Steel material, welding heat affected zone, and manufacturing method thereof

Abstract

The present invention relates to a steel material suitable for use in liquefied gas storage tanks and the like, a welding heat affected zone, and a method for manufacturing the same.

Description

Technical Field

[0001] The present disclosure relates to a steel material suitable for a liquefied gas storage tank, a weld heat-affected zone, and a manufacturing method thereof. Background Art

[0002] Recently, with the strengthening of environmental regulations, the demand for liquefied gas (environmentally friendly fuel) has increased. As a result, the construction of liquefied gas carriers that transport liquefied gas or gas-propelled ships using liquefied gas as fuel has increased. Among liquefied gases, liquefied natural gas and liquefied ethylene / ethane gas have very low liquefaction temperatures, and therefore, high-Ni steel is generally used for their storage. However, LPG, ammonia, CO2, etc. are gases with a liquefaction temperature of -60°C, and carbon steel with improved strength and low-temperature toughness has been used to reduce the cost of building tanks or ships that accommodate these gases.

[0003] Because existing liquefied gas transport tanks primarily utilize large A-type tank designs that don't require high pressure, there's no demand for high-strength or thick materials. However, recently, CO2 tanks and ammonia fuel tanks, which require high pressure during liquefaction, have been designed as C-type tanks, and there's a trend toward requiring high-strength and thick steel materials. Furthermore, since post-weld heat treatment (PWHT), required by the IGC specification, is essential for the construction of C-type tanks, ensuring the steel's properties after PWHT is also necessary.

[0004] PWHT is a process in which materials are held at high temperatures for a certain period of time to release residual stress after welding. During the PWHT process, the grain size and precipitate size of the low-temperature structure increase, which degrades material properties such as strength and toughness, making it difficult to ensure quality. Specifically, in the case of steel used in CO2 tanks, classification societies have recently required the quality of the crack tip opening displacement (CTOD) in the weld heat-affected zone (HAZ) after PWHT as a structural compatibility assessment item to improve the stability of tank designs.

[0005] Meanwhile, in the case of liquefied CO2, the liquefaction temperature varies according to the pressure, and therefore, the tank can be mainly divided into a low-pressure tank and a medium-pressure tank.

[0006] In the case of low-pressure tanks with a typical operating pressure of 5 to 10 bar, since the temperature should be lowered to -60°C for CO2 liquefaction, the steel used for the tanks must also have a tensile strength (TS) of 600 MPa to 650 MPa that can ensure low-temperature impact toughness below -60°C. In order to maintain the low temperature after liquefaction, separate cooling facilities are required, which increases initial installation costs and maintenance costs.

[0007] In the case of medium-pressure tanks, which typically operate at pressures of 18 to 20 bar, the liquefaction temperature rises to -35°C, facilitating the storage and transportation of liquefied CO₂. However, due to the high operating pressure, the steel used for medium-pressure tanks uses ultra-high-strength steel with a strength of 780 MPa or higher, rather than the conventional TS steel of 600 to 650 MPa. Because high-alloy steel is manufactured through a quenching-tempering (QT) heat treatment to ensure high strength, ensuring weld quality is difficult.

[0008] Patent Document 1 discloses that a steel material capable of ensuring low-temperature impact toughness at -40°C can be produced by heating a slab containing an alloy component comprising, by weight%, C: 0.06% to 0.12%, Si: 0.02% to 0.06%, Mn: 0.7% to 1.2%, Ti: 0.006% to 0.012%, Al: 0.002% to 0.01%, Cr: 0.3% to 0.5%, Mo: 0. 3% to 0.4%, V: 0.03% to 0.04%, N: 0.002% to 0.003%, S: 0.002% to 0.01%, P: 0.007% or less, and the balance Fe and impurities; the slab is then rolled at 800°C to 900°C with a cumulative reduction rate of 30% or more; the slab is air cooled; and the slab is further quenched at 910°C to 930°C; and then, the slab is tempered heat treated at 650°C to 680°C.

[0009] This steel provides a method for improving toughness by uniformly distributing Al-Ti-MnS complex oxide inclusions in the weld heat-affected zone and refining the HAZ grains, but the steel cannot guarantee the quality of the material after PWHT, and the tensile strength of the base material is low, at 550 MPa to 650 MPa, which makes it unsuitable as a steel for medium-pressure liquefied CO2 (LCO2).

[0010] Meanwhile, Patent Document 2 discloses a method for manufacturing a P690QL2 steel material for an LCO2 storage tank, the steel material containing, in weight %, C: 0.09% to 0.12%, Mn: 0.9% to 1.3%, Ti: 0.015% or less, S: 0.002% or less, Nb: 0.06 or less, V: 0.08 or less, Ni: 1.5% to 2.5%, Cr: 0.2% to 0.4%, Mo: 0.3% to 0.4%, Cu: 0.3% or less, Al: 0.03% to 0.05%, Ca: 0.0005% to 0.005%, N: 0.005% or less, H: 0.0003% or less, and Fe and impurities in the balance, and having a thickness of 50 mm. Patent Document 2 discloses that when the material is heated to a temperature of 1160° C. to 1210° C., rolled at 780° C. to 880° C., quenched at 890° C. to 920° C., and subjected to tempering heat treatment at a temperature of 630° C. to 680° C., a tensile strength of 780 MPa or more and an impact toughness of 47 J or more at -60° C. after PWHT can be ensured.

[0011] However, since the CTOD guarantee temperature and standard of the HAZ do not meet the level required by the present disclosure, no practical method is presented in which the material can be used as steel for medium-pressure LCO2 storage tanks.

[0012] Patent Document 3 discloses a steel plate comprising, in terms of weight %, C: 0.03% to 0.18%, Si: 0.30% or less, Mn: 0.3% to 1.6%, P: 0.015% or less, S: 0.005% or less, Cu: 2.0% or less, Ni: 1.0% to 7.0%, Al: 0.01% to 0.20%, Ca: 0.005% or less, N: 0.0100% or less, O: 0.0060% or less, and the balance being Fe and impurities, wherein the content ratio of Mn to Ni is Mn / Ni of 0.80 or less, and the carbon equivalent Ceq based on WES is 0.43. 0% to 0.900%, a tensile strength of 780 MPa or more and 930 MPa or less, a yield strength of 630 MPa or more and 750 MPa or less, and a yield strength ratio of 85% or less, a HAZ impact toughness evaluated at 0°C of 100 J or more, a plate thickness of 40 mm or more and 120 mm or less, and a Vickers hardness of 225 points or more measured at a position 1 / 4 of the plate thickness from the surface, and when the average value of the Vickers hardness from the smaller side to 20% is defined as Hvmin and the average value from the larger side to 20% is defined as Hvmax, Hvmin / Hvmax is 0.85 or less. The strength meets the range required for medium-pressure LCO2 steel, but the steel plate may not be suitable because the impact toughness cannot be guaranteed at about -40°C (the temperature at which CO2 liquefies). In addition, when Ceq and hardness are too high, the martensite fraction may become too high during quenching, which may make it difficult to ensure impact toughness and HAZ CTOD characteristics.

[0013] Therefore, in order to ensure an appropriate low-temperature phase fraction, there are requirements for the range of Ceq and the corresponding surface hardness value, but since this technology does not describe this matter, it is not suitable as a technology for manufacturing ultra-high strength steel for medium-pressure LCO2.

[0014] Therefore, the above-mentioned conventional technology does not present a practical method for manufacturing ultra-high strength steel for medium-pressure LCO2, which has a yield strength of 690 MPa or more and a tensile strength of 770 MPa to 940 MPa after post-weld heat treatment (PWHT), and has low-temperature impact toughness at -40°C or less and a CTOD quality of 0.1 mm or more.

[0015] (Patent Document 1) Korean Patent Publication No. 10-2022-7030534 (published on October 5, 2022)

[0016] (Patent Document 2) Chinese Patent Publication No. CN2022-11506158 (published on November 29, 2022)

[0017] (Patent Document 3) PCT Patent Publication No. WO2021-255858-023761 (published on December 23, 2021) Summary of the Invention

[0018] Technical issues

[0019] One aspect of the present disclosure relates to a weld heat affected zone having excellent strength and low temperature impact toughness and excellent low temperature CTODD quality after post weld heat treatment (PWHT) and a method of making the same.

[0020] Furthermore, another aspect of the present disclosure relates to a steel material having excellent strength and low-temperature impact toughness after post-weld heat treatment (PWHT) and a method of manufacturing the same.

[0021] The aspects of the present disclosure are not limited to the above matters. Other aspects of the present disclosure are described in the entire content of the specification, and according to the content described in the specification of the present disclosure, it will not be difficult for those skilled in the art to which the present disclosure belongs to understand the other aspects of the present disclosure.

[0022] Technical Solution

[0023] One embodiment of the present disclosure relates to a steel material comprising, in terms of weight %, C: 0.05% to 0.20%, Si: 0.05% to 0.50%, Mn: 0.5% to 2.0%, Al: 0.005% to 0.10%, P: 0.010% or less, S: 0.015% or less, Nb: 0.001% to 0.070%, V: 0.001% to 0.30 %, Ti: 0.001% to 0.030%, Cr: 0.01% to 1.0%, Mo: 0.01% to 1.0%, Cu: 0.01% to 0.60%, Ni: 1.0% to 4.0%, B: 0.001% to 0.005%, Ca: 0.0005% to 0.0040%, O: 0.001% or less, and the balance of Fe and other inevitable impurities,

[0024] Wherein Ceq defined in the following relational expression 1 satisfies the range of 0.50 to 0.70,

[0025] The microstructure contains 50% by area or more of tempered martensite,

[0026] The microstructure has a packet size of 17 μm or less, and

[0027] The maximum size of the CaO-Al2O3-X (where X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less, and the upper layer is the area from the surface layer of the steel material to t / 4 (where t is the thickness, unit: mm) from the surface in the thickness direction, and every 1 μm 2 The number of said inclusions per unit area is 50 or less,

[0028] [Relational expression 1]

[0029] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni+[Cu]) / 15

[0030] In relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] respectively represent the contents (weight %) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, and are replaced by 0 when these components are not intentionally added.

[0031] The microstructure may contain bainite and other unavoidable structures.

[0032] When the tempered martensite exceeds 80%, the cross-sectional Brinell hardness of the steel can be 220HB to 290HB.

[0033] In the steel material, the number of VC or VCN precipitates with a diameter of 5 nm to 15 nm among the precipitates observed in the cross section of the steel material may be 1 μm. 2 At least 5.

[0034] This microstructure can be observed within a prior austenite grain size with an average grain size of 30 μm or less.

[0035] The steel material may have a yield strength of 690 MPa or more, a tensile strength of 770 MPa to 940 MPa, and a Charpy impact absorption energy of 50 J or more at -40°C.

[0036] Another embodiment of the present disclosure relates to a method for manufacturing a steel material, comprising: manufacturing a steel billet satisfying the above-mentioned composition and Ceq, wherein the maximum size of CaO-Al2O3-X (wherein X is at least one of Ti, Mg, or S) complex oxide inclusions observed in the upper layer is 50 μm or less, the upper layer being the surface layer of the steel billet to a region t / 4 (wherein t is thickness, unit: mm) from the surface in the thickness direction, and the maximum size of the CaO-Al2O3-X (wherein X is at least one of Ti, Mg, or S) complex oxide inclusions observed in the upper layer is 50 μm or less .... 2 The number of said inclusions per unit area is 50 or less;

[0037] reheating the steel billet to a temperature range of 1050°C to 1300°C;

[0038] The steel is produced by hot rolling a steel slab at a finishing hot rolling temperature of 830° C. to 1050° C.;

[0039] heating the steel material to a temperature range of 820° C. to 950° C. and holding the heated steel material for 10 to 40 minutes, and then cooling the steel material to a temperature of 300° C. or less at a cooling rate of 3° C. / second to 80° C. / second based on the steel thickness t / 4 temperature; and

[0040] After cooling, the steel is heated to 550°C to 700°C and kept at this temperature for 5 to 60 minutes for a tempering heat treatment operation.

[0041] Manufacturing steel billets can include:

[0042] Adding a metal Ca wire rod into the molten steel after secondary refining so that the amount of Ca added is 0.015 kg / ton to 0.15 kg / ton; and

[0043] The cleaning bubbling is performed for 5 to 40 minutes so that after the Ca is added, the inert gas blowing rate in the ladle is 10 to 50 liters / minute.

[0044] The input speed of the Ca wire is 100 m / min to 300 m / min, and the number of inert gas blowing points in the ladle is 2.

[0045] The method for manufacturing the steel material may further include performing multi-stage stacked cooling in a temperature range of 200° C. or higher to room temperature after the hot rolling.

[0046] The method for manufacturing the steel material may further include performing post-weld heat treatment on the steel material.

[0047] The post weld heat treatment may be performed at 595° C. to 625° C. for 180 minutes (min) to 540 minutes.

[0048] When the temperatures of the tempering heat treatment and the post-weld heat treatment are the same as each other, LMP according to the following relational expression 2 may be 17.0 to 19.5.

[0049] [Relational expression 2]

[0050] LMP=T k (Log t+20)

[0051] Where T k = Kelvin temperature of tempering and PWHT, and t = tempering and PWHT time (hours)

[0052] When the temperatures of the tempering heat treatment and the post-weld heat treatment are different from each other, the LMP according to the following Relationship 3 may be 17.0 to 19.5.

[0053] [Relational expression 3]

[0054] T(logt eq +20)=T i (logt i +20)

[0055] Total LMP = T i [log(t i +t eq )+20]

[0056] Where T i : Kelvin temperature of PWHT, t i : PWHT time (hours), T: tempering Kelvin temperature, and t eq : Equivalent tempering time when converted to PWHT temperature (hours).

[0057] Another embodiment of the present disclosure relates to a weld heat affected zone satisfying the above composition and Ceq,

[0058] The maximum size of the CaO-Al2O3-X (where X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less, and the upper layer is the area from the surface layer of the steel slab to t / 4 (where t is the thickness, unit: mm) from the surface in the thickness direction, and the maximum size of the CaO-Al2O3-X (where X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less. 2 The number of said inclusions per unit area is 50 or less,

[0059] The grain boundary occupancy of cementite present in the grain boundaries is 10% or less, and the average size of the cementite is 1 μm or less, and

[0060] The cross-section hardness is 200HB to 300HB.

[0061] [Relational expression 1]

[0062] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15

[0063] In relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] respectively represent the contents (weight %) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, and are replaced by 0 when these components are not intentionally added.

[0064] Another embodiment of the present disclosure is directed to a method of manufacturing a weld heat affected zone, comprising:

[0065] Producing steel that meets the above composition and Ceq,

[0066] The maximum size of the CaO-Al2O3-X (where X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less, and the upper layer is the area from the surface layer of the steel slab to t / 4 (where t is the thickness, unit: mm) from the surface in the thickness direction, and the maximum size of the CaO-Al2O3-X (where X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less. 2 The number of said inclusions per unit area is 50 or less;

[0067] Welding of steel; and

[0068] Perform post-weld heat treatment on welded steel.

[0069] Beneficial effects

[0070] According to one aspect of the present disclosure, deterioration of physical properties can be minimized even after post weld heat treatment (PWHT), and thus, a steel material having excellent strength and low-temperature impact toughness and a method for manufacturing the same can be provided.

[0071] According to another aspect of the present disclosure, a weld heat affected zone having excellent low-temperature impact toughness and low-temperature CTOD properties after post-weld heat treatment (PWHT) and a method for manufacturing the same may be provided.

[0072] The advantages and effects of the present application are not limited to the foregoing, and they can be more easily understood in the course of describing specific embodiments of the present disclosure. DETAILED DESCRIPTION

[0073] The terms used in this specification are intended to describe specific embodiments and are not intended to limit the present disclosure. In addition, unless the relevant definition indicates the opposite meaning, the singular form used in this specification includes the plural form.

[0074] The meanings of “include” and “comprising” used in the specification specify a configuration and do not exclude the existence or addition of other configurations.

[0075] Unless otherwise defined, all terms used herein, including technical and scientific terms, have the same meaning as commonly understood by one of ordinary skill in the art to which the exemplary embodiments of the present disclosure belong. It should also be understood that terms, such as those defined in commonly used dictionaries, should be interpreted as having a meaning consistent with the relevant technical literature and the current disclosure.

[0076] Hereinafter, preferred embodiments of the present disclosure will be described. These embodiments of the present disclosure can be modified in various forms, and the scope of the present disclosure should not be interpreted as being limited to the embodiments described below. These embodiments are provided to further describe the present invention to those skilled in the art.

[0077] First, the composition of a steel material which is one example of the present disclosure will be described in detail.

[0078] The steel material contains, in terms of weight%, C: 0.05% to 0.20%, Si: 0.05% to 0.50%, Mn: 0.5% to 2.0%, Al: 0.005% to 0.10%, P: 0.010% or less, S: 0.0050% or less, Nb: 0.001% to 0.070%, V: 0.001% to 0.30%, Ti: 0.001% to 0.030%, Cr: 0.01% to 1.0%, Mo: 0.01% to 1.0%, Cu: 0.01% to 0.60%, Ni: 1.0% to 4.0%, B: 0.001% to 0.005%, Ca: 0.0005% to 0.0040%, O: 0.002% or less, and the balance being Fe and other inevitable impurities.

[0079] Ceq defined by the following relational expression 1 may be 0.50 to 0.70.

[0080] [Relational expression 1]

[0081] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni+[Cu]) / 15

[0082] In relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] respectively represent the contents (weight %) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, and are replaced by 0 when these components are not intentionally added.

[0083] The composition of the steel material may be the same as not only that of the steel billet but also that of the base material of the steel material before welding, the steel material after welding, and the steel material after welding heat treatment, and the composition of the weld heat affected zone (HAZ).

[0084] Hereinafter, the composition will be described in detail. Hereinafter, % and ppm described in relation to the alloy composition are based on weight unless otherwise specified.

[0085] Carbon (C): 0.05% to 0.20%

[0086] Carbon (C) is the most important element for ensuring basic strength, so carbon (C) needs to be included in the steel within an appropriate range, and in order to obtain this addition effect, 0.05% or more of carbon (C) can be added. Preferably, 0.10% or more of carbon (C) can be included. On the other hand, when the content of carbon (C) exceeds a certain level, the fraction of martensite becomes too high, the hardness of the weld heat affected zone (HAZ) increases, and the fraction of grain boundary cementite increases, so that the impact extension resistance during the impact test or CTOD test is reduced, and therefore the toughness value required in the present disclosure may not be ensured. Therefore, the present disclosure can be carried out with a carbon (C) content of 0.20%. The upper limit of a more preferred carbon (C) content can be 0.18%.

[0087] Silicon (Si): 0.05% to 0.50%

[0088] Silicon (Si) is a substitutional element that improves the strength of steel by solid solution strengthening, and has a strong deoxidizing effect, and therefore, silicon (Si) is an essential element for producing clean steel. Therefore, silicon (Si) may be contained in an amount of 0.05% or more, and preferably in an amount of 0.20% or more. On the other hand, when silicon (Si) is contained in a large amount, silicon (Si) may produce a martensite-austenite (Martensite-Austenite, MA) phase and excessively increase the matrix strength of the martensite phase, thereby deteriorating the low-temperature impact toughness, and therefore, the upper limit of its content may be contained as 0.50%. A more preferred upper limit of the silicon (Si) content may be 0.40%.

[0089] Manganese (Mn): 0.5% to 2.0%

[0090] Manganese (Mn) is a useful element that improves strength and hardenability by solid solution strengthening to produce a low-temperature transformation phase. Therefore, in order to ensure a yield strength of 690MPa or greater, it is preferably possible to include 0.5% or more manganese (Mn). A more preferred manganese (Mn) content may be 0.7% or greater. On the other hand, manganese (Mn) may form MnS (elongated non-metallic inclusions) with sulfur (S), thereby reducing toughness and acting as an impact starting point, which may be a factor that rapidly reduces the low-temperature impact toughness of the steel. Therefore, it is preferably possible to control the content of manganese (Mn) to 2.0% or less, and a more preferred manganese (Mn) content may be 1.5% or less.

[0091] Aluminum (Al): 0.005% to 0.10%

[0092] Aluminum (Al) and silicon (Si) are together one of the powerful deoxidizers in the steelmaking process, and it is preferable to contain 0.005% or more to achieve this effect. A more preferred lower limit of the aluminum (Al) content may be 0.01%. On the other hand, when the aluminum (Al) content is excessive, the fraction of Al2O3 in the oxide inclusions generated as a result of deoxidation excessively increases, and their size becomes coarse, and it may be difficult to remove the inclusions during refining, which may be a factor that reduces low-temperature impact toughness. Therefore, it is preferable to control the aluminum (Al) content to 0.10% or less. A more preferred aluminum (Al) content may be 0.07% or less.

[0093] Phosphorus (P): 0.010% or less (including 0%) and Sulfur (S): 0.0050% or less (including 0%)

[0094] Phosphorus (P) and sulfur (S) are elements that cause embrittlement at grain boundaries or form coarse inclusions that may cause embrittlement. Therefore, phosphorus (P) and sulfur (S) are preferably not contained. In order to improve brittle crack growth resistance, phosphorus (P) may be preferably contained in an amount of 0.010% or less, and sulfur (S) may be preferably contained in an amount of 0.0050% or less.

[0095] Niobium (Nb): 0.001% to 0.070%

[0096] Niobium (Nb) is an element that improves the strength of the base material by precipitating as NbC or NbCN. Furthermore, niobium (Nb), which undergoes solid solution during high-temperature reheating, precipitates significantly finely as NbC during rolling, suppressing the recrystallization of austenite and thus having the effect of refining the structure. Therefore, niobium (Nb) may be preferably included in an amount of 0.001% or more, with a more preferred niobium (Nb) content of 0.005% or more. On the other hand, excessive inclusion of niobium (Nb) may result in undissolved niobium (Nb) in the form of TiNb(C,N), which can inhibit low-temperature impact toughness. Therefore, the upper limit of the niobium (Nb) content may be preferably 0.070%. A more preferred upper limit of the niobium (Nb) content may be 0.065% or less.

[0097] Vanadium (V): 0.001% to 0.30%

[0098] Since vanadium (V) is almost completely dissolved again during reheating, the strengthening effect by precipitation or solution during subsequent rolling is minimal, but vanadium (V) has the effect of improving strength by precipitating significantly fine carbonitrides during tempering heat treatment or PWHT. In order to fully obtain this effect, it is necessary to add 0.001% or more of vanadium (V). A more preferred lower limit of the vanadium (V) content may be 0.01%. On the other hand, when the content is excessive, the surface hardness of the slab may be excessively increased due to the high hardenability, which may become a factor such as surface cracks during flange processing, and the manufacturing cost may also increase rapidly, which is not commercially advantageous. Therefore, the vanadium (V) content can be limited to 0.3% or less. A more preferred vanadium (V) content may be 0.25% or less.

[0099] Titanium (Ti): 0.001% to 0.030%

[0100] Titanium (Ti) is a component that significantly improves low-temperature toughness by precipitating as TiN during reheating and inhibiting the growth of original austenite grains at high temperatures. In order to achieve this effect, it may be preferred to include 0.001% or more of titanium (Ti). On the other hand, when titanium (Ti) is included in excess, low-temperature toughness may be reduced due to clogging of the casting nozzle or centerline inclusions. In addition, since titanium (Ti) combines with nitrogen (N) to form coarse TiN precipitates at the center of the thickness, the elongation of the product is reduced, and therefore, the uniform elongation is reduced during the forging process, resulting in surface cracks. Therefore, the content of titanium (Ti) may be 0.030% or less. The preferred upper limit of the titanium (Ti) content may be 0.025% or less, and the more preferred titanium (Ti) content may be 0.018% or less.

[0101] Chromium (Cr): 0.01% to 1.0%

[0102] Chromium (Cr) is a component that increases yield strength and tensile strength by increasing hardenability and forming a low-temperature transformation structure. Chromium (Cr) is also a component that has the effect of preventing strength reduction by slowing down the spheroidization rate of cementite. For this effect, 0.01% or more of chromium (Cr) may be contained. On the other hand, when the content of chromium (Cr) is excessive, coarse carbides rich in Cr, such as M 23 The size and fraction of C6 increase, which reduces the impact toughness of the steel, and the solubility of niobium (Nb) and the fraction of fine precipitates such as NbC in the steel decrease, which may reduce the strength of the product. Therefore, the upper limit of the chromium (Cr) content in the present disclosure may be 1.0%. The preferred upper limit of the chromium (Cr) content may be 0.8%.

[0103] Molybdenum (Mo): 0.01% to 1.0%

[0104] Molybdenum (Mo) is an element that increases grain boundary strength and has a significant solid solution strengthening effect in ferrite, effectively contributing to increased product strength and ductility. Furthermore, Mo prevents toughness degradation caused by grain boundary segregation of impurity elements such as phosphorus (P). To achieve this effect, 0.01% or more of Mo can be added. However, since Mo is an expensive element, excessive addition can significantly increase manufacturing costs, so the upper limit of the Mo content is preferably 1.0%.

[0105] Copper (Cu): 0.01% to 0.60%

[0106] Copper (Cu) can significantly improve the strength of the matrix by solid solution strengthening in ferrite, and also has the effect of suppressing corrosion in a wet hydrogen sulfide atmosphere, and is a favorable element in the present disclosure. For this effect, 0.01% or more of copper (Cu) may be included. A more preferred copper (Cu) content may be 0.03% or more. However, when the copper (Cu) content is excessive, the possibility of causing star cracks on the surface of the steel plate increases, and since copper (Cu) is an expensive element, there may be a problem of a significant increase in manufacturing costs. Therefore, the upper limit of the copper (Cu) content may preferably be 0.60%, and the more preferred upper limit of the copper (Cu) content may be 0.35%.

[0107] Nickel (Ni): 1.0% to 4.0%

[0108] Nickel (Ni) is an element that enhances stacking faults at low temperatures, makes the cross slip of dislocations easier, thereby improving impact toughness, and effectively contributes to improving strength by improving hardenability. For this effect, 1.0% or more of nickel (Ni) may be included. A preferred nickel (Ni) content may be 0.10% or greater. On the other hand, when excessive nickel (Ni) is added, manufacturing costs may also increase due to high costs, and therefore, the upper limit of nickel (Ni) content may preferably be 4.0%, and a more preferred upper limit of nickel (Ni) content may be 3.0%.

[0109] Calcium (Ca): 0.0005% to 0.0040%

[0110] When calcium (Ca) is added after deoxidation with aluminum (Al), calcium (Ca) suppresses the formation of MnS by combining with sulfur (S) that forms MnS inclusions, and at the same time forms spherical CaS to suppress the occurrence of cracks due to hydrogen-induced cracking. In order to sufficiently form CaS from sulfur (S) contained as an impurity, 0.0005% or more of calcium (Ca) may be preferably contained. However, when the amount added is excessive, CaS may be formed and the remaining calcium (Ca) may combine with oxygen (O) to generate coarse oxide inclusions, which may be elongated and broken during rolling, which may reduce lamellar tearing properties. Therefore, the upper limit of the calcium (Ca) content may preferably be 0.0040%.

[0111] Boron (B): 0.001% to 0.005%

[0112] Boron (B) is an effective element for improving the hardenability of hot-rolled steel sheets by delaying the transformation of austenite to ferrite during cooling. To achieve this effect, boron (B) is preferably added in an amount of 0.001% or greater. However, when the content exceeds 0.005%, the strength increases excessively, leading to a decrease in impact toughness. Furthermore, boron oxide may form, which may reduce the surface quality of the steel sheet.

[0113] Oxygen (O): 0.002% or less

[0114] Oxygen (O) combines with Ca or Al in molten steel to form CaO, Al2O3, or CaO-Al2O3 composite oxide inclusions. Oxide inclusions generated at high temperatures are crushed during the rough rolling operation of the slab and elongated in the rolling direction. They remain undissolved not only during the rolling and heat treatment processes of the parent metal but also in the HAZ during welding. Because the ends of the crushed oxide inclusions can serve as starting points for damage during impact and CTOD assessments due to the notch effect, it is recommended to remove as much of them as possible through flotation separation during secondary refining, so that the oxygen concentration in the molten steel can be preferably 20 ppm or less. The present disclosure proposes a secondary refining method that can extremely control oxide inclusions. After applying this method, the total oxygen content can be more preferably 15 ppm or less, and most preferably 10 ppm or less.

[0115] In addition to the above components, Fe and other unavoidable impurities may be included in the balance. However, since unintended impurities may inevitably be introduced from raw materials or the surrounding environment during the normal manufacturing process, it may not be possible to completely eliminate these unavoidable impurities. Since these impurities are known to anyone of ordinary skill in the art, not all of them are specifically mentioned in this specification.

[0116] Furthermore, the addition of other effective components in addition to the above components is not completely excluded.

[0117] The Ceq of the composition as defined by the following Relational Expression 1 may be 0.50 to 0.70.

[0118] [Relational expression 1]

[0119] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15

[0120] In relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] respectively represent the contents (weight %) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, and are replaced by 0 when these components are not intentionally added.

[0121] When Ceq of Relational Expression 1 is less than 0.50, the high yield strength value required in the present disclosure may not be ensured due to the decrease in martensite fraction, and the bainite fraction is excessively increased, so that after post-weld heat treatment (PWHT), the occupancy rate of grain boundary cementite in the weld heat-affected zone (HAZ), especially the coarse grained heat-affected zone (CGHAZ), increases, and thus, the appropriate impact toughness and HAZ CTOD value required in the present disclosure may not be ensured. In addition, when the Ceq value exceeds 0.70, the hardness may be excessively increased, and cleavage fracture may be induced, and thus, crack growth resistance is reduced, which may make it difficult to ensure appropriate impact toughness and CTOD characteristics in the weld heat-affected zone.

[0122] In the steel material of the present disclosure, the maximum size of the CaO-Al2O3-X (wherein X is at least one of Ti, Mg, or S) complex oxide inclusions observed in the upper layer, which is a region from the surface layer to a distance t / 4 (t: thickness) from the surface in the thickness direction of the steel material based on a cross section, and per 1 μm can be 50 μm or less. 2 The number of inclusions per unit area is 50 or less.

[0123] The complex oxide inclusions can be observed not only in the steel billet but also in the base material of the steel before welding, the steel after welding, and the steel after welding heat treatment, and in the heat affected zone (HAZ).

[0124] When the composite oxide inclusions are formed in the billet manufacturing operation, the shape and fraction do not change in the subsequent product manufacturing operation, and during rough rolling, the composite oxide inclusions may be crushed and elongated in the rolling direction. The crushed oxide inclusions may act as locations of stress concentration and may act as starting points when fracture occurs. Therefore, when the amount or size of oxide inclusions generated during the billet manufacturing process is large, the amount of crushed oxide inclusions after rough rolling increases, which becomes a major cause of deterioration in impact toughness and CTOD quality. Therefore, when the size of the composite oxide inclusions in the upper layer of the product (which is the region from the surface layer to t / 4) exceeds 50 μm or every 1 μm 2 When the number of exceeds 50, it may be difficult to ensure the low-temperature impact toughness at -40°C based on the base material and the HAZ CTOD quality at -35°C or lower.

[0125] The steel material of the present disclosure has tempered martensite or a tempered martensite-bainite mixed structure as a basic structure. Specifically, the steel material may contain 50% or more of tempered martensite by area fraction, and may contain bainite and unavoidable structures in the balance.

[0126] When the matrix structure is a tempered martensite-bainite mixed structure, the area fraction of the tempered martensite may be 50% to 80%.

[0127] When the matrix structure comprises a single structure of more than 80% tempered martensite, the cross-sectional hardness of the steel base material after PWHT may be 220 HB to 290 HB in terms of Brinell.

[0128] The inevitably occurring structure may be a martensite-austenite (MA) component or the like.

[0129] The lath group size of the low-temperature structure of tempered martensite or bainite may be 17 μm or less.

[0130] The average grain size of the prior austenite of the steel may be 30 μm or less.

[0131] Fine VC or VCN precipitates with a diameter of 5nm to 15nm observed on the cross section of the steel material may include precipitates of 1μm 2 5 or more precipitates.

[0132] The microstructural characteristics may be the same in the steel material before welding, the base material after welding, and the base material of the steel material subjected to heat treatment after welding.

[0133] The average grain size of the original austenite of the steel can be 30 μm or less. Post-weld heat treatment (PWHT) is generally performed at a high temperature of 595°C or higher to prevent cracks in the weld zone by stress relief after welding. In this case, the dislocation density is reduced and the cross-sectional hardness can be reduced, so that the integrity of the weld zone can be increased. However, during the long period of maintenance at high temperature, the cementite produced at the interface of the bainite lath decomposes, and carbon undergoes long-range diffusion into the austenite grain boundaries. Therefore, when the grain size increases, the grain boundary area that the cementite may occupy decreases, so that segregation increases, which may deteriorate the impact toughness. Therefore, it is necessary to control the original austenite grain size under appropriate low-temperature pressure. Therefore, the original austenite grain size is preferably 30 μm or less, and more preferably 25 μm or less.

[0134] As the temperature and time of PWHT increase, coarse cementite spheroidizes at the original austenite grain boundaries, and some of it forms film-like cementite. When such cementite forms at the grain boundaries, the grain boundary strength becomes lower than the shear strength, which ultimately becomes the main cause of intergranular fracture. Therefore, in order to prevent intergranular fracture after PWHT, martensite that only causes short-range diffusion should be included in an appropriate fraction or more. However, when the proportion of martensite is significantly high, the cross-sectional hardness becomes too high, and even if intergranular fracture occurs, the impact absorption energy value is reduced, which ultimately leads to toughness degradation. Therefore, the steel microstructure can preferably be formed by a mixed structure of tempered martensite and bainite, and in this case, the fraction of tempered martensite can preferably be 50% to 80% by area fraction.

[0135] However, when the basic structure of the steel microstructure is tempered martensite, that is, when martensite exceeds 80% (including 100%), the tempering temperature should be further increased to maintain the cross-sectional hardness of the base material at 200HB to 290HB. When the cross-sectional hardness range is deviated, it is difficult to ensure the low-temperature impact toughness and CTOD value required in the present disclosure.

[0136] The lath group size of the low-temperature structure of tempered martensite or bainite can be 17 μm or less. The lath group size is a group unit of laths having the same orientation relationship in a low-temperature transformation structure such as bainite or martensite, and can be understood as a domain wall that may have resistance when a physical crack extends (a property that the crack bends at the boundary when the crack extends). As an example of lath group size measurement, the lath group size can be measured as a grain boundary that can be determined when domains within an orientation relationship of 15 degrees are set to the same orientation using an electron backscattered diffraction (EBSD) device. Since bainite or martensite grows only within the grains of the original austenite crystal, it is smaller than the size of the original austenite grains. When the lath group size exceeds 17 μm, it is difficult to ensure sufficient impact toughness due to difficulty in ensuring sufficient crack extension resistance, and therefore, the lath group size may be preferably 17 μm or less.

[0137] On the other hand, during tempering and PWHT, the steel's strength decreases due to a decrease in dislocation density. When fine precipitates are used to appropriately compensate for this decrease, strength degradation during the heat treatment process can be reduced. Normalized or thermomechanically controlled process (TMCP) materials typically utilize precipitates such as NbC and NbCN, but these precipitates are generated at temperatures between 800°C and 900°C and are induced by deformation, making them unsuitable for quench-and-temper (QT) heat-treated steel. Therefore, vanadium carbides / carbonitrides, such as VC and VCN, that are generated in the 550°C to 700°C temperature range are readily utilized. When the diameter of a VC or VCN fine precipitate is less than 5 nm, its size can be significantly small, resulting in minimal strength-improving effect. However, when its size exceeds 15 nm, the fine precipitate is significantly coarse, reducing the strength-improving effect. Therefore, a fine precipitate size of 5 to 15 nm is suitable. On the other hand, when the density of VC and VCN is 1 μm 2 If the number of particles per unit area is less than 5, the precipitation strengthening effect will also decrease. Therefore, the number of particles per 1 μm 2 Five or more per unit area is preferred. The formation of precipitates is determined by the chemical potential and temperature determined by the steel composition. Precipitates of 5 to 15 nm in size can be generated in large quantities, but as the number increases, the precipitates aggregate and become larger in size. Therefore, precipitates of 5 to 15 nm in size do not form indefinitely. Therefore, even if there is no upper limit on the number of fine precipitates, it is not necessarily unclear to those skilled in the art.

[0138] Meanwhile, the microstructure of a weld heat affected zone (HAZ) as an example of the present disclosure will be described in detail The HAZ may be a HAZ after post weld heat treatment (PWHT), and specifically, may be a coarse grain heat affected zone (CGHAZ).

[0139] The grain boundary occupancy rate of cementite present at the grain boundaries of the weld heat affected zone (HAZ) may be 10% or less.

[0140] The grain boundary occupancy can be understood as the length ratio occupied by film cementite at the grain boundary, and is converted into an area ratio in three dimensions. As an example of measuring the occupancy, a specimen is etched with nitric acid (nital) and observed using an SEM or the like, and it is possible to distinguish and measure general grain boundaries and grain boundaries having film cementite therein.

[0141] Meanwhile, the average size of cementite may be 1 μm or less.

[0142] After post weld heat treatment (PWHT), in the HAZ, especially the CGHAZ microstructure, the area fraction of tempered martensite may preferably be 50% or greater. In addition to tempered martensite, the balance may be bainite and other structures. When the area fraction of tempered martensite is less than 50%, the occupancy of film-like cementite at the grain boundaries after PWHT increases, and since the grains become coarser, it is difficult to ensure a suitable CTOD value of 0.1 mm or greater at -35°C or lower. For the same reason, the grain boundary occupancy of cementite present at the grain boundaries after PWHT may preferably be 10% or less, and the average size of cementite present at the grain boundaries may preferably be 1 μm or less.

[0143] The average grain size of the prior austenite in the heat affected zone (HAZ) may be 200 μm or less. The HAZ microstructure is determined by the microstructure of the base material and the influence of welding conditions such as welding heat input / preheating temperature / interpass temperature. Specifically, since the CGHAZ is a portion where the grains become very coarse compared to the base material, the material may be inferior to the base material. The average grain size of the prior austenite according to the CGHAZ may preferably be 200 μm or less, and when the average grain size exceeds 200 μm, the hardenability increases excessively and the cross-sectional hardness increases, so the average grain size of the CGHAZ may preferably be 200 μm or less. The average grain size of the prior austenite may not be different before and after PWHT. In order to change the grain size, the prior austenite should be reversed to the austenite single phase region, and therefore, the temperature should be heated to 800°C or higher, but since the PWHT temperature is about 500°C to 600°C, the AGS itself may not change.

[0144] The grain boundary occupancy rate of cementite present at the grain boundaries of the heat-affected zone (HAZ) may be 10% or less.

[0145] The grain boundary occupancy rate can be understood as the length ratio occupied by film-like cementite at the grain boundaries, and in three dimensions, it becomes the area ratio. As an example of measuring the occupancy rate, after etching the sample with nital, the sample can be observed using an SEM or the like, and it is possible to distinguish and measure normal grain boundaries and grain boundaries in which film-like cementite is formed.

[0146] The steel material of the present disclosure can maintain the strength and low-temperature impact toughness of the parent material after PWHT. Specifically, the steel material can have a yield strength of 690 MPa or higher, a tensile strength of 770 MPa to 940 MPa, and a low-temperature impact toughness absorbed energy of 50 J or greater in a Charpy V-notch test of the parent material at -40°C or lower.

[0147] At the same time, after PWHT, the weld heat-affected zone, especially the CGHAZ, can ensure excellent low-temperature impact toughness and CTOD properties. Specifically, the low-temperature impact toughness absorbed energy in a Charpy V-notch test conducted at -40°C or lower can be 40J or greater, and the CTOD property at -35°C or lower can be 0.1mm or greater.

[0148] Hereinafter, one example of the method of manufacturing a steel material of the present disclosure will be described in detail.

[0149] This manufacturing method can produce a steel billet that satisfies the alloy composition and the Ceq range of Relational Expression 1 described above, and can produce a steel material by heating, hot rolling, reheating, quenching, and tempering the billet. Since the reheating-quenching-tempering process is performed, this can be referred to as a so-called RQT method.

[0150] Steels manufactured in this way can be welded and post-weld heat treated (PWHT).

[0151] Each process can be described in detail below.

[0152] Billet manufacturing

[0153] A steel slab is produced in which Ceq defined in the above alloy composition and relational expression 1 satisfies the range of 0.50 to 0.70, the maximum size of CaO-Al2O3-X (wherein X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer, which is a region from the surface layer to t / 4 (t: thickness, unit: mm) from the surface of the steel slab in the thickness direction, is 50 μm or less, and the maximum size of the CaO-Al2O3-X (wherein X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less. 2The number of inclusions per unit area is 50 or less.

[0154] As an example of a method for manufacturing a steel billet, a calcium (Ca) raw material may be added to secondary refined molten steel, and a clean bubbling process may be performed. As a specific example, a metallic Ca wire may be added to the molten steel after secondary refining, and the amount of Ca may be added so as to be 0.015 kg / ton to 0.15 kg / ton. The metallic Ca wire may be formed from a steel material coated with a Ca alloy, the steel material may have a thickness of 1.2 mm to 1.4 mm, and the metallic Ca wire may be added at a speed of 100 m / min to 300 m / min.

[0155] After Ca is added, cleaning bubbling may be performed so that the amount of inert gas blown into the ladle is 10 to 50 liters per minute, the number of inert gas blowing points in the ladle may be 2, and the cleaning bubbling time may be 5 to 40 minutes.

[0156] The present disclosure does not specifically limit the processes before secondary refining, and any conventional method can be applied. The present disclosure does not specifically limit the processes after the Ca addition and the cleaning bubbling process, and the molten steel can be cooled under conventional conditions to produce slabs. In addition, as the amount of Al2O3 in the molten steel increases, the generation and coarsening of inclusions containing both Ca and Al progress, which increases the fragmentation of inclusions during rolling. As a result, the low-temperature impact toughness and CTOD quality of the HAZ may not be ensured. Therefore, the total amount of inclusions in the molten steel before the addition of Ca can be limited to 2 ppm to 5 ppm.

[0157] Hereinafter, the Ca input process and the cleaning bubbling process will be described in detail.

[0158] (1) Calcium (Ca) input

[0159] When manufacturing steel billets, a Ca wire can be added to the molten steel after secondary refining, with the Ca addition amount being 0.015 kg / ton to 0.15 kg / ton. The Ca wire is formed from a steel material coated with a Ca alloy, and the steel material can have a thickness of 1.2 mm to 1.4 mm. The Ca wire can be added at a speed of 100 m / min to 300 m / min.

[0160] When Ca is added, if the thickness of the Ca alloy steel covering the metallic Ca wire is less than 1.2 mm, the Ca melts in the upper portion of the ladle, reducing the static pressure and resulting in poor Ca yield, which in turn increases the amount of Ca added. On the other hand, if the thickness exceeds 1.4 mm, the metallic Ca wire contacts the base of the ladle, causing problems with the ladle's refractory material melting therein, and thus, operational stability may not be ensured.

[0161] When adding Ca to molten steel, the rate at which the wire is added must be controlled to ensure Ca yield and the thickness of the metallic Ca wire. If the wire is added at a rate of less than 100 m / min, Ca melts in the upper portion of the ladle, reducing static pressure and resulting in poor Ca yield, which in turn increases the amount added. On the other hand, if the wire is added at a rate exceeding 300 m / min, the Ca wire contacts the base of the ladle, causing problems with melting the ladle's refractory material and potentially preventing stable operation. A preferred rate of addition is 120 to 300 m / min, and more preferably 140 to 180 m / min.

[0162] When the amount of Ca added is significantly low, MnS is generated in the central portion during solidification, and low-temperature impact toughness may not be ensured due to segregation defects, while when the amount is excessive, there is a problem that the refractory reacts with Al2O3 and accelerates the dissolution of the refractory, making it difficult to ensure operational stability. Therefore, considering the above problems, the amount of Ca added may be 0.015 kg / ton to 0.15 kg / ton, preferably 0.015 kg / ton to 0.10 kg / ton, and more preferably 0.050 kg / ton to 0.10 kg / ton.

[0163] (2) Cleaning the bubbler

[0164] After Ca is added, cleaning bubbling may be performed for 5 to 40 minutes so that the inert gas blowing rate in the ladle becomes 10 to 50 liters / minute. The number of inert gas blowing points in the ladle may be two.

[0165] If the amount of inert gas blown into the ladle is significantly low, the amount of Al2O3 clusters and complex inclusions containing both Ca and Al that are removed by adhering to the inert gas is low, resulting in poor cleanliness and the potential failure to ensure HAZ impact toughness. On the other hand, if the amount is too high, the stirring force may be strong, causing slag mixing and exposure of the molten steel surface, which may lead to poor cleanliness. Therefore, the inert gas blowing rate can be 10 to 50 liters / minute, preferably 15 to 40 liters / minute, and more preferably 20 to 30 liters / minute.

[0166] On the other hand, when there is only one inert gas blowing point in the ladle, uneven areas may exist in the molten steel, which may lead to poor removal of Al2O3 clusters and complex inclusions containing both Ca and Al, while when there are three or more inert gas blowing points, during gas blowing, the stirring force may be strong due to the occurrence of overlapping parts, causing slag mixing and at the same time causing the molten steel surface to be exposed, which may lead to poor cleanliness.

[0167] Although the amount of inert gas blown into the ladle is limited, if the cleaning bubbling time is significantly short, the amount of Al2O3 clusters and complex inclusions containing both Ca and Al removed by adhering to the inert gas may be low, which may lead to poor cleanliness and may make it impossible to ensure toughness. On the other hand, if the time is too long, the temperature drop of the molten steel may be large, and a temperature gradient may occur in the ladle, which may lead to poor cleanliness. The cleaning bubbling time can be 5 to 40 minutes, preferably 7 to 17 minutes, and more preferably 10 to 15 minutes.

[0168] Heating of steel billets

[0169] The steel slab may be heated to a temperature in the range of 1050°C to 1300°C.

[0170] Heating is performed by redissolving Ti or Nb carbonitrides or coarse TiNb(C,N) crystallites formed during casting, and heating and maintaining the austenite at a temperature above the recrystallization temperature after sizing. Therefore, the slab heating temperature is preferably 1150°C or higher to maximize the austenite grain size. On the other hand, if the temperature is too high, problems may arise due to high-temperature oxide scale, and excessive grain growth may occur, deteriorating low-temperature impact toughness. Therefore, the upper limit is preferably 1300°C.

[0171] Hot Rolling

[0172] The heated slab may be hot rolled at a finishing hot rolling temperature of 830°C to 1050°C.

[0173] 830°C corresponds to a range equal to or higher than the recrystallization temperature. Since the strength of the steel plate increases as the rolling temperature decreases, the complex inclusions generated during the refining process should conform to the deformation caused by rolling, and therefore, the complex inclusions may be crushed or split into small-sized inclusions, or inclusions such as MnS may be elongated. Since such crushed or split small-sized inclusions or elongated inclusions act as a direct cause of the initiation and expansion of low-temperature impact toughness, the finish hot rolling may be preferably performed at 830°C or higher, which is the recrystallization temperature at which work hardening does not occur. However, when the temperature exceeds 1050°C, austenite grain growth continues to occur even after the end of rolling, which may reduce the impact toughness due to the increase in the ductile-brittle transition temperature (DBTT).

[0174] The thickness of the steel plate after finish hot rolling may be 5 mm to 100 mm, more preferably 5 mm to 70 mm, and further preferably 5 mm to 50 mm.

[0175] The steel sheet produced by hot rolling can be air-cooled to room temperature. In principle, air cooling is performed after hot rolling is completed, but accelerated cooling to an appropriate temperature can be performed after rolling to suppress precipitate growth and shorten processing time.

[0176] In the present disclosure, when manufacturing a slab, when the content of dissolved hydrogen in molten steel is 1.3 ppm or greater, an operation of multi-step cooling from a temperature of 200° C. or higher to room temperature based on the surface temperature after hot rolling may be further included before reheating heat treatment. Multi-step cooling releases hydrogen dissolved in the steel, which can more effectively suppress internal microcracks caused by hydrogen and ultimately improve the low-temperature impact toughness of the base material and the weld heat-affected zone.

[0177] Reheating

[0178] After hot rolling, the steel plate may be reheated by heating the steel plate to a temperature of 820° C. to 950° C. and maintaining the steel plate for 10 minutes to 40 minutes.

[0179] When the heating temperature during reheating is lower than 820°C or the holding time is less than 10 minutes, carbides or grain boundary segregated impurity elements generated during cooling after rolling may not be smoothly redissolved and may remain even after cooling, which may be a major cause of deterioration in low-temperature impact toughness. On the other hand, when the heating temperature exceeds 950°C or the holding time exceeds 40 minutes, the austenite grain size may excessively increase due to grain growth, which may be a major cause of deterioration in toughness due to increased DBTT. Therefore, the reheating temperature may be preferably 820°C to 950°C, and more preferably 850°C to 890°C.

[0180] Quenching and tempering

[0181] The reheated steel plate may be quenched and tempered.

[0182] The cooling rate during the quenching process can be considered one of the most important process factors that determine the microstructure of the base material. In the present disclosure, after reheating the steel plate, the steel plate can be cooled to 300°C or less at a cooling rate of 3°C / second to 80°C / second based on the temperature at a point of t / 4 (t: thickness, unit: mm) of the steel material thickness. When the cooling rate ends at 300°C or higher, due to incomplete cooling transformation, the base material microstructure may contain martensite-austenite (MA) components that are not dissolved during subsequent tempering and PWHT processes, resulting in deterioration of low-temperature impact toughness. Therefore, the cooling end temperature may be preferably 300°C or less.

[0183] On the other hand, if the cooling rate is less than 3°C / sec, the base material's microstructure is formed from a mixed structure of ferrite and bainite or a single-phase structure of bainite, rather than a mixed structure of martensite and bainite or a single-phase structure of martensite as required by the present disclosure. While bainite becomes the primary microstructure, strength and low-temperature impact toughness before PWHT can be ensured, low-temperature impact toughness and CTOD quality are not easily ensured due to the formation of grain boundary cementite after PWHT. On the other hand, when the cooling rate exceeds 80°C / sec, the microstructure is formed from 100% martensite, and the cross-sectional hardness increases to a level that may be uncontrollable by tempering temperature alone. There is also a risk of plate fracture due to the difference in cooling rate between the surface layer and the center.

[0184] Therefore, the cooling rate during quenching is preferably 3° C. / second or more and 80° C. / second or less, and more preferably 5° C. / second or more and 30° C. / second or less.

[0185] The quenched steel sheet can be heated to a temperature of 550°C to 700°C and held for 5 to 60 minutes to perform a tempering heat treatment. If the tempering temperature is lower than 550°C or the holding time is less than 5 minutes, dislocation recovery does not occur properly during quenching. As a result, sufficient low-temperature impact toughness may not be ensured due to excessive strength of the base material. If the tempering temperature exceeds 700°C or the holding time exceeds 60 minutes, the grain boundaries (cementite) in the weld heat-affected zone may become coarsened, making it difficult to ensure toughness after PWHT.

[0186] Meanwhile, in addition to RQT, steel can also be manufactured by a method of directly quenching and tempering after hot rolling (so-called DQT).

[0187] The fabricated steel can be welded and can be post weld heat treated.

[0188] welding

[0189] The steel of the present disclosure can be used by performing post weld heat treatment (PWHT) after welding, and the steel of the present disclosure can ensure the strength and low temperature impact toughness of the base material after post weld heat treatment, and can ensure the low temperature impact toughness and low temperature CTOD characteristics of the weld heat affected zone (HAZ).

[0190] When welding steel, there is an unaffected base material, a molten part (weld metal) that is melted and joined, and a weld heat affected zone (HAZ) which is a part of the base material affected by the welding heat, and the molten part and the weld heat affected zone can be collectively referred to as the weld part. At the same time, according to the maximum heating temperature and histological characteristics, the weld heat affected zone can be divided into coarse-grained HAZ (CGHAZ), subcritical HAZ (SubcriticalHAZ, SCHAZ), critical HAZ (IntercriticalHAZ, ICHAZ), and fine-grained HAZ (FGHAZ). Among them, CGHAZ is the part in which the original austenite grain size is the largest in the heat affected zone (HAZ) during welding, and is maintained at a high temperature of 1000°C or higher and then cooled. When the grain size is large, the ductile-brittle transition temperature increases, and due to the increase in hardenability, when a hard structure is produced, the brittle fracture resistance decreases. Therefore, the CGHAZ is a portion where the low-temperature impact toughness and the toughness value of CTOD in the HAZ may be most deteriorated, and the characteristics of the CGHAZ may be considered critical.

[0191] Welding can be performed by a method that can be performed in the technical field to which the present disclosure pertains, and is not necessarily limited. In the present disclosure, the thermal cycle of welding can be performed by heating to a maximum temperature of 1250°C to 1400°C at a maximum heating rate of 50°C / second, holding the target temperature for 5 to 10 seconds, performing a primary cooling to 500°C at a cooling rate of 10°C / second to 20°C / second, and then performing a secondary cooling to 200°C at a cooling rate of 5°C / second to 15°C / second. The thermal cycle associated with the maximum temperature and cooling rate can be measured using a thermocouple at the CGHAZ position during actual welding and can be set based on this as described above.

[0192] During actual welding, such as FCAW and SAW, the thermal cycle of the weld is measured as consistently as possible by measuring the thermal cycle temperature of the weld heat-affected zone (HAZ). As a practical example of thermal cycling, the heat input for flux-cored arc welding (FCAW) can be 1.0 to 2.0 kJ / mm, and for submerged arc welding (SAW) can be 2.5 to 4.0 kJ / mm. In this case, both the preheat temperature and the interpass temperature can be 150 to 250°C.

[0193] Post-weld heat treatment (PWHT)

[0194] Post-weld heat treatment (PWHT) can be performed using methods known in the art and is not particularly limited. For example, PWHT can be performed by maintaining a temperature of 595° C. to 625° C. for 180 to 540 minutes (min) in accordance with ASME Sec. VIII. Meanwhile, the total LMP during the tempering heat treatment and the PWHT heat treatment according to the following Relational Expressions 2 and 3 may preferably be 17.0 to 19.5.

[0195] Specifically, when the tempering temperature and the PWHT temperature are the same as each other, the LMP is determined by Relational Expression 2, and when the tempering temperature and the PWHT temperature are different from each other, the LMP is determined by Relational Expression 3.

[0196] [Relational expression 2]

[0197] LMP=T k (Log t+20)

[0198] Where T k = Kelvin temperature of tempering and PWHT, t = tempering and PWHT time (hours)

[0199] [Relational expression 3]

[0200] T(logt eq +20)=T i (logt i +20)

[0201] Total LMP = T i [log(t i +t eq )+20]

[0202] Where T i : Kelvin temperature of PWHT, t i : PWHT time (hours), T: tempering Kelvin temperature, t eq : Equivalent tempering time when converted to PWHT temperature (hours)

[0203] When the LMP is less than 17.0, as described above, the dislocation density does not decrease, so that excessive strength failure may occur, or the cross-sectional hardness may be significantly high, which may deteriorate the impact toughness. On the other hand, when the LMP exceeds 19.5, the grain boundary occupancy of cementite after PWHT is significantly high, making it difficult to properly ensure impact toughness and CTOD quality. Therefore, the LMP may be preferably 17.0 to 19.5, and more preferably 18.0 to 19.3.

[0204] Embodiments of the invention

[0205] Hereinafter, embodiments of the present disclosure will be described. It should be understood that those skilled in the art can modify the following embodiments in various ways without departing from the scope of the present disclosure. The following embodiments are provided for understanding the present disclosure, and the scope of the present disclosure should not be limited to the following embodiments, but should be determined by the described claims and their equivalents.

[0206] (Example 1)

[0207] Under the conditions of Table 2, steel billets were manufactured using molten steel having the alloy composition of Table 1 (the contents of the component elements in Table 1 are in weight %, but the units of P, S and Ca are in ppm, and the remaining components are Fe and inevitable impurities) and Ceq of the above-mentioned relational expression 1, and then heated and hot-rolled to manufacture hot-rolled steel plates with a thickness of 50 mm.

[0208] The hot-rolled steel sheets manufactured above were reheated and quenched under the conditions disclosed in Table 3, and then tempered at 610° C. for 20 minutes.

[0209] At the same time, in order to confirm the excellent strength and low-temperature impact toughness of the steel after PWHT, a heat treatment process simulating PWHT was performed. In this case, the hot-rolled steel plate was subjected to a heat treatment process of tempering at 595°C for 180 minutes for three cycles.

[0210] The microstructure and physical properties of the base material were measured for each of the samples produced as described above, and the results are shown in Tables 4 and 5 below.

[0211] Here, electron backscatter diffraction (EBSD) is used to measure the grain size of the prior austenite and the fraction of each phase by collecting samples. In this case, transmission electron microscopy (TEM) is used to measure precipitates and inclusions. Precipitations and inclusions are determined by diffraction patterns and EDX mapping.

[0212] Yield strength / ultimate tensile strength was assessed by room temperature tensile testing, with a 0.2% offset applied to the yield strength. Furthermore, the impact toughness of each specimen was determined by averaging the absorbed energy values ​​measured three times by Charpy V-notch testing at the corresponding temperature. Brinell hardness was measured according to ISO 6506 Part 1.

[0213] [Table 1]

[0214]

[0215] [Table 2]

[0216]

[0217] [Table 3]

[0218]

[0219] [Table 4]

[0220]

[0221] [Table 5]

[0222]

[0223] As can be seen from Tables 1 to 5, in the case of Samples 1 to 5 and 17 to 20, which met the alloy compositions and manufacturing conditions set forth in the present disclosure, all samples were able to ensure excellent qualities, such as the excellent strength of the base material and excellent low-temperature impact toughness at -40°C, even after excellent post-weld heat treatment. In contrast, it can be seen that Samples 6 to 16, which met the alloy compositions set forth in the present disclosure but did not meet the manufacturing conditions, found it difficult to ensure the characteristics of the base material set forth in the present disclosure. Furthermore, Samples 21 to 25, which met the steel manufacturing conditions set forth in the present disclosure but did not meet the alloy compositions, found it difficult to ensure the characteristics of the base material set forth in the present disclosure.

[0224] (Example 2)

[0225] For the steel materials manufactured under the conditions of Table 3 above, welding and post-weld heat treatment were performed to determine the low-temperature toughness and CTOD properties of the CGHAZ in the heat-affected zone (HAZ).

[0226] To this end, for the steel materials manufactured according to Tables 1 to 3, welding simulations were performed under the welding thermal cycle conditions in Table 6, and a post-weld heat treatment process was performed at 595° C. for 180 minutes for 3 cycles. In Table 6, LMP is obtained from the aforementioned relational expression 3.

[0227] The microstructure and physical properties of the CGHAZ in each of the samples manufactured as described above were measured, and the results are shown in Tables 7 and 8 below.

[0228] Here, the phase fractions were measured by collecting samples and using electron backscatter diffraction (EBSD).

[0229] The impact toughness of each sample was determined by the average of three absorbed energy values ​​measured by the Charpy V-notch test at the corresponding temperature. The Brinell hardness was measured according to ISO 6506 Part 1, and the CTOD test was performed according to BS5762 (British Standard).

[0230] In addition, the cementite grain boundary occupancy was measured using transmission electron microscopy (TEM).The cementite grain boundary occupancy measures the ratio of film-like cementite to the total grain boundary length.

[0231] [Table 6]

[0232]

[0233] [Table 7]

[0234]

[0235] [Table 8]

[0236]

[0237] As can be seen from Tables 6 to 8 above, in the case of Samples 1 to 5 and 11 to 16, which satisfy the alloy composition and manufacturing conditions proposed in the present disclosure, it can be seen that excellent low-temperature impact toughness and good CTOD quality can be ensured in the CGHAZ. In contrast, Samples 6 to 10 satisfy the alloy composition proposed in the present disclosure, but do not satisfy the conditions for manufacturing steel, and specifically, since Samples 6 to 10 are cases in which the size of the complex oxide inclusions is coarse, it can be seen that the characteristics of the CGHAZ are poor. On the other hand, Samples 17 to 20 satisfy the alloy composition and manufacturing conditions proposed in the present disclosure, but do not satisfy the welding and post-weld heat treatment conditions, and therefore cannot ensure the characteristics of the CGHAZ.

[0238] Sample Nos. 21 to 25 satisfy the manufacturing conditions proposed by the present disclosure, but do not satisfy the alloy composition, and therefore, it can be seen that the appropriate quality required by the present disclosure is not satisfied.

[0239] As described above, the detailed description of the present disclosure describes the preferred embodiments of the present disclosure, but it is obvious that those skilled in the art can make various modifications within the scope of the present disclosure. Therefore, the scope of the rights of the present disclosure should not be limited to the described embodiments, but should be determined by the described claims and their equivalents.

Claims

1. A steel material comprising, in terms of weight %, C: 0.05% to 0.20%, Si: 0.05% to 0.50%, Mn: 0.5% to 2.0%, Al: 0.005% to 0.10%, P: 0.010% or less, S: 0.015% or less, Nb: 0.001% to 0.070%, V: 0.001% to 0.30%, Ti:

0. 001% to 0.030%, Cr: 0.01% to 1.0%, Mo: 0.01% to 1.0%, Cu: 0.01% to 0.60%, Ni: 1.0% to 4.0%, B: 0.001% to 0.005%, Ca: 0.0005% to 0.0040%, O: 0.001% or less, and the balance of Fe and other inevitable impurities, Wherein Ceq defined in the following relational expression 1 satisfies the range of 0.50 to 0.70, The microstructure contains 50% by area or more of tempered martensite, The lath size of the microstructure is 17 μm or less, and The maximum size of the CaO-Al2O3-X (where X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less, and the upper layer is the area from the surface layer of the steel material to t / 4 (where t is the thickness, unit: mm) from the surface in the thickness direction, and the maximum size of the CaO-Al2O3-X (where X is at least one of Ti, Mg or S) complex oxide inclusions observed in the upper layer is 50 μm or less. 2 The number of said inclusions per unit area is 50 or less, [Relational expression 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15 In relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] respectively represent the contents (weight %) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel material, and when these components are not intentionally added, they are replaced by 0. 2 . The steel material according to claim 1 , wherein the microstructure comprises bainite and other inevitable structures. 3 . The steel material according to claim 1 , wherein when the tempered martensite exceeds 80%, the cross-sectional Brinell hardness of the steel material is 220 HB to 290 HB.

4. The steel material according to claim 1, wherein in the steel material, the number of VC or VCN precipitates having a diameter of 5 nm to 15 nm among precipitates observed in a cross section of the steel material is 1 μm. 2 At least 5. The steel material according to claim 1 , wherein the microstructure is observed within a prior austenite grain size having an average grain size of 30 μm or less. The steel material according to claim 1 , wherein the steel material has a yield strength of 690 MPa or more, a tensile strength of 770 MPa to 940 MPa, and a Charpy impact absorption energy of 50 J or more at −40° C.

7. A method for manufacturing a steel material, comprising: A steel billet is manufactured, the steel billet comprising, in terms of weight %, C: 0.05% to 0.20%, Si: 0.05% to 0.50%, Mn: 0.5% to 2.0%, Al: 0.005% to 0.10%, P: 0.010% or less, S: 0.005% or less, Nb: 0.001% to 0.070%, V: 0.001% to 0.30%, Ti: 0.001% to 0.030%, Cr: 0.01% to 1.0%, Mo: 0.01% to 1.0%, Cu: 0.01% to 0.60%, Ni: 1.0% to 4.0%, B: 0.001% to 0.005%, Ca: 0.0005% to 0.0040%, O: 0.002% or less, and Fe and other inevitable impurities in the balance, wherein Ceq defined in the following relational expression 1 satisfies the range of 0.50 to 0.70, and the maximum size of CaO-Al2O3-X (wherein X is at least one of Ti, Mg, or S) complex oxide inclusions observed in the upper layer, which is the surface layer of the steel slab to a region t / 4 (wherein t is thickness, unit: mm) from the surface in the thickness direction, and per 1 μm is 50 μm or less. 2 The number of inclusions per unit area is 50 or less; reheating the steel billet to a temperature range of 1050° C. to 1300° C.; manufacturing a steel material by hot rolling the steel slab at a finishing hot rolling temperature of 830° C. to 1050° C.; heating the steel material to a temperature range of 820° C. to 950° C. and maintaining the heated steel material for 10 to 40 minutes, and then cooling the steel material to a temperature of 300° C. or less at a cooling rate of 3° C. / second to 80° C. / second based on the steel thickness t / 4 temperature; as well as After the cooling, the steel material is heated to 550° C. to 700° C. and kept at this temperature for 5 to 60 minutes for a tempering heat treatment operation. [Relational expression 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15 In relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] respectively represent the contents (weight %) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel material, and when these components are not intentionally added, they are replaced by 0.

8. The method for manufacturing a steel material according to claim 7, wherein manufacturing the steel billet comprises: Adding a metal Ca wire rod into the molten steel after secondary refining so that the amount of Ca added is 0.015 kg / ton to 0.15 kg / ton; as well as The cleaning bubbling is performed for 5 to 40 minutes so that after the Ca is added, the inert gas blowing rate in the ladle is 10 to 50 liters / minute. The input speed of the Ca wire is 100 m / min to 300 m / min, and the number of inert gas blowing points in the ladle is 2.

9. The method for manufacturing the steel material according to claim 7, further comprising: After the hot rolling, multi-stage stacked cooling is performed in a temperature range of 200° C. or higher to room temperature.

10. The method for manufacturing the steel material according to claim 7, further comprising: The steel is subjected to post-weld heat treatment. 11 . The method for manufacturing a steel material according to claim 10 , wherein the post-weld heat treatment is performed at 595° C. to 625° C. for 180 to 540 minutes (min).

12. The method for manufacturing a steel material according to claim 11, wherein when the temperatures of the tempering heat treatment and the post-weld heat treatment are the same as each other, LMP according to the following relational expression 2 is 17.0 to 19.5, [Relational expression 2] LMP=T k (Log t+20) Where T k = Kelvin temperature of tempering and PWHT, and t = tempering and PWHT time in hours.

13. The method for manufacturing a steel material according to claim 11, wherein when the temperatures of the tempering heat treatment and the post-weld heat treatment are different from each other, LMP according to the following Relationship 3 is 17.0 to 19.5, [Relational expression 3] T(logt eq +20)=T i (logt i +20) Total LMP = T i [log(t i +t eq )+20] Where T i : Kelvin temperature of PWHT, t i : PWHT time (hours), T: tempering Kelvin temperature, and t eq : Equivalent tempering time when converted to PWHT temperature (hours).

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