disappeared was measured as the surface crack length.TABLE 3PropertiesMicrostructureImpactSurfaceCentralPorosityPrecip-tough-layerportiondiffer-itateshard-TensileElonga-nessCrackSampleSteelportion(area % )ence(pernessstrengthtion(−50° C.,DepthNo.type(area %)BF(mm3 / g)μm2)(HB)(MPa)(%)J)(μm)Classification 1AP.F9280.00224202514723150Inventiveexample 1 2BP.F89110.00229215534593020Inventiveexample 2 3CP.F88120.01325207524642980Inventiveexample 3 4DP.F76240.01536205508643000Inventiveexample 4 5EP.F89110.00719203517593150Inventiveexample 5 6AP.F9190.012342875266630719Comparativeexample 1 7BP.F84160.36033207500123200Comparativeexample 2 8BP.F76240.017452765435925716Comparativeexample 3 9CP.F88120.015272745296128918Comparativeexample 410CP.F86140.0071928369321149Comparativeexample 511DM84160.006182745345428815Comparativeexample 612DP.F22780.00723206433573040Comparativeexample 713EP.F81190.015 418246164130Comparativeexample 814FM10000.01741215798 81217Comparativeexample 915GP.F9910.01535208434302890Comparativeexample 1016HP.F84160.009 1210451613010Comparativeexample 1117IB9820.0118921468755258Comparativeexample 1218JP.F30700.01638211431572860Comparativeexample 13* P.F: polygonal ferrite, M: martensite, B: bainite

[0121] As listed in Table 3, in the case of the inventive example satisfying the alloy composition and manufacturing conditions of the present disclosure, the microstructure characteristics suggested in the present disclosure were satisfied, and the properties targeted in the present disclosure was also ensured.

[0122] FIG. 1 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of inventive example 1 according to an embodiment of the present disclosure. In inventive example 1, as indicated in FIG. 1, polygonal ferrite was formed on the surface layer portion, and the ¼ point was formed as bainite, which was advantageous for ensuring cold formability and strength.

[0123] As for comparative examples 1 and 2, the cumulative reduction ratio exceeded the range of the present disclosure during forging. In comparative example 1, the reduction ratio was excessive during forging, such that forging cracks were created, and surface quality was poor during cold bending. Also, the surface layer portion hardness value was excessive. In comparative example 2, the reduction ratio was insufficient during forging, such that the retained voids were not sufficiently compressed at the center of the steel sheet, resulting in degradation of elongation.

[0124] As for comparative example 3, the reduction ratio exceeded the range of the present disclosure at the recrystallization temperature or lower. Accordingly, surface defects occurred.

[0125] As for comparative example 4, the strain rate during forging exceeded the range suggested in the present disclosure. Accordingly, the surface layer portion hardness value was excessive, and cracks also occurred.

[0126] As for comparative example 5, the finishing-rolling temperature was not reached during hot-rolling. Accordingly, strength was excessively high, low-temperature impact toughness was deteriorated, and cold formability was also deteriorated.

[0127] As for comparative example 6, the cooling rate exceeded the range suggested in the present disclosure during primary-cooling. Accordingly, martensite, a hard structure, was formed in the surface layer, which reduced cold formability, and surface defects occurred in the product during the cold bending process.

[0128] FIG. 2 is an image of a microstructure of a surface layer and a ¼ point in a thickness direction of comparative example 6, which does not satisfy an embodiment of the present disclosure. As indicated in FIG. 2, it may be confirmed that the surface layer portion was formed as a martensite structure.

[0129] As for comparative example 7, the cooling rate was not satisfied during secondary-cooling, and adequate strength was not ensured in the central portion, such that strength targeted in the present disclosure was not ensured.

[0130] As for comparative example 8, the tempering temperature was excessively high, and the formation of precipitates was not smooth, and accordingly, strength targeted in the present disclosure was not ensured.

[0131] As for comparative examples 9 and 10, the carbon content exceeded the range suggested in the present disclosure. Accordingly, in comparative example 9, martensite, a hard structure, was formed in the surface layer portion, such that it was difficult to ensure the target impact toughness. As for comparative example 10, the Mn content was also deteriorated, such that it was difficult to ensure appropriate strength in the central portion, resulting in degradation of strength.

[0132] As for comparative example 11, the Nb content did not satisfy the content suggested in the present disclosure. Accordingly, the formation of precipitates was not easy, and strength targeted was not ensured.

[0133] As for comparative example 12, the V content exceeded the range suggested in the present disclosure. Accordingly, the bainite structure was formed in the surface layer portion, and accordingly, low-temperature impact toughness was deteriorated.

[0134] As for comparative example 13, the C content did not satisfy the content suggested in the present disclosure. Accordingly, bainite was not properly formed in the central portion, and strength was also degraded.

[0135] The present disclosure has been described in detail through embodiments above, and other forms of embodiments may also be possible. Thus, the technical spirit and scope of the claims described below are not limited to the embodiments.

Examples

embodiment

[0115]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 steel sheet having a thickness of 133 mm was finally manufactured. In this case, as for the primary-reheating temperature, 1200° C. was commonly applied, as for the secondary-reheating temperature, 1100° C. was commonly applied, and as for the reheating time for primary-cooling and tempering time, 30 minutes was commonly applied.

TABLE 1SteelAlloy composition (weight %)typeCSiMnAlPSNbVTiCrMoCuNiCaA0.160.351.320.03090100.01700.0070.0110.060.090.050.1919B0.140.311.430.02782110.01100.0090.0070.080.080.070.2322C0.130.321.350.03177100.01800.0120.0050.120.100.050.2118D0.160.351.220.0288490.01700.0230.0100.10.090.100.2218E0.150.401.210.0308990.01500.0190.0110.090.060.080.1917F0.280.351.210.0307480.01900.0090.0090.120.080.090.2518G0.030.310.30.03080100.01300...

Claims

1. A steel sheet, comprising:by weight %, carbon (C): 0.10 to 0.25%, 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): 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V): 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities,wherein, in a microstructure, a surface layer portion, which is a region from a surface to a ⅛ point in a thickness direction, includes 80% or more of polygonal ferrite by area %, and a central portion, which is the other region from a ⅛ point to a ½ point in the thickness direction, includes 50% or more of bainite and a remainder of ferrite by area %,wherein a difference between a porosity of a region from a surface to a ⅜ point in the thickness direction and a porosity of a region from a ⅜ point to a ⅝ point in the thickness direction is 0.100 mm3 / g or less, andwherein a hardness value of the surface layer portion is 200 to 220 HB.

2. The steel sheet of claim 1, wherein the surface layer portion includes one or more of pearlite and bainite as a remainder structure.

3. The steel sheet of claim 1, wherein, in the steel sheet, one or more of fine NbC, NbCN, VC, CVN precipitates having a diameter of 5 to 50 nm are present in a density of 10 or more per 1 μm2.

4. The steel sheet of claim 1, wherein the steel sheet has a tensile strength of 510 to 690 MPa, an elongation (ZRA) in the thickness direction of 35% or more, and a Charpy impact absorption energy value at −50° C. of 50 J or more.

5. The steel sheet of claim 1, wherein the steel sheet has a maximum surface crack depth of 1 μm or less in a 180° bending test at room temperature.

6. The steel sheet of claim 1, wherein the steel sheet has a thickness of 133 to 233 mm.

7. A method of manufacturing a steel sheet, the method comprising:primary-reheating a steel slab including, by weight %, carbon (C): 0.10 to 0.25%, 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) 0.0015% or less, niobium (Nb): 0.001 to 0.030%, vanadium (V) 0.001 to 0.030%, titanium (Ti): 0.001 to 0.030%, chromium (Cr): 0.01 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, and a balance of Fe and inevitable impurities;forging the primary-reheated steel slab at a cumulative reduction ratio of 35 to 65% and a strain rate of 1.0 to 4.0 / s;secondary-reheating the forged steel slab;hot-rolling the secondary-reheated steel slab at a finishing-rolling temperature of 900 to 1100° C.;heating the hot-rolled steel sheet in a temperature range of 820 to 900° C., holding for 10 to 40 minutes, and primary-cooling the steel sheet to 700° C. at an average cooling rate of 0.1 to 5.0° C. / s based on a steel sheet surface temperature;secondary-cooling the primary-cooled steel sheet to room temperature at an average cooling rate of 10.0° C. / s or higher based on the steel sheet surface temperature; andperforming a tempering heat treatment of heating the secondary-cooled steel sheet in a temperature range of 550 to 700° C. and holding for 5 to 60 minutes,wherein, in the forging, a cumulative reduction ratio at a recrystallization temperature or lower is 20% or less.

8. The method of claim 7,wherein the primary-reheating is performed in a temperature range of 1100 to 1300° C., andwherein the secondary-reheating is performed in a temperature range of 1000 to 1200° C.

9. The method of claim 7,wherein, in the primary-reheating, a thickness of the steel slab is 650 to 750 mm,wherein, after the forging, a thickness of the steel slab is 350 to 450 mm, andwherein, after the hot-rolling, a thickness of the steel sheet is 133 to 233 mm.

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