Steel material, weld heat affected zone and manufacturing method thereof

Abstract

The objective is to develop a weld heat affected zone that not only has excellent strength and low-temperature impact toughness after post-weld heat treatment (PWHT) but also has excellent low-temperature CTOD quality, and a method for manufacturing the same. [Solution] The present invention relates to a sintered body containing, by weight %, C: 0.05-0.20%, Si: 0.05-0.50%, Mn: 0.5-2.0%, Al: 0.005-0.10%, P: 0.010% or less, S: 0.015% or less, Nb: 0.001-0.070%, V: 0.001-0.30%, Ti: 0.001-0.030%, Cr: 0.01-1.0%, Mo: 0.01-1.0%, Cu: 0.01-0.60%, Ni: 1.0-4.0%, B: 0.001-0.005%, Ca: 0.0005-0.0040%, O: 0.001-0.005%, B: 0.001-0.005%, Ca: 0.0005-0.0040%, and O: 0.001-0.070%. The maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxide inclusions observed in the upper layer, which is the region from the surface to t / 4 (t is thickness, unit mm) in the thickness direction from the surface of the steel material, is 50 μm or less, and the unit area is 1 μm. 2 The number of winning items is 50 or less.

Description

[Technical Field]

[0001] The present invention relates to a steel material, a weld heat affected zone, and a manufacturing method thereof, and more particularly to a steel material, a weld heat affected zone, and a manufacturing method thereof, which are applicable to liquefied gas storage tanks and the like, and which have excellent strength and low-temperature impact toughness after post-weld heat treatment (PWHT) as well as excellent low-temperature CTOD quality. [Background technology]

[0002] Recently, as environmental regulations have become stricter, demand for liquefied gas, an environmentally friendly fuel, has increased, leading to an increase in the construction of liquefied gas carriers that transport liquefied gas and ships that use liquefied gas as fuel. Among the liquefied gases mentioned above, liquefied natural gas and liquefied ethylene / ethane gas generally use high-Ni steel because their liquefaction temperatures are very low. However, gases such as LPG, ammonia, and CO2 have liquefaction temperatures of around -60°C, and carbon steel with improved strength and low-temperature toughness is increasingly being used to reduce the construction costs of tanks and ships that store these gases.

[0003] Existing tanks for transporting liquefied gases mainly use large Type A tank designs that are not subject to high pressure, so there is no requirement for high strength or thick materials. However, recently, CO2 tanks or ammonia fuel tanks, which require high pressure during liquefaction, are designed as Type C, and there is a trend toward requiring thicker steel as well as higher strength. Furthermore, since the construction of Type C tanks requires Post Weld Heat Treatment (PWHT) as required by the IGC Code, the steel must have guaranteed physical properties after PWHT.

[0004] The PWHT process involves holding the material 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, degrading material properties such as strength and toughness, making quality assurance difficult. In particular, for steel materials for CO2 tanks, recent ship classifications require crack tip opening displacement (CTOD) quality in the heat-affected zone (HAZ) after PWHT as a structural fitness evaluation item to improve the stability of tank design.

[0005] On the other hand, in the case of liquefied CO2, the liquefaction temperature varies depending on the pressure, and therefore tanks can be broadly divided into low-pressure tanks and medium-pressure tanks.

[0006] In the case of low-pressure tanks with a typical operating pressure of 5-10 bar, the temperature must be lowered to -60°C to liquefy CO2, so steel with a tensile strength (TS) of 600-650 MPa is used for the storage tank, which can guarantee low-temperature impact toughness below -60°C. Maintaining the low temperature after liquefaction requires separate cooling equipment, which increases the initial installation and maintenance costs.

[0007] In the case of medium-pressure tanks with an operating pressure of 18-20 bar, the liquefaction temperature is as high as -35°C, which has the advantage of making it easier to store and transport liquefied CO2. However, due to the high operating pressure, ultra-high strength steel of 780 MPa or more is used for medium-pressure tanks, rather than the existing TS 600-650 MPa steel. To ensure high strength, steel containing high alloy elements is manufactured using Quenching-Tempering (QT) heat treatment, making it difficult to ensure the quality of the welds.

[0008] Patent Document 1 describes a steel sheet containing, 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 They state that by heating a slab consisting of Fe and impurities to 1100-1200°C, rolling it at 800-900°C with a cumulative reduction of 30% or more, air-cooling it, quenching it again at 910-930°C, and then tempering it at 650-680°C, it is possible to manufacture steel that can guarantee low-temperature impact toughness at -40°C.

[0009] The above steel provides a method for improving toughness by refining HAZ grains through the uniform distribution of Al-Ti-MnS composite oxidized inclusions in the weld heat affected zone. However, this does not guarantee the quality of the material after PWHT, and the tensile strength of the base material is low at 550-650 MPa, making it unsuitable for use in medium-pressure liquefied CO2 (LCO2).

[0010] On the other hand, Patent Document 2 describes a manufacturing method for P690QL2 steel material for LCO2 storage tanks, which contains, by weight, C: 0.09-0.12%, Mn: 0.9-1.3%, Ti: 0.015% or less, S: 0.002% or less, Nb: 0.06% or less, V: 0.08% or less, Ni: 1.5-2.5%, Cr: 0.2-0.4%, Mo: 0.3-0.4%, Cu: 0.3% or less, Al: 0.03-0.05%, Ca: 0.0005-0.005%, N: 0.005% or less, H: 0.0003% or less, the balance being Fe and impurities, has a thickness of 50 mm or less, and is capable of guaranteeing low-temperature impact toughness. The above Patent Document 2 states that if a material is heated at a temperature of 1160 to 1210°C, rolled at 780 to 880°C, quenched at 890 to 920°C, and tempered at a temperature of 630 to 680°C, a tensile strength of 780 MPa or more and an impact toughness of 47 J or more at -60°C can be ensured after PWHT.

[0011] However, the CTOD guaranteed temperature and standard of the HAZ portion do not reach the level required by the present invention, and no practical method is presented for use as a steel material for medium-pressure LCO2 storage tanks.

[0012] Patent Document 3 discloses a steel sheet containing, by 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, with the remainder being Fe and impurities, Mn / Ni being the ratio of the Mn and Ni contents being 0.80 or less, WES standard carbon equivalent Ceq being 0.430 to 0.900%, and tensile strength being 780 MPa or more, 930 MPa or less The steel plate can be manufactured with a yield strength of 630 MPa or less and 750 MPa or less, a yield ratio of 85% or less, impact toughness of the HAZ (heat-affected zone) at 0°C of 100 J or more, a thickness of 40 mm or more and 120 mm or less, and Vickers hardness measured at 225 or more points at 1 / 4 of the plate thickness on the surface. When the average of the smallest 20% of the Vickers hardness values ​​(Hvmin) and the average of the largest 20% of the Vickers hardness values ​​(Hvmax) is defined as Hvmin and Hvmax, respectively, the ratio Hvmin / Hvmax is 0.85 or less. While the strength meets the requirements for medium-pressure LCO2 steel, it may not be suitable because it does not guarantee impact toughness at temperatures near -40°C, the temperature at which CO2 liquefies. Furthermore, if the Ceq and hardness are too high, the martensite fraction during quenching may become too high, making it difficult to ensure impact toughness and CTOD characteristics in the HAZ.

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

[0014] Therefore, the above-mentioned prior art does not provide a practical method for manufacturing ultra-high strength steel for medium-pressure LCO2, which has a yield strength of 690 MPa or more, a tensile strength of 770 to 940 MPa after PWHT (Post Weld Heat Treatment), low-temperature impact toughness at -40°C or less, and a CTOD quality of 0.1 mm or more. [Prior art documents] [Patent documents]

[0015] [Patent Document 1] Korean Patent Publication No. 10-2022-7030534 [Patent Document 2] Chinese Patent Application Publication No. 11506158 [Patent Document 3] International Publication No. 2021 / 255858 Summary of the Invention [Problem to be solved by the invention]

[0016] The present invention aims to develop a weld heat affected zone that is excellent in strength and low-temperature impact toughness after post-weld heat treatment (PWHT) as well as low-temperature CTOD quality, and a method for manufacturing the same.

[0017] Another object of the present invention is to develop a steel material that is excellent in strength and low-temperature impact toughness after post-weld heat treatment (PWHT) and a method for producing the same.

[0018] The object of the present invention is not limited to the above-mentioned matters. Further object of the present invention is described in the overall content of the specification, and a person having ordinary skill in the art to which the present invention pertains will have no difficulty in understanding the further object of the present invention from the content described in the specification of the present invention. [Means for solving the problem]

[0019] The steel material of the present invention contains, by 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%, and O: 0. The steel material has a maximum oxide inclusion size of 50μm or less, with the remainder consisting of Fe and other unavoidable impurities, and the Ceq defined by the following relational expression 1 satisfies the range of 0.50 to 0.70, the microstructure contains tempered martensite at an area fraction of 50% or more, the packet size of the microstructure is 17μm or less, and the maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxide inclusions observed in the upper layer, which is the region from the surface to t / 4 (t is thickness, unit mm) in the thickness direction from the surface of the steel material, is 50μm or less, and the unit area is 1μm 2 The number of winning items is 50 or less.

[0020] [Equation 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15

[0021] In the above Relational Formula 1, [C], [Mn], [Cr], [Mo], [V], [Ni], and [Cu] respectively represent the contents (wt%) of C, Mn, Cr, Mo, V, Ni, and Cu contained in the steel, and when these elements are not intentionally added, 0 is substituted.

[0022] The microstructure may include bainite and other unavoidable structures.

[0023] When the tempered martensite content exceeds 80%, the Brinell hardness of the cross section of the steel material is 220 to 290HB.

[0024] The steel material has a structure in which, among precipitates observed in a cross section of the steel material, VC or VCN precipitates having a diameter of 5 to 15 nm are present in a size of 1 μm or less.2 There are 5 or more winning pieces.

[0025] The microstructure is observed within prior austenite grains having an average grain size of 30 μm or less.

[0026] The steel material has a yield strength of 690 MPa or more, a tensile strength of 770 to 940 MPa, and a Charpy impact absorption energy at -40°C of 50 J or more.

[0027] The method for producing a steel material of the present invention satisfies the above-mentioned composition and Ceq, and the maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxide inclusions observed in the upper layer, which is the region from the surface of the steel slab to t / 4 (t is thickness, unit mm) in the thickness direction, is 50 μm or less, and the maximum size of the CaO-Al2O3-X complex oxide inclusions ... 2 the steel slab to a temperature range of 820 to 950°C, and hold the temperature for 10 to 40 minutes, and then cool the steel to a temperature of 300°C or less at a cooling rate of 3°C / s to 80°C / s, based on a temperature standard of t / 4 of the thickness of the steel; and after the cooling, a tempering heat treatment step of heating the steel to 550 to 700°C and holding the temperature for 5 to 60 minutes.

[0028] The step of producing the steel slab includes a step of adding a metal Ca wire to the molten steel after secondary refining so that the Ca addition amount is 0.015 to 0.15 kg / ton, and a step of performing clean bubbling for 5 to 40 minutes after adding the Ca so that the inert gas injection amount into the ladle is 10 to 50 liters / min, wherein the Ca wire injection speed is 100 to 300 meters / min, and there are two inert gas injection points in the ladle.

[0029] After the hot rolling, the method may further include a step of multi-stage stack cooling in a temperature range of 200° C. or more to room temperature.

[0030] The step of performing a post-weld heat treatment on the steel material may further be included.

[0031] The post-weld heat treatment can be carried out by holding the temperature at 595°C to 625°C for 180 to 540 minutes (min).

[0032] When the tempering heat treatment and the post-weld heat treatment are performed at the same temperature, the LMP according to the following relational expression 2 is 17.0 to 19.5.

[0033] [Equation 2] LMP=T k (Log t+20)

[0034] where T k = Kelvin temperature of tempering and PWHT, t = tempering and PWHT time (hours)

[0035] When the tempering heat treatment and the post-weld heat treatment are performed at different temperatures, the LMP according to the following relational expression 3 is 17.0 to 19.5.

[0036] [Equation 3] T(logt eq +20)=T i (logt i +20) Total LMP=T i [log(t i +t eq )+20]

[0037] where T i : PWHT temperature in Kelvin, t i : PWHT time (hours), T: Tempering temperature in Kelvin, t eq : Equivalent tempering time (hours) when converted to PWHT temperature

[0038] The weld heat affected zone of the present invention satisfies the above-mentioned composition and Ceq, and has a maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxide inclusions of 50 μm or less in maximum size and 1 μm in unit area, observed in the upper layer, which is the region from the surface to t / 4 (t is thickness, unit mm) in the thickness direction from the surface. 2 The number of particles per grain is 50 or less, the grain boundary occupancy ratio of cementite present at the grain boundaries is 10% or less, the average size of cementite is 1 μm or less, and the hardness of the cross section is 200 to 300HB.

[0039] [Equation 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15

[0040] In Relational Formula 1, [C], [Mn], [Cr], [Mo], [V], [Ni], and [Cu] represent the contents (wt%) of C, Mn, Cr, Mo, V, Ni, and Cu contained in the steel, respectively. If these elements are not intentionally added, 0 is substituted.

[0041] The method for producing a weld heat affected zone of the present invention satisfies the above-mentioned composition and Ceq, and the maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxidized inclusions observed in the upper layer, which is the region from the surface to t / 4 (t is thickness, unit mm) in the thickness direction from the surface, is 50 μm or less, and the maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxidized inclusions is 50 μm or less, and the maximum size of CaO-Al2O3-X ... 2 The method includes the steps of preparing steel materials having 50 or less pieces per unit; welding the steel materials; and post-welding heat treating the welded steel materials. [Effects of the Invention]

[0042] According to the present invention, it is possible to provide a steel material having excellent strength and low-temperature impact toughness with minimal deterioration in physical properties even after post-weld heat treatment (PWHT), and a method for manufacturing the same.

[0043] Furthermore, the present invention can provide a weld heat affected zone that is excellent in low-temperature impact toughness and low-temperature CTOD characteristics after post-weld heat treatment (PWHT), and a method for producing the same.

[0044] The various beneficial advantages and effects of the present invention are not limited to the above-mentioned contents, but can be more easily understood in the course of describing specific embodiments of the present invention. DETAILED DESCRIPTION OF THE INVENTION

[0045] The terminology used herein is for the purpose of describing the invention and is not intended to limit the invention. Furthermore, as used herein, the singular forms "a," "an," and "the" include the plural forms unless the relevant definition clearly dictates otherwise.

[0046] The meaning of "comprises" as used herein is to specify features and does not exclude the presence or addition of other features.

[0047] Unless otherwise defined, all terms, including technical and scientific terms, used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this invention belongs. Dictionary-defined terms are to be interpreted as having a meaning consistent with the relevant technical literature and the presently disclosed content.

[0048] The following describes preferred embodiments of the present invention. The embodiments of the present invention can be modified in various forms, and the scope of the present invention should not be construed as being limited to the embodiments described below. The embodiments are provided to further explain the present invention to those skilled in the art.

[0049] First, the composition of the steel material of the present invention will be described in detail.

[0050] In weight percent, it contains C: 0.05-0.20%, Si: 0.05-0.50%, Mn: 0.5-2.0%, Al: 0.005-0.10%, P: 0.010% or less, S: 0.0050% or less, Nb: 0.001-0.070%, V: 0.001-0.30%, Ti: 0.001-0.030%, Cr: 0.01-1.0%, Mo: 0.01-1.0%, Cu: 0.01-0.60%, Ni: 1.0-4.0%, B: 0.001-0.005%, Ca: 0.0005-0.0040%, O: 0.002% or less, and the remainder is Fe and other unavoidable impurities.

[0051] Ceq defined by relational expression 1 is 0.50 to 0.70.

[0052] [Equation 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15

[0053] In Relational Formula 1, [C], [Mn], [Cr], [Mo], [V], [Ni], and [Cu] represent the contents (wt%) of C, Mn, Cr, Mo, V, Ni, and Cu contained in the steel, respectively. If these elements are not intentionally added, 0 is substituted.

[0054] The composition of the above steel material is the same as that of not only the steel slab, but also the base material of the steel material before welding, the steel material after welding, and the post-weld heat-treated steel material, as well as the composition of the weld heat-affected zone (HAZ).

[0055] The above composition will be explained in detail below. Unless otherwise specified, % and ppm shown in the alloy composition are based on weight.

[0056] Carbon (C): 0.05~0.20% Carbon (C) is the most important element for ensuring basic strength, so it must be contained within an appropriate range in steel. To achieve this effect, 0.05% or more of carbon (C) can be added. Preferably, 0.10% or more of carbon (C) can be added. However, if the carbon (C) content exceeds a certain level, the martensite fraction becomes excessively high, increasing the hardness of the weld heat-affected zone (HAZ), and the grain boundary cementite fraction increases, reducing impact propagation resistance during impact tests and CTOD tests, making it impossible to achieve the toughness value required in the present invention. Therefore, in the present invention, the carbon (C) content can be set to 0.20%, with a more preferred upper limit of 0.18%.

[0057] Silicon (Si): 0.05 to 0.50% Silicon (Si) is a substitutional element that improves the strength of steel through solid solution strengthening and has a strong deoxidizing effect, making it an essential element in the production of clean steel. Therefore, silicon (Si) may be contained in an amount of 0.05% or more, preferably 0.20% or more. However, if silicon (Si) is contained in a large amount, it may form an island martensite (MA, martensite-austenite) phase, excessively increasing the matrix strength of the martensite phase and potentially deteriorating low-temperature impact toughness. Therefore, the upper limit of silicon (Si) content can be set at 0.50%. A more preferred upper limit of silicon (Si) content is 0.40%.

[0058] Manganese (Mn): 0.5 to 2.0% Manganese (Mn) is a useful element that improves strength through solid solution strengthening and improves hardenability by forming low-temperature transformation phases. Therefore, to ensure a yield strength of 690 MPa or more, a manganese (Mn) content of 0.5% or more is preferred. A manganese (Mn) content of 0.7% or more is more preferred. However, manganese (Mn) can form elongated non-metallic inclusions, MnS, together with sulfur (S), which can reduce toughness and act as impact initiation points, thereby rapidly reducing the low-temperature impact toughness of steel. Therefore, the manganese (Mn) content is preferably controlled to 2.0% or less, and more preferably 1.5% or less.

[0059] Aluminum (Al): 0.005 to 0.10% Aluminum (Al), along with silicon (Si), is one of the powerful deoxidizing agents in the steelmaking process, and to achieve this effect, a content of 0.005% or more is preferred. A more preferred lower limit for the aluminum (Al) content is 0.01%. However, excessive aluminum (Al) content can lead to excessive increases in the Al2O3 fraction in oxidized inclusions formed as a result of deoxidation, resulting in coarse inclusion size and making removal of the inclusions difficult during refining, which can result in reduced low-temperature impact toughness. Therefore, it is preferred to control the aluminum (Al) content to 0.10% or less. A more preferred aluminum (Al) content is 0.07% or less.

[0060] Phosphorus (P): 0.010% or less (including 0%) and sulfur (S): 0.0050% or less (including 0%) Phosphorus (P) and sulfur (S) are elements that induce embrittlement at grain boundaries or form coarse inclusions, which induce embrittlement. Therefore, it is preferable that they are not contained. In order to improve brittle crack propagation resistance, it is preferable that the phosphorus (P) content be 0.010% or less and the sulfur (S) content be 0.0050% or less.

[0061] Niobium (Nb): 0.001 to 0.070% Niobium (Nb) is an element that precipitates in the form of NbC or NbCN to improve the strength of the base material. Furthermore, niobium (Nb) that dissolves during high-temperature reheating precipitates very finely in the form of NbC during rolling, suppressing austenite recrystallization and thus having the effect of refining the structure. Therefore, niobium (Nb) is preferably contained in an amount of 0.001% or more, and a more preferred niobium (Nb) content is 0.005% or more. On the other hand, if niobium (Nb) is contained in excess, undissolved niobium (Nb) is generated in the form of TiNb(C,N), which can impair low-temperature impact toughness. Therefore, the upper limit of the niobium (Nb) content is preferably 0.070%. A more preferred upper limit of the niobium (Nb) content is 0.065% or less.

[0062] Vanadium (V): 0.001 to 0.30% Since most of the vanadium (V) dissolves again during reheating, its strengthening effect due to precipitation or dissolution during subsequent rolling is minimal. However, it precipitates as very fine carbonitrides during tempering and PWHT, thereby improving strength. To fully achieve this effect, 0.001% or more vanadium (V) must be added. A more preferable lower limit for the vanadium (V) content is 0.01%. However, excessive vanadium (V) content not only increases the surface hardness of the slab excessively due to its high hardening ability, which may cause surface cracks during flanging, but also sharply increases production costs, making it commercially unprofitable. Therefore, the vanadium (V) content can be limited to 0.3% or less. A more preferable vanadium (V) content is 0.25% or less.

[0063] Titanium (Ti): 0.001 to 0.030% Titanium (Ti) precipitates as TiN during reheating and inhibits the growth of prior austenite grains at high temperatures, significantly improving low-temperature toughness. To achieve this effect, a titanium (Ti) content of 0.001% or more is preferred. However, excessive titanium (Ti) content can lead to clogging of the continuous casting nozzle and crystallization in the center, reducing low-temperature toughness. Furthermore, titanium (Ti) combines with nitrogen (N) to form coarse TiN precipitates in the center of the thickness, reducing the elongation of the product. This can reduce uniform elongation during the forging process and cause surface cracks. Therefore, the titanium (Ti) content is 0.030% or less. The preferred upper limit of the titanium (Ti) content is 0.025% or less, with a more preferred titanium (Ti) content being 0.018% or less.

[0064] Chromium (Cr): 0.01 to 1.0% Chromium (Cr) is an element that increases the yield strength and tensile strength by increasing hardenability and forming a low-temperature transformation structure. It is also an element that has the effect of preventing a decrease in strength by slowing down the spheroidization rate of cementite. For this effect, 0.01% or more of chromium (Cr) can be contained. On the other hand, if the chromium (Cr) content is excessive, M 23 The size and fraction of Cr-rich coarse carbides such as C6 increase, reducing the impact toughness of the steel, and the solid solubility of niobium (Nb) in the steel and the fraction of fine precipitates such as NbC decrease, potentially resulting in a decrease in product strength. Therefore, in the present invention, the upper limit of the chromium (Cr) content can be set to 1.0%. The preferred upper limit of the chromium (Cr) content is 0.8%.

[0065] Molybdenum (Mo): 0.01 to 1.0% Molybdenum (Mo) is an element that increases grain boundary strength and has a significant effect on solid solution strengthening in ferrite, effectively contributing to increased strength and ductility of products. Molybdenum (Mo) also has the effect of preventing a decrease in toughness due to the 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. Therefore, the upper limit of the Mo content is preferably 1.0%.

[0066] Copper (Cu): 0.01 to 0.60% Copper (Cu) is an advantageous element in the present invention because it not only significantly improves the strength of the matrix phase through solid solution strengthening in ferrite, but also has the effect of suppressing corrosion in a wet hydrogen sulfide atmosphere. To achieve these effects, copper (Cu) can be contained in an amount of 0.01% or more. A more preferred copper (Cu) content is 0.03% or more. However, excessive copper (Cu) content increases the possibility of inducing star cracks on the surface of the steel sheet, and since copper (Cu) is an expensive element, this can lead to problems such as significantly increasing manufacturing costs. Therefore, the upper limit of the copper (Cu) content is preferably 0.60%, and more preferably 0.35%.

[0067] Nickel (Ni): 1.0-4.0% Nickel (Ni) is an element that effectively contributes to improving strength by increasing stacking faults at low temperatures and facilitating dislocation cross-slip, thereby improving impact toughness and hardening ability. To achieve these effects, nickel (Ni) can be contained in an amount of 1.0% or more. A preferred nickel (Ni) content is 0.10% or more. However, excessive nickel (Ni) addition can increase manufacturing costs due to its high cost. Therefore, the upper limit of the nickel (Ni) content is preferably 4.0%, and more preferably 3.0%.

[0068] Calcium (Ca): 0.0005-0.0040% Adding calcium (Ca) after deoxidation with aluminum (Al) inhibits the formation of MnS by combining with sulfur (S), which forms MnS inclusions, and also inhibits the occurrence of cracks due to hydrogen-induced cracking by forming spherical CaS. To ensure that sulfur (S) contained as an impurity is fully converted into CaS, calcium (Ca) content of 0.0005% or more is preferred. However, if the amount of calcium added is excessive, the calcium (Ca) remaining after forming CaS combines with oxygen (O) to form coarse oxidized inclusions, which are elongated and fractured during rolling, potentially reducing lamellar tear resistance. Therefore, the upper limit of calcium (Ca) content is preferably 0.0040%.

[0069] Boron (B): 0.001 to 0.005% B is an element that is effective in improving the hardenability of hot-rolled steel sheets by delaying the transformation of austenite to ferrite during cooling transformation. To achieve this effect, it is preferable to add 0.001% or more. However, if the content exceeds 0.005%, the strength increases excessively, which can lead to a decrease in impact toughness. In addition, boron oxides can be formed, which can degrade the surface quality of the steel sheet.

[0070] Oxygen (O): 0.002% or less Oxygen combines with Ca and Al in molten steel to form CaO, Al2O3, or CaO-Al2O3 composite oxide inclusions. Oxide inclusions formed at high temperatures are crushed and elongated in the rolling direction during the rough rolling of slabs, remaining undissolved not only during the rolling / heat treatment process of the base metal but also in the HAZ during welding. Because the edges of crushed oxide inclusions can act as initiation points for fracture during impact and CTOD evaluation due to the notch effect, they are preferably removed by flotation separation in the secondary refining process. Therefore, the oxygen concentration in molten steel is preferably 20 ppm or less. The present invention proposes a secondary refining method that can minimize the amount of oxide inclusions. After applying this process, the total oxygen content is preferably 15 ppm or less, and most preferably 10 ppm or less.

[0071] In addition to the above components, the remainder consists of Fe and other inevitable impurities. However, since unintended impurities may be unavoidably mixed in from raw materials or the surrounding environment during normal manufacturing processes, it is not possible to completely eliminate them. These impurities are known to anyone with ordinary skill in the art, and therefore, the entire contents of these impurities are not specifically mentioned in this specification.

[0072] Furthermore, the addition of other effective ingredients in addition to the ingredients mentioned above is not completely excluded.

[0073] The composition has a Ceq of 0.50 to 0.70 as defined by the following relational expression 1.

[0074] [Equation 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15

[0075] In Relational Formula 1, [C], [Mn], [Cr], [Mo], [V], [Ni], and [Cu] represent the contents (wt%) of C, Mn, Cr, Mo, V, Ni, and Cu contained in the steel, respectively. If these elements are not intentionally added, 0 is substituted.

[0076] If Ceq according to Relation 1 is less than 0.50, the martensite fraction decreases, making it impossible to ensure the high yield strength required in the present invention, and the bainite fraction increases excessively, increasing the proportion of grain boundary cementite in the weld heat affected zone (HAZ), particularly the coarse grained heat affected zone (CGHAZ), after post-weld heat treatment (PWHT), making it impossible to ensure the appropriate impact toughness and HAZ CTOD required in the present invention. Furthermore, if Ceq exceeds 0.70, the hardness becomes too high, potentially inducing cleavage fracture, reducing crack propagation resistance and making it difficult to ensure the appropriate impact toughness and CTOD properties in the weld heat affected zone.

[0077] The steel material of the present invention has a maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxide inclusions observed in the upper layer, which is the region from the surface of the steel material to t / 4 (t: thickness) in the thickness direction, of 50 μm or less in a unit area of ​​1 μm. 2 The number of winning pieces can be 50 or less.

[0078] Complex oxide inclusions can be observed in the base metal and heat-affected zone (HAZ) of steel slabs as well as pre-welded, welded, and post-weld heat-treated steel.

[0079] When complex oxide inclusions are formed during the steel slab manufacturing process, their shape and fraction do not change during the subsequent product manufacturing process, but they can be crushed during rough rolling and elongate in the rolling direction. Crushed oxide inclusions can act as points where stress concentrates and can act as initiation points for fractures. Therefore, if the amount or size of oxide inclusions generated during the steel slab manufacturing process is large, the number of crushed oxide inclusions after rough rolling will increase, which is the main cause of reduced impact toughness and CTOD quality. Therefore, if the size of complex oxide inclusions exceeds 50 μm or is less than 1 μm in the upper layer of the product, which is the surface to t / 4 region, 2If the number of particles per unit area exceeds 50, it may be difficult to ensure low-temperature impact toughness at -40°C and HAZ CTOD quality at temperatures below -35°C based on the base material.

[0080] The steel material of the present invention has a base structure of tempered martensite or a mixed structure of tempered martensite and bainite. Specifically, the steel material contains 50% or more tempered martensite in terms of area fraction, with the remainder containing bainite and an unavoidably occurring structure.

[0081] When the matrix structure is a mixed structure of tempered martensite and bainite, the area fraction of tempered martensite is 50 to 80%.

[0082] When the base structure is composed of a single structure of more than 80% tempered martensite, the hardness of the cross section of the base material of the steel after PWHT is 220 to 290HB on the Brinell scale.

[0083] The structure that inevitably occurs may be island martensite (MA) or the like.

[0084] The packet size of the tempered martensite or bainite low-temperature structure is 17 μm or less.

[0085] The average grain size of the prior austenite in the steel is 30 μm or less.

[0086] Fine VC precipitates with a diameter of 5 to 15 nm or VCN precipitates with a diameter of 1 μm observed in the cross section of the steel material 2 Each precipitate may contain five or more precipitates.

[0087] The microstructural characteristics may be similar not only in the steel before welding and in the base material after welding, but also in the base material of a steel that has been heat treated after welding.

[0088] The average grain size of prior austenite in steel is 30 μm or less. Postweld heat treatment (PWHT) is typically performed at temperatures above 595°C to prevent cracking in the weld zone through stress relief after welding. This reduces dislocation density and potentially reduces cross-sectional hardness, thereby improving the integrity of the weld. However, during prolonged high-temperature holding, cementite formed at the bainite lath interface decomposes, resulting in long-range diffusion of carbon to the austenite grain boundaries. As a result, if the grain size increases, the grain boundary area available for cementite decreases, increasing segregation and potentially degrading impact toughness. Therefore, it is necessary to control the grain size of prior austenite under appropriate low-temperature and high pressure conditions. Therefore, the prior austenite grain size is preferably 30 μm or less, and more preferably 25 μm or less.

[0089] As the PWHT temperature and time increase, coarse cementite becomes spheroidized at prior austenite grain boundaries, with some forming as film-like cementite. When such cementite forms at grain boundaries, the grain boundary strength becomes lower than the shear strength, ultimately becoming the primary cause of intergranular fracture. Therefore, to prevent intergranular fracture after PWHT, an appropriate fraction of martensite, which undergoes only short-range diffusion, must be present. However, if the martensite fraction is too high, the cross-sectional hardness increases, resulting in low impact energy absorption even if transgranular fracture occurs, ultimately resulting in reduced toughness. Therefore, the microstructure of the steel is preferably composed of a mixed structure of tempered martensite and bainite, with the tempered martensite fraction preferably being 50-80% by area.

[0090] However, when the microstructure of the steel material is tempered martensite, i.e., when the martensite content exceeds 80% (including 100%), the tempering temperature must be increased to maintain the cross-sectional hardness of the base material at 200 to 290HB. If the cross-sectional hardness falls outside this range, it becomes difficult to ensure the low-temperature impact toughness and CTOD value required by the present invention.

[0091] The packet size of the low-temperature structure of tempered martensite or bainite is 17 μm or less. Packet size is a grouping unit of laths with the same orientation in low-temperature transformation structures such as bainite or martensite. It can be understood as a domain boundary that can physically resist crack propagation (i.e., the property of cracks breaking at the boundary when the crack propagates). For example, packet size can be measured at grain boundaries, which can be confirmed by using EBSD (Electro-Backscattered Diffraction) equipment when domains with an orientation relationship of within 15 degrees are set as having the same orientation. Bainite and martensite grow only within prior austenite grains, so their size is smaller than that of the prior austenite grains. A packet size of 17 μm or less is preferable because it is difficult to ensure sufficient crack propagation resistance and therefore sufficient impact toughness when the packet size exceeds 17 μm.

[0092] On the other hand, the tempering and PWHT processes reduce the dislocation density of steel, resulting in a decrease in strength. The use of fine precipitates to appropriately compensate for this can reduce the strength degradation phenomenon during heat treatment. Typically, normalized and thermo-mechanical control processed (TMCP) steels utilize precipitates such as NbC and NbCN. However, these precipitates are formed at temperatures between 800 and 900°C and are formed through deformation induction, making them unsuitable for quenching-tempering (QT) heat-treated steels. Therefore, vanadium carbides / carbonitrides such as VC and VCN, which are formed in the temperature range of 550 to 700°C, are more readily utilized. Fine VC or VCN precipitates with a diameter of less than 5 nm may be too small to provide much of a strength improvement, while those with a diameter greater than 15 nm are excessively coarse, resulting in a reduced strength improvement effect. Therefore, it is appropriate that the size of fine precipitates is 5 to 15 nm. On the other hand, when the density of VC and VCN is 1 μm 2 If the number of fine precipitates per grain is less than five, the precipitation strengthening effect described above is reduced, so five or more is preferred. The formation of precipitates is determined by the chemical potential and temperature, which are determined by the steel composition. Although a large number of precipitates of 5 to 15 nm in size may occur, as the number increases, the precipitates agglomerate and increase in size, preventing the formation of 5 to 15 nm in size indefinitely. Therefore, the fact that there is no upper limit on the number of fine precipitates is not necessarily obvious to those skilled in the art.

[0093] Meanwhile, the microstructure of the weld heat affected zone (HAZ) of the present invention will be described in detail. The HAZ is the HAZ after post-weld heat treatment (PWHT), and particularly the CGHAZ (Coarse Grained Heat Affected Zone).

[0094] In the heat-affected zone (HAZ), the grain boundary occupancy ratio of cementite at grain boundaries is 10% or less. The grain boundary occupancy ratio can be understood as the length ratio of film-like cementite that occupies the grain boundary, and in three dimensions, it is the area ratio. As an example of occupancy ratio measurement, a test piece is nital etched and then observed using a SEM, etc., to distinguish between general grain boundaries and grain boundaries where film-like cementite has formed.

[0095] On the other hand, the average size of cementite is 1 μm or less.

[0096] The microstructure of the HAZ, particularly the CGHAZ, after post-weld heat treatment (PWHT) preferably has an area fraction of tempered martensite of 50% or more. The remainder, in addition to tempered martensite, is bainite and other structures. If the area fraction of tempered martensite is less than 50%, the proportion of film-like cementite at the grain boundaries after PWHT increases, resulting in coarsening of the crystal grains, making it difficult to ensure an appropriate CTOD value of 0.1 mm or more at temperatures below -35°C. For the same reason, the grain boundary proportion of cementite present at the grain boundaries after PWHT is preferably 10% or less, and the average size of the cementite present at the grain boundaries is preferably 1 μm or less.

[0097] The average grain size of prior austenite in the weld heat-affected zone (HAZ) is 200 μm or less. The microstructure of the HAZ is determined by the microstructure of the base metal and welding conditions such as welding heat input, preheat temperature, and lamination temperature. In particular, the CGHAZ is a region where grains become significantly coarser than the base metal, potentially resulting in material degradation compared to the base metal. Based on the CGHAZ, the average grain size of prior austenite is preferably 200 μm or less. Exceeding this limit excessively increases hardenability and cross-sectional hardness, so the average grain size of the CGHAZ is preferably 200 μm or less. The average grain size of prior austenite may not change before and after PWHT. To change the grain size, reverse transformation to the austenite single-phase region is required, requiring a temperature increase of 800°C or more. However, since the PWHT temperature is in the 500-600°C range, there is no likely to be any change in the AGS itself.

[0098] In the heat-affected zone (HAZ), the grain boundary occupancy ratio of cementite present at grain boundaries is 10% or less. The grain boundary occupancy ratio can be understood as the length ratio of film-like cementite that occupies the grain boundary, and in three dimensions, it is the area ratio. As an example of occupancy ratio measurement, a test piece is nital etched and then observed using a SEM, etc., to distinguish between ordinary grain boundaries and grain boundaries where film-like cementite has formed.

[0099] The steel material of the present invention can ensure the strength and low-temperature impact toughness of the base material after PWHT. Specifically, the steel material has a yield strength of 690 MPa or more, a tensile strength of 770 to 940 MPa, and a low-temperature impact toughness absorption energy of 50 J or more in a Charpy V-Notch test performed at -40°C or below.

[0100] Meanwhile, the weld heat-affected zone (CGHAZ) after PWHT, especially the CGHAZ, can maintain excellent low-temperature impact toughness and CTOD properties. Specifically, in a Charpy V-Notch test conducted at temperatures below -40°C, the low-temperature impact toughness absorbed energy can be 40 J or more, and the CTOD property at temperatures below -35°C can be 0.1 mm or more.

[0101] An example of the method for producing a steel material according to the present invention will be described in detail below.

[0102] The manufacturing method involves manufacturing a steel slab that satisfies the above-mentioned alloy composition and the range of Ceq in Relational Equation 1, and then heating, hot rolling, reheating, quenching, and tempering the steel slab to manufacture the steel material. This involves a reheating-quenching-tempering process, which can be called the RQT method.

[0103] The steel material thus manufactured can be subjected to welding and post-weld heat treatment (PWHT).

[0104] Each step will be described in detail below.

[0105] Steel slab manufacturing The alloy composition and Ceq defined by the relational formula 1 described above satisfy the range of 0.50 to 0.70, and the maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) complex oxide inclusions observed in the upper layer, which is the region from the surface of the steel slab to t / 4 (t: thickness, unit: mm) in the thickness direction, is 50 μm or less, and the unit area is 1 μm 2 Produce steel slabs with 50 or fewer pieces per slab.

[0106] As an example of a method for producing steel slabs, a calcium (Ca) raw material may be added to secondary refined molten steel and a cleaning bubbling process may be performed. As a specific example, after secondary refinement, metal Ca wire may be added to the molten steel so that the amount of Ca added is 0.015 to 0.15 kg / ton. The metal Ca wire is made of steel surrounding a Ca alloy, and the thickness of the steel may be 1.2 to 1.4 mm, and the Ca wire addition speed is 100 to 300 meters / min.

[0107] After adding Ca, cleaning bubbling can be performed in the ladle so that the inert gas is blown into the ladle at a rate of 10 to 50 liters / min. There may be two inert gas blowing points in the ladle, and the cleaning bubbling time is 5 to 40 minutes.

[0108] In the present invention, the process before the secondary refining is not particularly limited, and any conventional method can be applied. In the present invention, the process after the Ca addition and cleaning bubbling process is also not particularly limited, and slabs can be produced by cooling the molten steel under conventional conditions. Furthermore, if the amount of Al2O3 in the molten steel increases, the generation and coarsening of inclusions containing both Ca and Al progresses, and the number of crushed inclusions increases during rolling, making it impossible to ensure the low-temperature impact toughness and CTOD quality of the HAZ. Therefore, the total amount of inclusions in the molten steel before Ca addition can be limited to 2 to 5 ppm.

[0109] The Ca introduction step and the cleaning bubbling step will be described in detail below.

[0110] (1) Addition of calcium (Ca) When manufacturing steel slabs, metal Ca wire can be added to the molten steel after secondary refining so that the Ca input amount is 0.015 to 0.15 kg / ton. Metal Ca wire is composed of steel surrounding a Ca alloy, and the thickness of the steel is 1.2 to 1.4 mm, and the input speed of the Ca wire is 100 to 300 meters / min.

[0111] When Ca is added, if the thickness of the steel surrounding the Ca alloy in the metal Ca wire is less than 1.2 mm, the Ca melts at the top of the ladle, reducing the effect of ferrostatic pressure, reducing the Ca recovery rate and increasing the amount added.On the other hand, if the thickness exceeds 1.4 mm, the Ca wire contacts the base of the ladle, causing problems such as melting and damage to the ladle's refractory, making it impossible to ensure stable operation.

[0112] When adding Ca to molten steel, the speed at which the wire is added to the molten steel must be controlled to ensure a high Ca recovery rate, along with the thickness of the metallic Ca wire. If the wire is added at a speed less than 100 meters / min, the Ca melts at the top of the ladle, reducing the effect of ferrostatic pressure, resulting in a poor Ca recovery rate and an increased amount of wire added. On the other hand, if the wire is added at a speed greater than 300 meters / min, the Ca wire may reach the base of the ladle, causing problems such as melting and damage to the ladle's refractory, resulting in operational stability being compromised. The preferred wire addition speed is 120 to 300 meters / min, and more preferably 140 to 180 meters / min.

[0113] If the amount of Ca added is too small, MnS is generated in the center during solidification, making it impossible to ensure low-temperature impact toughness due to segregation defects, and if the amount is excessive, it reacts with Al2O3 in the refractory, accelerating the erosion of the refractory and making it impossible to ensure stable operation. Therefore, taking the above-mentioned problems into consideration, the amount of Ca added can be set to 0.015 to 0.15 kg / ton, preferably 0.015 to 0.10 kg / ton, and more preferably 0.050 to 0.10 kg / ton.

[0114] (2) Clean bubbling After adding Ca, cleaning bubbling can be carried out for 5 to 40 minutes so that the inert gas injection rate into the ladle is 10 to 50 liters / min. There are two points in the ladle where inert gas is injected.

[0115] If the inert gas injection rate into the ladle is too low, the amount of Al2O3 clusters and composite inclusions containing both Ca and Al that adhere to and are removed by the inert gas will be reduced, resulting in a deterioration in cleanliness and the inability to ensure impact toughness in the HAZ. On the other hand, if the amount is excessive, the stirring force will be strong, causing bare metal on the molten steel surface and slag contamination, resulting in a deterioration in cleanliness. Therefore, the inert gas injection rate may be 10 to 50 liters / min, preferably 15 to 40 liters / min, and more preferably 20 to 30 liters / min.

[0116] On the other hand, if there is only one inert gas injection point in the ladle, non-uniform regions will exist in the molten steel, and the ability to remove Al2O3 clusters and composite inclusions containing both Ca and Al will be reduced. If there are three or more points, overlapping areas will occur during gas injection, increasing the stirring force, causing bare metal on the molten steel surface and slag contamination, resulting in reduced cleanliness.

[0117] If the cleaning bubbling time is too short despite limiting the amount of inert gas injected into the ladle, the amount of Al2O3 clusters and composite inclusions containing both Ca and Al that adhere to and are removed by the inert gas will be reduced, resulting in a deterioration in cleanliness and failure to ensure toughness. On the other hand, if the time is too long, the temperature in the molten steel will drop significantly, creating a temperature gradient in the ladle and potentially deteriorating cleanliness. The cleaning bubbling time can be 5 to 40 minutes, preferably 7 to 17 minutes, and more preferably 10 to 15 minutes.

[0118] Steel slab heating The steel slab can be heated to a temperature range of 1050 to 1300°C.

[0119] The heating is preferably performed at a slab heating temperature of 1150°C or higher to redissolve Ti or Nb carbonitrides or TiNb(C,N) coarse crystals formed during casting, and maximize the austenite grain size by heating and maintaining the austenite at or above its recrystallization temperature after sizing rolling. However, if the temperature is too high, problems may occur due to oxide scale at high temperatures, and excessive grain growth may cause deterioration of low-temperature impact toughness. Therefore, the upper limit of the heating temperature is 1300°C.

[0120] hot rolling The heated steel slab can be hot rolled to a finish hot rolling temperature of 830 to 1050°C.

[0121] 830°C corresponds to the region above the recrystallization temperature. Complex inclusions formed during the refining process must accommodate deformation during rolling as the strength of the steel increases with lower rolling temperatures. As a result, they may be crushed or segmented into smaller inclusions, or inclusions such as MnS may become elongated. These small inclusions, which are crushed or segmented or elongated, act as the direct cause of the initiation and propagation of low-temperature impact toughness. Therefore, it is preferable to perform finish hot rolling at 830°C or above, which is the recrystallization temperature at which work hardening does not occur. However, if the temperature exceeds 1050°C, austenite grain growth continues even after rolling, potentially reducing impact toughness due to an increase in the Ductile-Brittle Transition Temperature (DBTT).

[0122] The thickness of the steel sheet after finish hot rolling may be 5 to 100 mm, preferably 5 to 70 mm, and more preferably 5 to 50 mm.

[0123] Steel sheets manufactured 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 the growth of precipitates and shorten the process time.

[0124] In the present invention, when the amount of dissolved hydrogen in the molten steel during slab production is 1.3 ppm or more, the method may further include a step of multi-stage cooling from a temperature of 200°C or more to room temperature based on the surface temperature after hot rolling before reheating heat treatment. When multi-stage cooling is performed, dissolved hydrogen in the steel is released, thereby more effectively suppressing internal microcracks caused by hydrogen, and ultimately improving the low-temperature impact toughness of the base metal and the weld heat-affected zone.

[0125] reheating After hot rolling, the steel sheet can be reheated by heating it to a temperature of 820 to 950°C and holding it there for 10 to 40 minutes.

[0126] If the reheating temperature is less than 820°C or the holding time is less than 10 minutes, carbides formed during cooling after rolling and impurity elements segregated at grain boundaries may not be redissolved smoothly and may remain after cooling, which may be the main cause of deterioration in low-temperature impact toughness. On the other hand, if the reheating temperature exceeds 950°C or the holding time exceeds 40 minutes, the grain size of the prior austenite increases excessively due to grain growth, which may be the main cause of deterioration in toughness due to an increase in DBTT. Therefore, the reheating temperature is preferably 820 to 950°C, and more preferably 850 to 890°C.

[0127] Quenching & Tempering The reheated steel plate can be subjected to quenching and tempering.

[0128] The cooling rate during quenching is one of the most important process factors that determines the microstructure of the base material. In the present invention, after reheating, the steel can be cooled to 300°C or below at a cooling rate of 3°C / s to 80°C / s, based on the temperature at t / 4 of the thickness (t: thickness, unit: mm). If the cooling rate ends at 300°C or above, the microstructure of the base material may contain island martensite (MA, martensite-austenite constituent) due to incomplete cooling transformation. This does not dissolve in the subsequent tempering and PWHT processes, causing a deterioration in low-temperature impact toughness. Therefore, the cooling end temperature is preferably 300°C or below.

[0129] On the other hand, if the cooling rate is less than 3°C / s, the base metal structure will be composed of a ferrite-bainite mixed structure or a bainite single-phase structure, rather than the martensite-bainite mixed structure or martensite single-phase structure required in the present invention. If the bainite structure is the main structure, strength and low-temperature impact toughness before PWHT can be ensured, but due to the formation of grain boundary cementite after PWHT as described above, it is difficult to ensure low-temperature impact toughness and CTOD quality. Furthermore, if the cooling rate exceeds 80°C / s, the microstructure will be composed of 100% martensite, and cross-sectional hardness will increase to a level that cannot be controlled by the tempering temperature alone. There is also a risk of plate fracture due to deviations in the cooling rate between the surface and center.

[0130] Therefore, the cooling rate during quenching is preferably 3°C / s or more and 80°C / s or less, and more preferably 5°C / s or more and 30°C / s or less.

[0131] Tempering can be performed on quenched steel sheets by heating them to a temperature of 550-700°C and holding for 5-60 minutes. If the tempering temperature is below 550°C or the holding time is less than 5 minutes, dislocation recovery does not occur properly during quenching, making the base material too strong and preventing sufficient low-temperature impact toughness. If the tempering temperature exceeds 700°C or the holding time exceeds 60 minutes, the grain boundaries (cementite) in the heat-affected zone may coarsen, making it difficult to ensure toughness after PWHT.

[0132] On the other hand, in addition to RQT, it can also be produced by a method in which direct quenching and tempering are performed after hot rolling (so-called DQT).

[0133] The manufactured steel material can be welded and subjected to post-weld heat treatment.

[0134] welding The steel material of the present invention can be used after being welded and subjected to post-weld heat treatment (PWHT). The steel material of the present invention can ensure the strength and low-temperature impact toughness of the base metal after post-weld heat treatment, and can also ensure the low-temperature impact toughness and low-temperature CTOD properties of the weld heat-affected zone (HAZ).

[0135] When steel is welded, the resulting structure is the unaffected base metal, the fused zone (weld metal) that melts and joins the weld, and the heat-affected zone (HAZ), which is the portion of the base metal affected by the welding heat. The fused zone and the heat-affected zone are collectively referred to as the "weld." The heat-affected zone is typically classified into coarse-grain HAZ (CGHAZ), subcritical HAZ (SCHAZ), intercritical HAZ (ICHAZ), and fine-grain HAZ (FGHAZ) based on the maximum heating temperature and microstructural characteristics. The CGHAZ is the region of the heat-affected structure (HAZ) that has the largest prior austenite grain size during welding. It is maintained at temperatures above 1000°C and then cooled. Larger grain sizes increase the ductile-brittle transition temperature and further increase hardenability. Therefore, if a hard structure is formed, brittle fracture resistance decreases. Therefore, the CGHAZ is the region of the HAZ where low-temperature impact toughness and toughness values ​​such as CTOD are most likely to deteriorate, making the properties of the CGHAZ crucial.

[0136] Welding can be performed by any method practicable in the technical field to which the present invention pertains, and is not necessarily limited thereto. In the present invention, the thermal cycle of welding can be performed by increasing the temperature to a maximum temperature of 1250 to 1400°C at a maximum heating rate of 50°C / s, holding at the target temperature for 5 to 10 seconds, performing primary cooling to 500°C at a cooling rate of 10 to 20°C / s, and then secondary cooling to 200°C at a cooling rate of 5 to 15°C / s. The thermal cycle related to the maximum temperature and cooling rate can be set based on the thermal history measured with a thermocouple at the CGHAZ position during actual welding.

[0137] The thermal history of welding was realized as similar as possible by measuring the thermal history of the heat affected zone (HAZ) during actual welding such as FCAW and SAW. As practical examples of the thermal history, the heat input for FCAW (Flux Cored Arc Weld) is 1.0-2.0 kJ / mm, and for SAW (Submerged Arc Weld) it is 2.5-4.0 kJ / mm. In this case, the preheating temperature and interpass temperature are both 150-250°C.

[0138] Post-weld heat treatment (PWHT) The post-weld heat treatment (PWHT) can be carried out by any method practicable in the technical field to which the present invention pertains, and is not particularly limited. For example, PWHT can be carried out in accordance with ASME Sec. VIII standard by holding the temperature at 595°C to 625°C for 180 to 540 minutes (min). Meanwhile, it is preferable that the combined LMP in the tempering heat treatment and PWHT heat treatment is 17.0 to 19.5 according to the following relations 2 and 3.

[0139] Specifically, when the tempering temperature and the PWHT temperature are the same, the LMP is determined by relational expression 2, and when the tempering temperature and the PWHT temperature are different, the LMP is determined by relational expression 3.

[0140] [Equation 2] LMP=T k (Log t+20)

[0141] where T k = Kelvin temperature of tempering and PWHT, t = tempering and PWHT time (hours)

[0142] [Equation 3] T(logt eq +20)=T i (logt i +20) Total LMP=T i [log(t i +t eq )+20]

[0143] where T i : PWHT temperature in Kelvin, t i : PWHT time (hours), T: Tempering temperature in Kelvin, t eq : Equivalent tempering time (hours) when converted to PWHT temperature

[0144] If the LMP is less than 17.0, the dislocation density will not be reduced as described above, which may result in excessive strength or a decrease in impact toughness due to excessive cross-sectional hardness. On the other hand, if the LMP is more than 19.5, the grain boundary occupancy ratio of cementite after PWHT will be too high, making it difficult to adequately ensure impact toughness and CTOD quality. Therefore, the LMP is preferably between 17.0 and 19.5, more preferably between 18.0 and 19.3. [Example]

[0145] Hereinafter, examples of the present invention will be described. It goes without saying that various modifications of the following examples are possible for those skilled in the art without departing from the scope of the present invention. The following examples are provided for the purpose of understanding the present invention, and the scope of the present invention should not be limited to the following examples, but should be defined by the claims below as well as equivalents thereof.

[0146] Example 1 Molten steel having the alloy elements shown in Table 1 (in Table 1, the content units of the component elements are wt.%, but the content units of P, S, and Ca are ppm. The remaining components are Fe and inevitable impurities) and Ceq of the above-mentioned Relational Formula 1 was used to produce steel slabs under the conditions shown in Table 2, which were then heated and hot-rolled to produce hot-rolled steel sheets with a thickness of 50 mm.

[0147] The produced hot-rolled steel sheets were reheated and quenched under the conditions disclosed in Table 3, and then tempered at 610°C for 20 minutes.

[0148] To confirm the excellent strength and low-temperature impact toughness of the steel material even after post-weld heat treatment, a heat treatment process simulating post-weld heat treatment was carried out. In this heat treatment process, the tempered hot-rolled steel sheet was subjected to heat treatment at 595°C for 180 minutes in three cycles.

[0149] For each test piece produced as described above, the microstructure and physical properties of the base material were measured, and the results are shown in Tables 4 and 5 below.

[0150] The prior austenite grain size and the fraction of each phase were measured using electron backscattered diffraction (EBSD) on specimens. Precipitates and inclusions were also measured using transmission electron microscopy (TEM). The precipitates and inclusions were identified using diffraction patterns and EDX mapping.

[0151] Yield and tensile strength were evaluated by room temperature tensile testing, with a 0.2% offset applied to yield strength. The impact toughness of each specimen was calculated by Charpy V-Notch testing, where the average absorbed energy was measured three times at each temperature. Brinell hardness was measured in accordance with ISO 6506 part 1.

[0152] [Table 1]

[0153] [Table 2]

[0154] [Table 3]

[0155] [Table 4]

[0156] [Table 5]

[0157] As can be seen from Tables 1 to 5 above, in the case of test specimen numbers 1 to 5 and 17 to 20, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, excellent quality can be ensured, including excellent base metal strength and excellent low-temperature impact toughness at -40°C, even after excellent post-weld heat treatment. In contrast, test specimen numbers 6 to 16 satisfy the alloy composition proposed by the present invention, but do not satisfy the manufacturing conditions, and it is understood that it is difficult to ensure the base metal properties proposed by the present invention. Furthermore, it is understood that test specimen numbers 21 to 25 satisfy the manufacturing conditions of the steel material proposed by the present invention, but do not satisfy the alloy composition, and it is difficult to ensure the base metal properties proposed by the present invention.

[0158] Example 2 The steel materials manufactured under the conditions in Table 3 above were subjected to welding and post-weld heat treatment, and the low-temperature toughness and CTOD properties of the CGHAZ in the weld heat-affected zone (HAZ) were confirmed.

[0159] Specifically, for this purpose, welding simulation was performed on the steel materials manufactured according to Tables 1 to 3 under the welding heat history conditions in Table 6, and a post-weld heat treatment process was performed at 595°C for 180 minutes in three cycles. In Table 6, the LMP was derived using the above-mentioned relational expression 3.

[0160] The microstructure and physical properties of the CGHAZ of each test piece prepared as described above were measured, and the results are shown in Tables 7 and 8 below.

[0161] Here, the phase fraction was measured by taking a test piece and using EBSD (Electron Back Scattered Diffraction).

[0162] The impact toughness of each specimen was calculated by Charpy V-Notch testing, where the average absorbed energy was measured three times at each temperature. Brinell hardness was measured according to ISO 6506 part 1, and CTOD tests were performed according to BS5762 (British Standard).

[0163] Additionally, the grain boundary occupancy ratio of cementite was measured using transmission electron microscopy (TEM). The grain boundary occupancy ratio of cementite was measured by measuring the proportion of film-like cementite in the total grain bondary length.

[0164] [Table 6]

[0165] [Table 7]

[0166] [Table 8]

[0167] As can be seen from Tables 6 to 8 above, in the case of test specimen numbers 1 to 5 and 11 to 16, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, not only excellent low-temperature impact toughness but also good CTOD quality can be ensured in the CGHAZ. In contrast, test specimen numbers 6 to 10 satisfy the alloy composition proposed by the present invention, but do not satisfy the conditions for manufacturing the steel material, and in particular, the size of the complex oxide inclusions is large, and it can be seen that the physical properties of the CGHAZ deteriorate. On the other hand, test specimen numbers 17 to 20 satisfy the alloy composition and the conditions for manufacturing the steel material proposed by the present invention, but do not satisfy the welding and post-weld heat treatment conditions, and the physical properties of the CGHAZ could not be ensured.

[0168] It can be seen that test pieces Nos. 21 to 25 satisfy the manufacturing conditions proposed by the present invention, but do not satisfy the alloy composition and do not satisfy the appropriate quality required by the present invention.

[0169] As described above, the detailed description of the present invention has been given with reference to the preferred embodiment of the present invention, but it goes without saying that a person skilled in the art to which the present invention pertains can make various modifications without departing from the scope of the present invention. Therefore, the scope of the present invention should not be limited to the described embodiment, but should be defined by the claims below as well as their equivalents.

Claims

1. The alloy contains, by 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 remainder being Fe and other inevitable impurities, Ceq defined by the following relational expression 1 satisfies the range of 0.50 to 0.70, The microstructure contains tempered martensite at an area fraction of 50% or more, The packet size of the microstructure includes 17 μm or less, CaO-Al observed in the upper layer, which is the region from the surface of the steel material to t / 4 (t is thickness, unit: mm) in the thickness direction. ２ O ３ -X (X is one or more of Ti, Mg, and S) complex oxidized inclusions with a maximum size of 50 μm or less and a unit area of ​​1 μm ２ A steel material characterized in that the number of hits is 50 or less. [Relationship 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15 In the above Relational Formula 1, [C], [Mn], [Cr], [Mo], [V], [Ni], and [Cu] respectively represent the contents (wt%) of C, Mn, Cr, Mo, V, Ni, and Cu contained in the steel, and when these components are not intentionally added, 0 is substituted.

2. 2. The steel material according to claim 1, wherein the microstructure includes bainite and other unavoidable structures.

3. 3. The steel material according to claim 1, wherein when the tempered martensite exceeds 80%, the Brinell hardness of the cross section of the steel material is 220 to 290 HB.

4. The steel material has a VC or VCN precipitate having a diameter of 5 to 15 nm and a size of 1 μm or less among precipitates observed in the cross section of the steel material. ２ 2. The steel material according to claim 1, wherein the number of particles per one particle is five or more.

5. 2. The steel material according to claim 1, wherein the microstructure is observed within prior austenite grains having an average grain size of 30 μm or less.

6. 2. 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 to 940 MPa, and a Charpy impact absorption energy at -40°C of 50 J or more.

7. In weight percent, 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, the remainder being Fe and other inevitable impurities, and Ceq defined by the following relational expression 1 satisfies the range of 0.50 to 0.70, and CaO-Al observed in the upper layer portion, which is the region from the surface of the steel slab to t / 4 (t is thickness, unit: mm) in the thickness direction ２ O ３ -X (X is one or more of Ti, Mg, and S) complex oxidized inclusions have a maximum size of 50 μm or less and a unit area of ​​1 μm ２ producing steel slabs having 50 or fewer pieces per slab; reheating the steel slab to a temperature range of 1050 to 1300°C; hot rolling the steel slab at a finish hot rolling temperature of 830 to 1050°C to produce a steel product; Heating the steel to a temperature range of 820 to 950°C, maintaining the temperature for 10 to 40 minutes, and then cooling the steel to a temperature of 300°C or less at a cooling rate of 3°C / s to 80°C / s based on a temperature standard of t / 4 of the thickness of the steel; After the cooling, the steel material is heated to 550 to 700° C. and held for 5 to 60 minutes for tempering. [Relationship 1] Ceq=[C]+[Mn] / 6+([Cr]+[Mo]+[V]) / 5+([Ni]+[Cu]) / 15 In the above Relational Formula 1, [C], [Mn], [Cr], [Mo], [V], [Ni], and [Cu] respectively represent the contents (wt%) of C, Mn, Cr, Mo, V, Ni, and Cu contained in the steel, and when these components are not intentionally added, 0 is substituted.

8. The step of producing the steel slab comprises: After the secondary refining, adding a metal Ca wire to the molten steel so that the Ca addition amount is 0.015 to 0.15 kg / ton; After adding the Ca, the method includes a step of bubbling the inert gas into the ladle for 5 to 40 minutes at a rate of 10 to 50 liters / min; 8. The method for manufacturing steel material according to claim 7, wherein the Ca wire is introduced at a rate of 100 to 300 meters / min, and the number of inert gas injection points in the ladle is two.

9. The method for manufacturing a steel material according to claim 7, further comprising the step of multi-stage stack cooling in a temperature range from 200° C. or more to room temperature after the hot rolling.

10. The method of claim 7, further comprising the step of subjecting the steel to post-weld heat treatment.

11. The method for manufacturing a steel material according to claim 10, wherein the post-weld heat treatment is carried out by holding the steel 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 tempering heat treatment and the post-weld heat treatment are performed at the same temperature, the LMP according to the following relational expression 2 is 17.0 to 19.

5. [Relationship 2] LMP=T ｋ (Log t+20) Here, T ｋ = Kelvin temperature of tempering and PWHT, t = tempering and PWHT time (hours)

13. The method for producing a steel material according to claim 11, wherein when the tempering heat treatment and the post-weld heat treatment are performed at different temperatures, the LMP according to the following relational expression 3 is 17.0 to 19.

5. [Relationship 3] T(logt) ｅｑ +20) = T ｉ (logt ｉ +20) Total LMP=T ｉ [log(t ｉ +t ｅｑ )+20] Here, T ｉ : Kelvin temperature of PWHT, t ｉ : PWHT time (hour), T: tempering temperature in Kelvin, t ｅｑ : Equivalent tempering time (hours) converted to PWHT temperature

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