Steel, welded heat-affected zone, and method of manufacture thereof

Abstract

The present invention relates to a steel for a liquefied gas storage tank, a weld heat-affected zone thereof, and a manufacturing method thereof.

Description

Technical Field

[0001] This disclosure relates to steel used in liquefied gas storage tanks, welded heat-affected zones, and methods for their manufacture. Background Technology

[0002] Recently, with the strengthening of environmental regulations, the demand for liquefied gases (LNGs), which are environmentally friendly fuels, is increasing. Consequently, the construction of LNG transport containers or propulsion vessels using LNGs as fuel is on the rise. In the case of liquefied gases, particularly liquefied natural gas (LNG) and liquefied ethylene / ethane, high-Ni steel is typically used due to their very low liquefaction temperatures. However, for gases such as LPG, ammonia, and CO2, which have liquefaction temperatures of approximately -60°C, and to reduce the construction costs of tanks and vessels used for gas storage, the use of carbon steel with improved strength or low-temperature toughness is becoming more prevalent.

[0003] Since conventional tanks used for transporting liquefied gases primarily employ large Type A tank designs without the application of high pressure, there is no need for high-strength or heavy-duty materials. However, recently, CO2 tanks or ammonia fuel tanks, which require high pressure during liquefaction, have been designed as Type C tanks, necessitating not only higher reinforcement of the steel but also greater thickness. Furthermore, since the post-weld heat treatment (PWHT) required by IGC specifications during the drying process of Type C tanks is essentially applied, the steel must also maintain the physical properties after PWHT.

[0004] PWHT is a process of holding the material at a high temperature for a certain period of time to eliminate residual stress after welding. Because the grain size and precipitate size of the low-temperature microstructure increase during PWHT, material quality, such as strength or toughness, deteriorates, making quality assurance difficult. In particular, in the case of steel used in CO2 tanks, classification societies have recently required the crack tip opening displacement (CTOD) quality in the post-weld heat-affected zone (HAZ) as a structural compatibility assessment item to improve the design stability of the tank.

[0005] Meanwhile, liquefied CO2 has different liquefaction temperatures depending on the pressure. Therefore, tanks can be mainly divided into low-pressure tanks and medium-pressure tanks.

[0006] Since low-pressure tanks, typically operating at pressures of 5 to 10 bar, need to be cooled to -60°C for CO2 liquefaction, steel with a tensile strength (TS) grade of 600 to 650 MPa, capable of guaranteeing cryogenic impact toughness at -60°C or lower, is used for the storage tanks. To maintain the cryogenic temperature after liquefaction, separate cooling maintenance equipment is required, which is a major reason for the increased initial installation and maintenance costs.

[0007] Medium-pressure tanks operating at 18 to 20 bar require higher liquefaction temperatures of -35°C, thus potentially facilitating the storage and transport of liquefied CO2. However, due to the high operating pressures, ultra-high-strength steel with a TS grade of 780 MPa or higher is used for medium-pressure tanks, replacing conventional steel with a TS grade of 600 to 650 MPa. Because steels containing high alloy components are manufactured through quenching-tempering (QT) heat treatment to ensure high strength, ensuring weld quality is not always straightforward.

[0008] Patent document 1 discloses that steel capable of guaranteeing low-temperature impact toughness at -40°C can be manufactured by heating a slab of an alloy composition to 1100°C to 1200°C, wherein the alloy composition comprises, by weight %: 0.06% to 0.12% C, 0.02% to 0.06% Si, 0.7% to 1.2% Mn, 0.006% to 0.012% Ti, 0.002% to 0.01% Al, and 0.3% to 0.5% [unclear - possibly referring to a specific component or component]. Cr, 0.3% to 0.4% Mo, 0.03% to 0.04% V, 0.002% to 0.003% N, 0.002% to 0.01% S, 0.007% or less P, and the remainder Fe and impurities; rolled at 800°C to 900°C with a cumulative reduction of 30% or greater; air-cooled; re-quenched at 910°C to 930°C; and tempered at 650°C to 680°C.

[0009] Although methods have been provided to improve toughness by refining the HAZ grains through uniformly distributed Al-Ti-MnS composite oxide inclusions in the weld heat-affected zone, the steel cannot guarantee the material quality after PWHT and has a tensile strength of only 550 MPa to 650 MPa in the base material, making it unsuitable as a steel for medium-pressure liquefied CO2 (LCO2).

[0010] Meanwhile, Patent Document 2 discloses a method for manufacturing P690QL2 steel for LCO2 storage tanks, wherein the steel comprises, by weight percent: 0.09% to 0.12% C, 0.9% to 1.3% Mn, 0.015% or less Ti, 0.002% or less S, 0.06% or less Nb, 0.08% or less V, 1.5% to 2.5% Ni, 0.2% to 0.4% Cr, 0.3% to 0.4% Mo, 0.3% or less Cu, 0.03% to 0.05% Al, 0.0005% to 0.005% Ca, 0.005% or less N, 0.0003% or less H, and the remainder being Fe and impurities. The steel has a thickness of 50 mm or less and can guarantee low-temperature impact toughness. Patent document 2 discloses that when a 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 tempered at 630°C to 680°C, a tensile strength of 780 MPa or greater and an impact toughness of 47 J or greater at -60°C can be ensured after PWHT.

[0011] However, since the CTOD guarantee temperature and standards of the HAZ section do not meet the levels required in this disclosure, no practical measures are proposed for obtaining steel for use in medium-pressure LCO2 storage tanks.

[0012] Patent document 3 discloses a method for manufacturing steel plates comprising, by weight percent: 0.03% to 0.18% C, 0.30% or less Si, 0.3% to 1.6% Mn, 0.015% or less P, 0.005% or less S, 2.0% or less Cu, 1.0% to 7.0% Ni, 0.01% to 0.20% Al, 0.005% or less Ca, 0.0100% or less N, 0.0060% or less O, and the remainder being Fe and impurities. The steel plate is characterized by having a Mn / Ni ratio of 0.80 or less (the ratio of Mn to Ni content), a WES-based carbon equivalent (Ceq) of 0.430% to 0.900%, a tensile strength of 780 MPa or greater and 930 MPa or less, and a tensile strength of 630 MPa or greater and 750 MPa or less. A yield strength of MPa or less, a yield ratio of 85% or less, HAZ partial impact toughness of 100 J or more at 0°C, a plate thickness of 40 mm or more and 120 mm or less, and an Hvmin / Hvmax ratio of 0.85 or less, where Hvmin is the average of the higher 20% values ​​of the Vickers hardness on the lower side and Hvmax is the average of the lower 20% values ​​of the Vickers hardness on the higher side when measuring Vickers hardness at 225 or more points at 1 / 4 of the plate thickness on a plate surface of 40 mm or more and 120 mm or less. While the strength meets the required range for steels used in medium-pressure LCO2 applications, impact toughness cannot be guaranteed at approximately -40°C (where CO2 liquefies), therefore, this steel may not be suitable. Furthermore, when Ceq and hardness are too high, the martensite fraction increases excessively during quenching, and it may be difficult to ensure impact toughness and HAZ partial CTOD properties.

[0013] Therefore, in order to ensure an appropriate low-temperature phase fraction, a range of Ceq and its corresponding surface hardness values ​​are required. However, since this is not described in the literature, it is not suitable as a technique for manufacturing ultra-high-strength steel for medium-pressure LCO2.

[0014] Therefore, the aforementioned conventional techniques do not offer a practical solution for manufacturing ultra-high strength steel for medium-pressure LCO2, which has a yield strength of 690 MPa or greater after post-weld heat treatment (PWHT), a tensile strength of 770 MPa to 940 MPa, low-temperature impact toughness at -40°C or lower, and a CTOD quality of 0.1 mm or greater.

[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 this disclosure is to provide a weld heat-affected zone with excellent low-temperature CTODD quality, as well as excellent strength and low-temperature impact toughness after post-weld heat treatment (PWHT), and a method thereof for manufacturing the same.

[0020] Another aspect of this disclosure is to provide steel with excellent strength and low-temperature impact toughness after post-weld heat treatment (PWHT) and a method for manufacturing the same.

[0021] The purpose of this disclosure is not limited to the foregoing description. Other purposes of this disclosure are described throughout the specification and will be readily apparent to any person skilled in the art from the description herein.

[0022] Technical solution

[0023] According to one aspect of this disclosure, the weld heat-affected zone comprises, by weight percent: 0.05% to 0.20% C, 0.05% to 0.50% Si, 0.5% to 2.0% Mn, 0.005% to 0.10% Al, 0.010% or less P, 0.0050% or less S, 0.001% to 0.070% Nb, and 0.001% to 0.30% V. 0.001% to 0.030% Ti, 0.01% to 1.0% Cr, 0.01% to 1.0% Mo, 0.01% to 0.60% Cu, 1.0% to 4.0% Ni, 0.001% to 0.005% B, 0.0005% to 0.0040% Ca, 0.002% or less O, and the remainder Fe and other unavoidable impurities.

[0024] Ceq, defined by relation 1, satisfies a value in the range of 0.50 to 0.70.

[0025] The largest size of the CaO-Al2O3-X (X is one or more of Ti, Mg, and S) composite oxide inclusions observed in the upper layer is 50 μm or smaller, in the region t / 4 of the surface (where t is the thickness in mm) from the surface layer to the thickness direction, and per 1 μm2 The number of complex oxide inclusions per unit area is 50 or fewer.

[0026] The cementite present in the grain boundaries has a grain boundary occupancy of 10% or less, and the average size of the cementite is 1 μm or less.

[0027] The cross-sectional hardness is 200 HB to 300 HB.

[0028] [Relation 1]

[0029]

[0030] Where [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] are the contents (by weight%) of C, Mn, Cr, Mo, V, Ni and Cu in the steel, respectively, and 0 is used when the component is not intentionally included.

[0031] The weld heat-affected zone can be a coarse-grained heat-affected zone (CGHAZ).

[0032] The microstructure of the weld heat-affected zone may contain 50% or more tempered martensite by area fraction, with the remainder being bainite.

[0033] The average grain size of the original austenite in the weld heat-affected zone can be 200 μm or smaller.

[0034] The weld heat-affected zone can have a Charpy impact absorption energy of 40 J or greater at -40°C and a CTOD characteristic of 0.1 mm or greater at -35°C.

[0035] According to another aspect of this disclosure, a method for manufacturing a weld heat-affected zone includes: preparing steel in which the composition and Ceq are satisfied.

[0036] The largest size of the CaO-Al2O3-X (X is one or more of Ti, Mg, and S) composite oxide inclusions observed in the upper layer is 50 μm or smaller, in the region t / 4 of the surface (where t is the thickness in mm) from the surface layer to the thickness direction, and per 1 μm 2 The number of inclusions per unit area is 50 or fewer;

[0037] Welding of steel; and

[0038] The welded steel is subjected to post-weld heat treatment.

[0039] The welding heat regime can be as follows: heating to a maximum temperature of 1250°C to 1400°C at a heating rate of up to 50°C / second, holding at the target temperature for 5 to 10 seconds, first cooling to 500°C at a cooling rate of 10°C / second to 20°C / second, and then second cooling to 200°C at a cooling rate of 5°C / second to 15°C / second.

[0040] The welding can be FCAW welding with a heat input of 1.0 kJ / mm to 2.0 kJ / mm under welding conditions with a preheating temperature and interpass temperature of 150°C to 250°C, or SAW welding with a heat input of 2.0 kJ / mm to 4.0 kJ / mm under welding conditions with a preheating temperature and interpass temperature of 150°C to 250°C.

[0041] Post-weld heat treatment can be performed by holding at 595°C to 625°C for 180 to 540 minutes.

[0042] The steel may contain 50% or more (including 100%) of tempered martensite by area fraction, as well as bainite and other unavoidable microstructures in the remainder, and the lath group size of the tempered martensite and bainite may be 17 μm or smaller.

[0043] When the tempered martensite content is greater than 80%, the steel can have a Brinell hardness of 220 HB to 290 HB.

[0044] Among the precipitates observed in the cross-section of the steel, the steel can be 1 μm in size. 2 It has 5 or more VC or VCN precipitates with a diameter of 5 nm to 15 nm.

[0045] The average grain size of the original austenite in steel can be 30 μm or smaller.

[0046] Steel can be manufactured by methods including the following:

[0047] Manufacturing steel billets;

[0048] The steel billet is reheated to a temperature in the range of 1050°C to 1300°C;

[0049] Steel billets are hot-rolled at a fine hot-rolling temperature of 830°C to 1050°C to manufacture steel products;

[0050] The steel is heated to a temperature ranging from 820°C to 950°C, held at that temperature for 10 to 40 minutes, and then cooled to 300°C or lower at a cooling rate of 3°C / sec to 80°C / sec, based on the temperature at t / 4 of the steel thickness; and

[0051] After cooling, the steel is subjected to tempering heat treatment, which involves heating it to 550°C to 700°C and holding it at that temperature for 5 to 60 minutes.

[0052] The manufacture of steel billets may include:

[0053] After the second refining, metallic Ca wire is added to the molten steel, such that the amount of Ca added is between 0.015 kg / ton and 0.15 kg / ton; and

[0054] After adding Ca, perform cleaning bubbling for 5 to 40 minutes to ensure that the inert gas flow rate in the ladle is 10 to 50 liters per minute.

[0055] The Ca wire addition rate is 100 m / min to 300 m / min, and the number of inert gas blowing points in the ladle is two.

[0056] When the tempering heat treatment and post-weld heat treatment temperatures are the same, the LMP, according to the following relationship 2, can be between 17.0 and 19.5:

[0057] [Relationship 2]

[0058]

[0059] Where T k =Tempering and PWHT temperatures in Kelvin, and t =Tempering and PWHT time (hours).

[0060] When the tempering heat treatment and post-weld heat treatment temperatures are different, the LMP can be between 17.0 and 19.5 according to the following relationship 3:

[0061] [Relationship 3]

[0062]

[0063]

[0064] Where T i The temperature of PWHT in Kelvin, t i PWHT time (hours), T: tempering temperature in Kelvin, and t eq Equivalent to tempering time (in hours) converted to PWHT temperature.

[0065] According to another aspect of this disclosure, regarding steels that satisfy the above composition and Ceq,

[0066] It contains 50% or more tempered martensite by area as its microstructure.

[0067] The lath groups of the microstructure have a size of 17 μm or smaller, and

[0068] The maximum size of the CaO-Al2O3-X (X is one or more of Ti, Mg, and S) composite oxide inclusions observed in the upper layer is 50 μm or less, and this upper layer is located in the region t / 4 (where t is the thickness in mm) from the steel surface in the thickness direction from the surface layer, and every 1 μm 2 The number of complex oxide inclusions per unit area is 50 or fewer:

[0069] [Relation 1]

[0070]

[0071] Where [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] are the contents (by weight%) of C, Mn, Cr, Mo, V, Ni and Cu in the steel, respectively, and 0 is used when the component is not intentionally included.

[0072] According to another aspect of this disclosure, a method for manufacturing steel includes: manufacturing a steel billet in which the composition and Ceq are satisfied, and the maximum size of CaO-Al2O3-X (X is one or more of Ti, Mg and S) composite oxide inclusions observed in the upper portion is 50 μm or less, said upper portion being located in a region t / 4 (where t is the thickness in mm) from the surface of the billet in the thickness direction, and every 1 μm 2 The number of complex oxide inclusions per unit area is 50 or less;

[0073] The steel billet is reheated to a temperature in the range of 1050°C to 1300°C;

[0074] Steel billets are hot-rolled at a fine hot-rolling temperature of 830°C to 1050°C to manufacture steel products;

[0075] The steel is heated to a temperature ranging from 820°C to 950°C, held at that temperature for 10 to 40 minutes, and then cooled to 300°C or lower at a cooling rate of 3°C / sec to 80°C / sec, based on the temperature at t / 4 of the steel thickness; and

[0076] After cooling, the steel is subjected to tempering heat treatment, which involves heating it to 550°C to 700°C and holding it at that temperature for 5 to 60 minutes.

[0077] Beneficial effects

[0078] According to one aspect of this disclosure, a weld heat-affected zone having low-temperature impact toughness and low-temperature CTOD properties after post-weld heat treatment (PWHT) and a method thereof can be provided.

[0079] According to another aspect of this disclosure, there is a method for manufacturing steel that exhibits excellent strength and low-temperature impact toughness even after post-weld heat treatment (PWHT) by minimizing the deterioration of physical properties.

[0080] The various beneficial advantages and effects of this disclosure are not limited to those described above, and can be more readily understood in the process of describing specific example implementations in this disclosure. Detailed Implementation

[0081] The terminology used in this specification is for describing the content of this disclosure and not for limiting it. Furthermore, unless the applicable limitation has the meaning expressly contrary to it, the singular form used in this specification includes the plural form as well.

[0082] The term "includes / contains" as used in this specification means to include the configuration and does not exclude the existence or addition of other configurations.

[0083] Unless otherwise defined, all terms used in this specification, including technical and scientific terms, shall have the same meaning as commonly understood by one of ordinary skill in the art to which this disclosure pertains. Predefined terms are to be interpreted as consistent with relevant technical literature and the disclosure herein.

[0084] Preferred exemplary embodiments of this disclosure will be described below. These exemplary embodiments may be modified in various ways, and the scope of this disclosure should not be construed as limited to the exemplary embodiments set forth below. These exemplary embodiments are provided to describe this disclosure in more detail to those skilled in the art to which this disclosure pertains.

[0085] First, the composition of steel will be described in detail as an example of the content of this disclosure.

[0086] The steel contains, by weight percent: 0.05% to 0.20% C, 0.05% to 0.50% Si, 0.5% to 2.0% Mn, 0.005% to 0.10% Al, 0.010% or less P, 0.0050% or less S, 0.001% to 0.070% Nb, 0.001% to 0.30% V, 0.001% to 0.030% Ti, 0.01% to 1.0% Cr, 0.01% to 1.0% Mo, 0.01% to 0.60% Cu, 1.0% to 4.0% Ni, 0.001% to 0.005% B, 0.0005% to 0.0040% Ca, 0.002% or less O, and the remainder Fe and other unavoidable impurities.

[0087] The Ceq value, defined by relation 1, can be between 0.50 and 0.70:

[0088] [Relation 1]

[0089]

[0090] Where [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] are the contents (by weight%) of C, Mn, Cr, Mo, V, Ni and Cu in the steel, respectively, and 0 is used when the component is not intentionally included.

[0091] The composition of the steel can be the same as that of the steel billet, the steel before welding, the steel after welding, the base material of the steel that has undergone post-weld heat treatment, and the composition of the weld heat-affected zone (HAZ).

[0092] The composition will be described in detail below. Unless otherwise specifically indicated, all percentages and ppm relating to the alloy composition described below are by weight.

[0093] Carbon (C): 0.05% to 0.20%

[0094] Since carbon (C) is the most important element for ensuring basic strength, it needs to be included in steel within an appropriate range, and to achieve the desired additive effect, 0.05% or more of carbon (C) should be added. Preferably, 0.10% or more of carbon (C) may be included. However, when the carbon (C) content exceeds a certain level, the martensite fraction increases excessively, the hardness of the weld heat-affected zone (HAZ) increases, and the fraction of grain boundary cementite increases, resulting in reduced impact propagation resistance in impact tests or CTOD tests, and potentially failing to ensure the toughness value required in this disclosure. Therefore, the carbon (C) content in this disclosure can be 0.20%, and more preferably, the upper limit of the carbon (C) content can be 0.18%.

[0095] Silicon (Si): 0.05% to 0.50%

[0096] Silicon (Si), as a substitutional element, improves the strength of steel through solid solution strengthening and has a strong deoxidizing effect, making it an essential element for manufacturing clean steel. Therefore, silicon (Si) can be included in the composition at 0.05% or more, preferably 0.20% or more. However, when silicon (Si) is included in large quantities, a martensite-austenite (MA) phase is formed, and the matrix strength of the martensite phase may be excessively increased, thereby deteriorating low-temperature impact toughness. Therefore, the upper limit of its content can be 0.50%. The upper limit of silicon (Si) content can more preferably be 0.40%.

[0097] Manganese (Mn): 0.5% to 2.0%

[0098] Manganese (Mn) is an element that improves strength and hardenability through solid solution strengthening to produce a low-temperature transformation phase. Therefore, to ensure a yield strength of 690 MPa or greater, it is preferable to contain 0.5% or more manganese (Mn). The manganese (Mn) content can preferably be 0.7% or greater. However, since manganese (Mn) can form MnS (which is an elongated non-metallic inclusion) with sulfur (S), thereby deteriorating toughness and acting as an impact initiation point, it may be a cause of rapid deterioration of the low-temperature impact toughness of steel. Therefore, it is preferable to control the manganese (Mn) content to 2.0% or less, and the manganese (Mn) content can preferably be 1.5% or less.

[0099] Aluminum (Al): 0.005% to 0.10%

[0100] Aluminum (Al) and silicon (Si) are both powerful deoxidizers in steelmaking processes, and to achieve this effect, it is preferable to include 0.005% or more aluminum. The lower limit of the aluminum (Al) content can be more preferably 0.01%. However, when the aluminum (Al) content is excessive, the fraction of Al2O3 in the oxide inclusions generated as a result of deoxidation increases excessively, leading to coarsening and making it difficult to remove the inclusions during refining. Therefore, this may be a cause of deterioration in low-temperature impact toughness. Therefore, it is preferable to control the aluminum (Al) content to 0.10% or less. The aluminum (Al) content can be more preferably 0.07% or less.

[0101] Phosphorus (P): 0.010% or less (inclusive) and Sulfur (S): 0.0050% or less (inclusive)

[0102] Phosphorus (P) and sulfur (S) are elements that cause brittleness at grain boundaries or lead to the formation of coarse inclusions. Therefore, it is preferable that phosphorus (P) and sulfur (S) are not included. To improve resistance to brittle crack propagation, it is preferable to include 0.010% or less of phosphorus (P) and 0.0050% or less of sulfur (S).

[0103] Niobium (Nb): 0.001% to 0.070%

[0104] Niobium (Nb) is an element that precipitates in the form of NbC or NbCN and improves the strength of the substrate. Furthermore, since niobium (Nb) that achieves a solid solution state during reheating at high temperatures precipitates very finely in the form of NbC during rolling to suppress austenite recrystallization, it has the effect of refining the microstructure. Therefore, it is preferable to contain 0.001% or more of niobium (Nb), and the content of niobium (Nb) can be more preferably 0.005% or more. However, when niobium (Nb) is contained in excess, undissolved niobium (Nb) is produced in the form of TiNb(C,N), which deteriorates the low-temperature impact toughness; therefore, it is preferable that the upper limit of the niobium (Nb) content is 0.070%. The upper limit of the niobium (Nb) content can be more preferably 0.065% or less.

[0105] Vanadium (V): 0.001% to 0.30%

[0106] Since vanadium (V) is almost entirely re-solidified during reheating, its strengthening effect through precipitation or solution during subsequent rolling is not significant. However, it precipitates as very fine carbonitrides during tempering heat treatment or PWHT to improve strength. To achieve this effect, vanadium (V) needs to be added at 0.001% or more. The lower limit of vanadium (V) content can be more preferably 0.01%. However, when the content is excessive, vanadium excessively increases the surface layer hardness of the slab due to its high hardenability, thus acting as a factor in surface cracking during flange processing and rapidly increasing manufacturing costs, which is detrimental to commercial purposes. Therefore, the vanadium (V) content can be limited to 0.3% or less. The vanadium (V) content can be more preferably 0.25% or less.

[0107] Titanium (Ti): 0.001% to 0.030%

[0108] Titanium (Ti) is a component that precipitates as TiN during reheating to suppress the growth of the original austenite grains at high temperatures and significantly improve low-temperature toughness. To achieve this effect, it is preferable to include 0.001% or more of titanium (Ti). However, when titanium (Ti) is excessively included, low-temperature toughness may decrease due to clogging of the soft casting nozzle or crystallization in the central portion. Furthermore, since titanium (Ti) combines with nitrogen (N) to form coarse TiN precipitates in the thickness center portion, the elongation of the product deteriorates, which may lead to a loss of uniform elongation during stamping, resulting in surface cracking. Therefore, the titanium (Ti) content can be 0.030% or less. The upper limit of the titanium (Ti) content is preferably 0.025% or less, and more preferably 0.018% or less.

[0109] Chromium (Cr): 0.01% to 1.0%

[0110] Chromium (Cr) is a component that increases hardenability to form a low-temperature transformation structure, thereby increasing yield strength and tensile strength. Furthermore, it slows down the spheroidization rate of cementite, thus preventing a decrease in strength. For these effects, 0.01% or more of chromium (Cr) may be included. However, when the chromium (Cr) content is excessive, coarse carbides rich in Cr, such as M... 23 The increased size and fraction of C6 degrades the impact toughness of the steel, reduces the solid solubility of niobium (Nb) and the fraction of fine precipitates such as NbC, and thus the reduced product strength may be problematic. Therefore, the upper limit of chromium (Cr) content in this disclosure can be 1.0%. The upper limit of chromium (Cr) content is preferably 0.8%.

[0111] Molybdenum (Mo): 0.01% to 1.0%

[0112] Molybdenum (Mo) is an element that increases grain boundary strength and has a significant solid solution strengthening effect in ferrite, effectively contributing to increased strength and ductility of products. Furthermore, Mo has the effect of preventing the reduction in toughness caused by grain boundary segregation of impurity elements such as phosphorus (P). For these effects, 0.01% or more of Mo can be added. However, Mo is an expensive element, and excessive addition can significantly increase manufacturing costs; therefore, a maximum Mo content of 1.0% is preferred.

[0113] Copper (Cu): 0.01% to 0.60%

[0114] Copper (Cu) is an advantageous element in this disclosure because it can significantly improve the strength of the matrix phase through solid solution strengthening in ferrite and also has the effect of inhibiting corrosion in a wet hydrogen sulfide atmosphere. For these effects, 0.01% or more of copper (Cu) may be included. The copper (Cu) content may more preferably be 0.03% or more. However, when the copper (Cu) content is excessive, copper (Cu) is more likely to cause star-shaped cracks on the surface of the steel sheet, and as an expensive element, it may lead to a problem of significantly increased manufacturing costs. Therefore, the upper limit of the copper (Cu) content is preferably 0.60%, and more preferably 0.35%.

[0115] Nickel (Ni): 1.0% to 4.0%

[0116] Nickel (Ni) is an element that effectively contributes to the increase of stacking defects at low temperatures to promote cross-slip of dislocations, thereby improving impact toughness and hardenability to improve strength. For these effects, 1.0% or more of nickel (Ni) may be included. The nickel (Ni) content is preferably 0.10% or greater. However, when excessive nickel (Ni) is added, manufacturing costs may increase due to high costs; therefore, the upper limit of the nickel (Ni) content is preferably 4.0%, and more preferably 3.0%.

[0117] Calcium (Ca): 0.0005% to 0.0040%

[0118] When calcium (Ca) is added after deoxidation with aluminum (Al), the calcium combines with sulfur (S) to form MnS inclusions, thus inhibiting MnS formation. Simultaneously, spherical CaS is also formed, thereby suppressing crack formation due to hydrogen-induced cracking. To ensure sufficient formation of CaS from the sulfur (S) contained as an impurity, it is preferable to include 0.0005% or more of calcium (Ca). However, when the amount added is excessive, the remaining calcium (Ca) after CaS formation combines with oxygen (O) to form coarse oxide inclusions. These coarse oxide inclusions are stretched and fractured during rolling, contributing to the deterioration of lamellar tear resistance. Therefore, the upper limit of the calcium (Ca) content is preferably 0.0040%.

[0119] Boron (B): 0.001% to 0.005%

[0120] Boron (B) is an element used to improve the hardenability of hot-rolled steel sheets by delaying the transformation from austenite to ferrite during cooling. To achieve this effect, 0.001% or more of boron is preferably added; however, when the content exceeds 0.005%, the strength increases excessively, which deteriorates the impact toughness. Furthermore, boron oxide is formed, thereby deteriorating the surface quality of the steel sheet.

[0121] Oxygen (O): 0.002% or less

[0122] O combines with Ca or Al in the molten steel and exists as complex oxide inclusions of CaO, Al2O3, or CaO-Al2O3. Oxide inclusions generated at high temperatures are broken up during the slab roughing step and stretched in the rolling direction, remaining in the HAZ without dissolving during welding and during the rolling / heat treatment of the substrate. Since the ends of the broken oxide inclusions may act as fracture initiation points when assessing impact and CTOD due to the notch effect, it is preferable to remove them as much as possible through flotation separation in the second refining process; therefore, the oxygen concentration in the molten steel is preferably 20 ppm or less. In this disclosure, a second refining method is proposed that allows for extremely precise control of oxide inclusions, and after applying this method, the total oxygen content can more preferably be 15 ppm or less, and most preferably 10 ppm or less.

[0123] In addition to the components mentioned above, Fe and other unavoidable impurities may also be present. However, since unintended impurities may inevitably be introduced from raw materials or the surrounding environment during ordinary manufacturing processes, these may not be completely eliminated. Since these impurities are known to any person skilled in the art, not all of them are specifically mentioned in this specification.

[0124] Furthermore, the possibility of adding other effective components besides those mentioned above is not entirely excluded.

[0125] The Ceq constituting the following relation 1 can be between 0.50 and 0.70:

[0126] [Relation 1]

[0127]

[0128] Where [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] are the contents (by weight%) of C, Mn, Cr, Mo, V, Ni and Cu in the steel, respectively, and 0 is used when the component is not intentionally included.

[0129] When Ceq in Equation 1 is less than 0.50, the high yield strength required in this disclosure may not be ensured due to the reduced martensite fraction. Furthermore, the excessive increase in bainite fraction leads to an increased proportion of grain boundary cementite in the weld heat-affected zone (HAZ) after post-weld heat treatment (PWHT), particularly in the coarse-grained HAZ (CGHAZ), which may further compromise the appropriate impact toughness and HAZ CTOD value required in this disclosure. Additionally, when the Ceq value is greater than 0.70, the excessive increase in hardness can lead to cleavage fracture, thus reducing resistance to crack propagation and making it difficult to ensure appropriate impact toughness and CTOD properties in the weld heat-affected zone.

[0130] As an example of this disclosure, the steel may have a maximum size of 50 μm or less of CaO-Al2O3-X (X being one or more of Ti, Mg, and S) complex oxide inclusions observed in the upper portion, and 50 or fewer inclusions per 1 μm. 2 The number of composite oxide inclusions per unit area, wherein the upper portion is within t / 4 (t: thickness) from the steel surface in the direction from the surface layer to the thickness.

[0131] Complex oxide inclusions can be observed in steel before welding, steel after welding, steel substrates undergoing post-weld heat treatment, the weld heat-affected zone (HAZ), and steel billets.

[0132] When complex oxide inclusions are formed during the billet manufacturing process, they can then be broken up during rough rolling without changing their form and fraction during the product manufacturing process, and stretched in the rolling direction. Since broken oxide inclusions can act as stress concentration points and as initiation points for fracture, the amount or size of oxide inclusions generated during billet manufacturing increases after rough rolling, becoming a major cause of deterioration in impact toughness and CTOD quality. Therefore, when the size of complex oxide inclusions in the upper part of the product (ranging from the surface layer to t / 4) is greater than 50 μm or per 1 μm... 2 When the quantity is greater than 50, it may be difficult to ensure the low-temperature impact toughness at -40°C based on the substrate and the HAZ CTOD quality at -35°C or lower.

[0133] The steel disclosed herein has a matrix structure of tempered martensite or a mixture of tempered martensite and bainite. Specifically, the steel comprises 50% or more tempered martensite by area fraction, with the remainder consisting of bainite and unavoidable microstructures.

[0134] When the matrix structure is a mixture of tempered martensite and bainite, the area fraction of tempered martensite can be 50% to 80%.

[0135] When the matrix structure is formed by a single structure of more than 80% tempered martensite, the hardness of the base section of the steel after PWHT can be 220 HB to 290 HB, based on Brinell hardness.

[0136] The structure that inevitably occurs can be martensitic-austenitic (MA).

[0137] The lath group size of the low-temperature structure of tempered martensite or bainite can be 17 μm or smaller.

[0138] The average grain size of the original austenite in steel can be 30 μm or smaller.

[0139] Fine VC or VCN precipitates with diameters ranging from 5 nm to 15 nm were observed in the cross-section of the steel, with a density of approximately 1 μm. 2 It includes five or more precipitates.

[0140] The microstructure characteristics can be the same in the base material of steel that has undergone post-weld heat treatment, as well as in the steel before and after welding.

[0141] The average grain size of the original austenite in the steel can be 30 μm or less. Post-weld heat treatment (PWHT) is typically performed at a high temperature of 595°C to prevent weld cracking through stress relief after welding. Here, dislocation density and cross-sectional hardness can be reduced, thus increasing weld integrity. However, during prolonged holding at high temperatures, cementite generated at the bainite lath interface decomposes, leading to long-range carbon diffusion to the austenite grain boundaries. Therefore, as the grain size increases, the area of ​​grain boundaries occupied by cementite decreases, potentially increasing segregation and deteriorating impact toughness. Therefore, it is necessary to control the original austenite grain size through appropriate low-temperature high-pressure treatment. Thus, the original austenite grain size is preferably 30 μm or less, more preferably 25 μm or less.

[0142] As the temperature and time of PWHT (Pulsed Wood Hardening) increase, coarse cementite spheroidizes within the original austenite grain boundaries, with some forming as film-like cementite. When cementite forms within the grain boundaries, the grain boundary strength becomes lower than the shear strength, becoming a major cause of intergranular fracture. Therefore, to prevent intergranular fracture after PWHT, martensite, which only causes short-range diffusion, should be included in an appropriate fraction or more. However, when the martensite ratio is too high, the cross-sectional hardness increases excessively, resulting in a decrease in impact absorption energy even in the event of intergranular fracture, which degrades toughness. Therefore, it is preferable that the microstructure of the steel consists of a mixture of tempered martensite and bainite, wherein the fraction of tempered martensite is preferably 50% to 80% by area.

[0143] However, when the matrix structure of the steel microstructure is tempered martensite, i.e., martensite is greater than 80% (including 100%), the tempering temperature should be further increased to maintain the cross-sectional hardness of the substrate to 200 HB to 290 HB. When the cross-sectional hardness is outside this range, it is difficult to ensure the low-temperature impact toughness and CTOD value required in this disclosure.

[0144] The lath group size of a low-temperature structure of tempered martensite or bainite can be 17 μm or smaller. The lath group size can be understood as a domain boundary that may possess resistance to physical crack propagation (crack bending characteristics at the boundary during crack propagation), serving as a cluster of laths with the same orientation relationship in a low-temperature transformation structure such as bainite or martensite. As an example of measuring the lath group size, it can be measured as the grain boundary plane that can be determined by setting domains within a 15° orientation relationship to the same orientation using an electron backscattered diffraction (EBSD) instrument. Since bainite or martensite grows only within the original austenite grains, it has a smaller grain size than the original austenite. When the lath group size is greater than 17 μm, it is difficult to ensure sufficient impact toughness due to the difficulty in ensuring sufficient crack propagation resistance; therefore, 17 μm or smaller is preferred.

[0145] Simultaneously, the dislocation density of steel decreases during tempering and PWHT, resulting in a reduction in strength. To appropriately compensate for this, the use of fine precipitates can reduce strength degradation during heat treatment. In the case of normalized materials or materials processed by thermo-mechanical control processes (TMCP), precipitates such as NbC and NbCN are commonly used; however, these precipitates are generated at temperatures between 800°C and 900°C and are transformation-induced, making them unsuitable for quench-temper (QT) heat-treated steels. Therefore, vanadium carbides / carbonitrides such as VN and VCN, generated at temperatures in the range of 550°C to 700°C, are readily used. When the diameter of fine precipitates of VC or VCN is less than 5 nm, their size is too small, making the strength improvement effect potentially insignificant; while when the size is greater than 15 nm, excessive coarsening occurs, thus reducing the strength improvement effect. Therefore, the appropriate size of the fine precipitates is 5 nm to 15 nm. Meanwhile, when the density of VC and VCN is per 1 μm 2When the number of precipitates per unit area is less than five, the aforementioned precipitation strengthening effect also decreases; therefore, a number of five or more is preferred. The formation of precipitates depends on the chemical potential and temperature determined by the steel composition. Although precipitates with a size of 5 nm to 15 nm may appear in large quantities, as the number increases, the precipitates aggregate to increase their size; therefore, precipitates with a size of 5 nm to 15 nm will not form indefinitely. Therefore, although those skilled in the art are unaware of the upper limit to the number of fine precipitates, they are not ignorant of it.

[0146] Furthermore, the microstructure of a weld heat-affected zone (HAZ) of an example of this disclosure will be described in detail. The HAZ can be a HAZ following post-weld heat treatment (PWHT), and in particular, it can be a HAZ of a coarse-grained heat-affected zone (CGHAZ).

[0147] The grain boundary occupancy of cementite present in the weld heat-affected zone (HAZ) can be 10% or less.

[0148] Grain boundary occupancy can be understood as the ratio of the length occupied by the film-like cementite within the grain boundary, and in three dimensions as the area ratio. As an example of occupancy measurement, the sample is etched and observed using SEM or similar methods to measure ordinary grain boundaries and grain boundaries that produce film-like cementite for distinction.

[0149] Meanwhile, the average size of the cementite can be 1 μm or smaller.

[0150] The microstructure of the HAZ, particularly CGHAZ, following post-weld heat treatment (PWHT) preferably has 50% or more tempered martensite by area fraction. Besides tempered martensite, the remainder can be bainite and other microstructures. When the area fraction of tempered martensite is less than 50%, the occupancy of film-like cementite increases in the grain boundaries after PWHT, and the grains coarsen, making it difficult to ensure a suitable CTOD value of 0.1 mm or larger at -35°C or lower. For the same reason, the grain boundary occupancy of cementite present in the PWHT grain boundaries is preferably 10% or less, and the average size of the cementite present in the grain boundaries is preferably 1 μm or less.

[0151] The average grain size of the original austenite in the weld heat-affected zone (HAZ) can be 200 μm or smaller. The HAZ microstructure is determined by the microstructure of the substrate and the welding conditions (e.g., weld heat input / preheating temperature / delamination temperature). In particular, since the CGHAZ is a cross-section in which the grains are very coarse compared to the substrate, the material quality may be degraded compared to the substrate. The average grain size of the original austenite based on the CGHAZ is preferably 200 μm or smaller, and when the size is larger than this range, the hardenability increases excessively, thereby increasing the cross-sectional hardness; therefore, the average grain size of the CGHAZ is preferably 200 μm or smaller. The average grain size of the original austenite before and after PWHT may not differ. To change the grain size, a reverse transformation to the austenite single-phase range should be performed; therefore, heating to 800°C or higher should be performed. However, since the PWHT temperature is about 500°C to 600°C, there may be no change in the AGS itself.

[0152] The grain boundary occupancy of cementite present in the weld heat-affected zone (HAZ) can be 10% or less.

[0153] Grain boundary occupancy can be understood as the ratio of the length occupied by the film-like cementite within the grain boundary, and in three dimensions as the area ratio. As an example of occupancy measurement, the sample is etched and observed using SEM or similar methods to measure ordinary grain boundaries and grain boundaries that produce film-like cementite for distinction.

[0154] The steel disclosed herein ensures the strength and low-temperature impact toughness of the substrate after PWHT. Specifically, the steel may have a yield strength of 690 MPa or greater and a tensile strength of 770 MPa to 940 MPa, as well as a low-temperature impact toughness absorption energy of 50 J or greater when subjected to a Charpy V-notch test at -40°C or lower.

[0155] Meanwhile, the weld heat-affected zone after PWHT, especially the CGHAZ, ensures excellent low-temperature impact toughness and CTOD properties. Specifically, the low-temperature impact toughness absorption energy during Charpy V-notch testing at -40°C can be 40 J or greater, and the CTOD physical properties at -35°C or lower can be 0.1 mm or greater.

[0156] The following section will describe in detail an example of a method for manufacturing steel according to the present disclosure.

[0157] This manufacturing method may include producing a steel billet that satisfies the alloy composition and the Ceq range of Equation 1 described above, and then heating, hot rolling, reheating, quenching, and tempering the billet to produce steel. This includes a reheating-quenching-tempering process, which may be referred to as the so-called RQT process.

[0158] Steel manufactured in this way can be subjected to welding and post-weld heat treatment (PWHT).

[0159] The process will be described in detail below.

[0160] Steel billet manufacturing

[0161] A steel billet is manufactured in which the alloy composition described above and the Ceq value defined in Relation 1 in the range of 0.50 to 0.70 are satisfied. The maximum size of the CaO-Al₂O₃-X (X is one or more of Ti, Mg, and S) composite oxide inclusions observed in the upper portion is 50 μm or less. This upper portion is located in the region t / 4 of the surface (where t is the thickness in mm) from the surface layer in the thickness direction, and per 1 μm... 2 The number of composite oxide inclusions per unit area is 50 or fewer.

[0162] As an example of its use in manufacturing steel billets, calcium (Ca) raw materials are added to refined molten steel, and a cleaning bubbling process can be performed. As a specific example, after a second refining process, metallic Ca wire is added to the molten steel, with the amount of Ca added ranging from 0.015 kg / ton to 0.15 kg / ton. The metallic Ca wire is formed from steel coated with a Ca alloy, the thickness of which can be from 1.2 mm to 1.4 mm, and the addition rate of the Ca wire can be from 100 m / min to 300 m / min.

[0163] After adding Ca, cleaning bubbling can be performed to purge the inert gas in the ladle at a rate of 10 liters / minute to 50 liters / minute. There can be two inert gas purge points in the ladle, and the cleaning bubbling time can range from 5 minutes to 40 minutes.

[0164] In this disclosure, the process prior to the second refining is not particularly limited, and ordinary methods can be applied. Similarly, the process after Ca addition and the cleaning bubbling process is not particularly limited, and the molten steel can be cooled under ordinary conditions to produce a slab. Furthermore, when the amount of Al2O3 in the molten steel increases, the formation and coarsening of inclusions containing both Ca and Al occur, thereby increasing the amount of inclusions that break during rolling. This may not guarantee the low-temperature impact toughness and CTOD quality of the HAZ portion. Therefore, the total amount of inclusions in the molten steel before Ca addition can be limited to 2 ppm to 5 ppm.

[0165] The process of adding Ca and cleaning the bubbling process will be described in detail below.

[0166] (1) Addition of calcium (Ca)

[0167] During the billet manufacturing process, after the second refining, metallic Ca wire is added to the molten steel at an amount of 0.015 kg / ton to 0.15 kg / ton. The metallic Ca wire is formed from steel coated with a Ca alloy, the thickness of which can be 1.2 mm to 1.4 mm, and the addition rate of the Ca wire can be 100 m / min to 300 m / min.

[0168] During the addition of Ca, when the thickness of the Ca wire covering the Ca alloy steel is less than 1.2 mm, Ca melts in the upper part of the ladle, thereby reducing the constant pressure of iron and deteriorating the actual Ca yield, thus increasing the addition amount. However, when the thickness is greater than 1.4 mm, the Ca wire even contacts the bottom of the ladle, leading to the dissolution and loss of the ladle's refractory material, and therefore, operational stability may not be guaranteed.

[0169] When Ca is added to molten steel, the rate at which the wire is added should be controlled to ensure the actual Ca yield while maintaining the thickness of the Ca wire. When the addition rate is less than 100 m / min, Ca melts in the upper part of the ladle, reducing the effect of constant iron pressure and thus deteriorating the actual Ca yield, necessitating an increase in the addition rate. However, when the addition rate exceeds 300 m / min, the Ca wire may even contact the bottom of the ladle, leading to the dissolution and loss of the ladle's refractory material, which may compromise operational stability. Furthermore, the addition rate is preferably between 120 m / min and 300 m / min, and more preferably between 140 m / min and 180 m / min.

[0170] When the amount of Ca added is too small, MnS will appear at the center, and low-temperature impact toughness may not be guaranteed due to segregation defects. When the amount is too large, Ca reacts with Al2O3 in the refractory material, thereby accelerating the dissolution loss of the refractory material, and thus, operational stability may not be guaranteed. Therefore, considering the above problems, the amount of Ca added can be from 0.015 kg / ton to 0.15 kg / ton, preferably from 0.015 kg / ton to 0.10 kg / ton, and more preferably from 0.050 kg / ton to 0.10 kg / ton.

[0171] (2) Clean the bubbling

[0172] After adding Ca, cleaning bubbling can be performed for 5 to 40 minutes to ensure that the inert gas blowing rate in the ladle is 10 to 50 liters per minute. The number of inert gas blowing points in the ladle can be two.

[0173] When the inert gas flow rate in the ladle is too low, the amount of Al2O3 clusters adhering to the inert gas and removed, as well as composite inclusions containing both Ca and Al, decreases, thus deteriorating the cleanliness and potentially compromising the impact toughness of the HAZ portion. However, when the flow rate is excessive, the stirring force increases, and exposed molten metal appears on the surface of the molten steel, along with the incorporation of slag, further deteriorating the cleanliness. Therefore, the inert gas flow rate can be from 10 liters / min to 50 liters / min, preferably from 15 liters / min to 40 liters / min, and more preferably from 20 liters / min to 30 liters / min.

[0174] Meanwhile, when there is only one inert gas blowing point in the ladle, there is an uneven region in the ladle in the molten steel, and the removal capacity of Al2O3 clusters and composite inclusions containing both Ca and Al is deteriorated. When there are three or more, the stirring force is increased due to the overlap during gas blowing. Therefore, the appearance of exposed molten metal on the surface of the molten steel and the incorporation of slag are shown, which deteriorates the cleanliness.

[0175] When the cleaning bubbling time is too short, despite the limitation on the amount of inert gas blown into the ladle, the amount of Al2O3 clusters adhering to the inert gas and removed, as well as composite inclusions containing both Ca and Al, decreases, thus deteriorating the cleanliness and potentially compromising toughness. However, when the time is too long, it increases the temperature drop in the molten steel, creating a temperature gradient in the ladle, which also deteriorates the cleanliness. The cleaning bubbling time can be from 5 to 40 minutes, preferably from 7 to 17 minutes, and more preferably from 10 to 15 minutes.

[0176] billet heating

[0177] Steel billets can be heated to temperatures ranging from 1050°C to 1300°C.

[0178] Preferably, the slab is heated at 1150°C or higher to maximize the austenite grain size by re-solutioning the Ti or Nb carbonitrides, TiNb(C,N) coarse crystals, etc., formed during casting, and then rolling to a specific size to heat the austenite to the recrystallization temperature or higher, followed by holding at that temperature. However, when the temperature is too high, problems may arise due to oxide scale at high temperatures, and grains may overgrow, thereby deteriorating low-temperature impact toughness. Therefore, the upper limit can be 1300°C.

[0179] Hot rolling

[0180] Heated steel billets can be hot rolled at a fine hot rolling temperature of 830℃ to 1050℃.

[0181] 830°C corresponds to the recrystallization temperature or higher. As the strength of the steel sheet increases with decreasing rolling temperature, the composite inclusions generated during refining must be deformed by rolling; therefore, they are broken or segmented into inclusions of smaller size, or inclusions such as MnS may be elongated. Since the inclusions themselves act as a direct cause of the initiation and extension of low-temperature impact toughness, fine hot rolling at 830°C or higher (where work hardening does not occur) is preferred. However, when the temperature is above 1050°C, austenite grain growth continues even after rolling is complete, and therefore, impact toughness may deteriorate due to the increased ductile-brittle transition temperature (DBTT).

[0182] The thickness of the steel plate after hot rolling can be from 5 mm to 100 mm, more preferably from 5 mm to 70 mm, and even more preferably from 5 mm to 50 mm.

[0183] Steel sheets manufactured by hot rolling can be air-cooled to room temperature. While air cooling is the standard procedure after hot rolling, accelerated cooling to a suitable temperature can be performed after rolling to inhibit precipitate growth and reduce process time.

[0184] In this disclosure, during slab manufacturing, when the dissolved hydrogen content in the molten steel is 1.3 ppm or greater, multi-stage cumulative cooling from 200°C or higher to room temperature may be included based on the surface temperature after hot rolling and before reheating heat treatment. When cooling is performed in a multi-stage cumulative manner, the dissolved hydrogen in the steel is released, thereby further and more effectively suppressing hydrogen-induced internal microcracks, thus ultimately improving the low-temperature impact toughness of the substrate and the weld heat-affected zone.

[0185] Reheating

[0186] After hot rolling, the steel plate can be reheated to a temperature of 820°C to 950°C and held for 10 to 40 minutes.

[0187] During reheating, when the heating temperature is below 820°C or the holding time is less than 10 minutes, carbides or grain boundary segregation impurities generated during cooling after rolling cannot be well returned to a solid solution state, meaning they may remain even after cooling, and this could be a major cause of low-temperature impact toughness degradation. Meanwhile, when the heating temperature is above 950°C or the holding time is longer than 40 minutes, the original austenite grain size becomes excessive due to grain growth, which could be a major cause of toughness degradation due to increased DBTT. Therefore, the reheating temperature is preferably 820°C to 950°C, more preferably 850°C to 890°C.

[0188] Quenching and tempering

[0189] It can be used to quench and temper steel plates that have been reheated.

[0190] The cooling rate during quenching is one of the most important process factors determining the microstructure of the substrate. In this disclosure, after reheating, the temperature at point t / 4 (t: thickness in mm) based on the steel thickness can be cooled to 300°C at a cooling rate of 3°C / s to 80°C / s. When cooling is completed at 300°C or higher, the microstructure of the substrate may contain martensitic-austenitic (MA) components due to incomplete cooling transformation, and these components do not dissolve during subsequent tempering and PWHT processes, thus contributing to the deterioration of low-temperature impact toughness. Therefore, the cooling end temperature is preferably 300°C or lower.

[0191] Meanwhile, when the cooling rate is less than 3°C / second, the matrix structure of the substrate is formed by a mixed structure of martensite and bainite, or a mixed structure of ferrite and bainite, or a single-phase bainite structure as required in this disclosure, instead of a single-phase martensite structure. When bainite is the main structure, the strength and low-temperature impact toughness before PWHT can be ensured. However, due to the formation of grain boundary cementite after PWHT, it is not easy to ensure low-temperature impact toughness and CTOD quality. At the same time, when the cooling rate is greater than 80°C / second, the microstructure is formed by 100% martensite, the cross-sectional hardness increases, making it impossible to control by tempering temperature alone, and there may be a risk of plate fracture due to the deviation in cooling rate at the center of the surface layer.

[0192] Therefore, the cooling rate during quenching is preferably 3°C / second or greater and 80°C / second or less, more preferably 5°C / second or greater and 30°C / second or less.

[0193] The quenched steel sheet is heated to a temperature of 550°C to 700°C and then subjected to a tempering heat treatment for 5 to 60 minutes. When the tempering temperature is below 550°C or the holding time is less than 5 minutes, dislocation recovery during quenching cannot occur properly, resulting in an excessive increase in the strength of the substrate, which may make it impossible to ensure sufficient low-temperature impact toughness. When the temperature is above 700°C or the holding time is longer than 60 minutes, the grain boundaries (cementite) in the weld heat-affected zone may coarsen, making it difficult to guarantee the toughness after PWHT.

[0194] In addition to RQT, steel can also be manufactured by directly quenching and tempering after hot rolling (the so-called DQT).

[0195] The manufactured steel is welded and then subjected to post-weld heat treatment.

[0196] welding

[0197] The steel disclosed herein can be used after post-weld heat treatment (PWHT) following welding, and the steel disclosed herein can ensure the strength and low-temperature impact toughness of the base material after PWHT, and ensure the low-temperature impact toughness and low-temperature CTOD characteristics of the weld heat-affected zone (HAZ).

[0198] When steel is welded, there exists a heat-affected zone (HAZ) consisting of an unaffected base material, a molten and bonded portion (weld metal), and the base material affected by the welding heat. The welded portion and the HAZ are collectively referred to as the welded area. Based on the maximum heating temperature and microstructure, the HAZ can be classified into coarse-grained HAZ (CGHAZ), subcritical HAZ (SCHAZ), intercritical HAZ (ICHAZ), and fine-grained HAZ (FGHAZ). Specifically, by maintaining a high temperature of 1000°C or higher, the CGHAZ is cooled from the HAZ during welding to a region with the highest original austenite grain size. With a larger grain size, the ductile-brittle transition temperature increases, and hardenability also increases; therefore, when a hard structure is formed, the brittle fracture resistance decreases. Thus, the CGHAZ has the lowest toughness values ​​among the HAZs, such as low-temperature impact toughness and CTOD, and CGHAZ characteristics can be significant.

[0199] Welding can be performed using methods known in the art to which this disclosure pertains, and is not necessarily limited to any particular method. In this disclosure, the welding thermal regime can be as follows: heating to a maximum temperature of 1250°C to 1400°C at a heating rate of up to 50°C / second, holding at the target temperature for 5 to 10 seconds, performing a first cooling to 500°C at a cooling rate of 10°C / second to 20°C / second, and then performing a second cooling to 200°C at a cooling rate of 5°C / second to 15°C / second. Based on the process of measuring the thermal regime at the CGHAZ position using thermocouples during actual welding, the thermal regime related to the maximum temperature and cooling rate can be set as follows.

[0200] During actual welding processes such as FCAW and SAW, the welding thermal regime is implemented as consistently as possible by actually measuring the thermal regime of the weld heat-affected zone (HAZ). As a practical example of the thermal regime, the heat input for flux-cored wire arc welding (FCAW) can be from 1.0 kJ / mm to 2.0 kJ / mm, and for submerged arc welding (SAW) it can be from 2.5 kJ / mm to 4.0 kJ / mm. In this paper, both the preheating temperature and the interpass temperature can be from 150°C to 250°C.

[0201] Post-weld heat treatment (PWHT)

[0202] Post-weld heat treatment (PWHT) can be performed by methods known in the art to which this disclosure pertains, and there are no particular limitations. As an example, according to ASME Sec. VIII, PWHT can be performed by holding at 595°C to 625°C for 180 to 540 minutes. Meanwhile, the total LMP in the tempering heat treatment and PWHT heat treatment according to the following Relationships 2 and 3 is preferably 17.0 to 19.5.

[0203] Specifically, when the tempering temperature and the PWHT temperature are the same, the LMP is determined by Equation 2, while when the tempering temperature and the PWHT temperature are different, the LMP is determined by Equation 3:

[0204] [Relationship 2]

[0205]

[0206] Where T k =Tempering and PWHT temperatures in Kelvin, and t =Tempering and PWHT time (hours).

[0207] [Relationship 3]

[0208]

[0209]

[0210] Where T i The temperature of PWHT in Kelvin, t i PWHT time (hours), T: tempering temperature in Kelvin, and t eq Equivalent to tempering time (in hours) converted to PWHT temperature.

[0211] When LMP is less than 17.0, as mentioned above, the dislocation density cannot be reduced, which may lead to excessive strength defects, or the cross-sectional hardness may be too high, which may degrade impact toughness. However, when LMP is greater than 19.5, the grain boundary occupancy of cementite after PWHT is too high, making it difficult to properly ensure impact toughness and CTOD quality. Therefore, LMP is preferably 17.0 to 19.5, and more preferably 18.0 to 19.3.

[0212] Invention Embodiments

[0213] In the following sections, embodiments of this disclosure will be described. It will be apparent that those skilled in the art to which this disclosure pertains can make various modifications to the following embodiments without departing from the scope of this disclosure. The following embodiments are provided for understanding this disclosure, and the scope of this disclosure is not limited to the following embodiments but rather to the described claims and their equivalents.

[0214] (Example 1)

[0215] Under the conditions in Table 2, molten steel having the alloy composition in Table 1 (in Table 1, the content of the component elements is in wt% but P, ​​S and Ca are in ppm, and the remaining components are Fe and unavoidable impurities) and Ceq of the above relation 1 is made into a billet, and heated and hot-rolled to produce a hot-rolled steel plate with a thickness of 50 mm.

[0216] The manufactured hot-rolled steel sheet was reheated and quenched under the conditions disclosed in Table 3, and then tempered at 610°C for 20 minutes.

[0217] Meanwhile, to determine the excellent strength and low-temperature impact toughness even after post-weld heat treatment, a heat treatment process simulating post-weld heat treatment was performed. In this case, the heat treatment process was carried out by subjecting the tempered hot-rolled steel sheet to three cycles at 595°C for 180 minutes each.

[0218] The microstructure and physical properties of the substrates of each sample manufactured as described above were measured, and the results are shown in Tables 4 and 5 below.

[0219] In this paper, the grain size and phase fraction of the original austenite were measured by collecting samples and using electron backscatter diffraction (EBSD). Precipitates and inclusions were then measured using transmission electron microscopy (TEM). Precipitates and inclusions were determined by diffraction patterns and EDX mapping.

[0220] Yield strength / tensile strength was assessed by room temperature tensile testing, with a 0.2% offset applied to the yield strength. Furthermore, impact toughness was measured for each specimen by Charcot V-notch testing, using the average of three absorbed energy values ​​measured at the corresponding temperature. Brinell hardness was measured according to tests according to ISO 6506 Part 1.

[0221] [Table 1]

[0222]

[0223] [Table 2]

[0224]

[0225] [Table 3]

[0226]

[0227] [Table 4]

[0228]

[0229] [Table 5]

[0230]

[0231] As can be seen from Tables 1 to 5, all specimens numbered 1 to 5 and 17 to 20 that meet the alloy composition and manufacturing conditions proposed in this disclosure ensure excellent quality through base material strength and excellent low-temperature impact toughness at -40°C, even after excellent post-weld heat treatment.

[0232] However, for samples 6 to 16 that meet the alloy composition specified in this disclosure but do not meet the manufacturing conditions, it is difficult to ensure the properties of the substrate specified in this disclosure. Furthermore, for samples 21 to 25 that meet the steel manufacturing conditions specified in this disclosure but do not meet the alloy composition, it is difficult to ensure the properties of the substrate specified in this disclosure.

[0233] (Example 2)

[0234] The steels manufactured under the conditions in Table 3 were welded and subjected to post-weld heat treatment, and the low-temperature toughness and CTOD characteristics of the CGHAZ in the heat-affected zone (HAZ) were determined.

[0235] Specifically, steel manufactured according to Tables 1 to 3 was subjected to welding simulation under the welding heat regime conditions specified in Table 6, and then subjected to a post-weld heat treatment process of 3 cycles at 595°C for 180 minutes. In Table 6, LMP is derived from the aforementioned Equation 3.

[0236] The microstructure and physical properties of CGHAZ were measured in the various specimens manufactured as described above, and the results are shown in Tables 7 and 8.

[0237] In this paper, the phase fraction is measured by collecting samples and using electron backscatter diffraction (EBSD).

[0238] Impact toughness of each specimen was measured using the Charpy V-notch test, with the average of three absorbed energy values ​​measured at the corresponding temperature. Brinell hardness was measured by performing tests according to ISO 6506 Part 1, and CTOD tests were performed according to BS 5762 (British Standard).

[0239] In addition, transmission electron microscopy (TEM) was used to measure the cementite grain boundary occupancy. The cementite grain boundary occupancy is measured as the ratio of the total grain boundary length occupied by film-like cementite.

[0240] [Table 6]

[0241]

[0242] [Table 7]

[0243]

[0244] [Table 8]

[0245]

[0246] As can be seen from Tables 6 to 8, the specimens numbered 1 to 5 and 11 to 16 that meet the alloy composition and manufacturing conditions proposed in this disclosure ensured good CTOD quality and excellent low-temperature impact toughness in CGHAZ.

[0247] However, while the alloy composition meets the requirements of this disclosure, the steel manufacturing conditions are not met, and in particular, samples 6 to 10, which have large-sized composite oxide inclusions, exhibit poor physical properties of CGHAZ. Meanwhile, samples 17 to 20, which meet the alloy composition and steel manufacturing conditions of this disclosure but do not meet the welding and post-weld heat treatment conditions, cannot guarantee the physical properties of CGHAZ.

[0248] Samples 21 to 25, which met the manufacturing conditions set forth in this disclosure but did not meet the alloy composition requirements, were found to have failed to meet the appropriate quality requirements as specified in this disclosure.

[0249] While exemplary embodiments have been shown and described above, it will be apparent to those skilled in the art that modifications and variations may be made without departing from the scope of this disclosure as defined by the appended claims.

Claims

1. A weld heat-affected zone comprising, by weight percent: 0.05% to 0.20% C, 0.05% to 0.50% Si, 0.5% to 2.0% Mn, 0.005% to 0.10% Al, 0.010% or less P, 0.0050% or less S, 0.001% to 0.070% Nb, 0.001% to 0.30% V, 0.0 0.01% to 0.030% Ti, 0.01% to 1.0% Cr, 0.01% to 1.0% Mo, 0.01% to 0.60% Cu, 1.0% to 4.0% Ni, 0.001% to 0.005% B, 0.0005% to 0.0040% Ca, 0.002% or less O, and the remainder Fe and other unavoidable impurities. Ceq, defined by relation 1, satisfies a value in the range of 0.50 to 0.

70. The largest size of the CaO-Al2O3-X (X is one or more of Ti, Mg, and S) composite oxide inclusions observed in the upper layer is 50 μm or smaller, in the region t / 4 of the surface (where t is the thickness in mm) from the surface layer to the thickness direction, and per 1 μm 2 The number of the composite oxide inclusions per unit area is 50 or fewer. The cementite present in the grain boundaries has a grain boundary occupancy of 10% or less, and the average size of the cementite is 1 μm or less. The cross-sectional hardness is 200 HB to 300 HB. [Relation 1] Where [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] are the contents (by weight%) of C, Mn, Cr, Mo, V, Ni and Cu in the steel, respectively, and 0 is used when the component is not intentionally included.

2. The weld heat-affected zone according to claim 1, wherein the weld heat-affected zone is a coarse-grained heat-affected zone (CGHAZ).

3. The weld heat-affected zone according to claim 1, wherein the microstructure of the weld heat-affected zone comprises 50% or more tempered martensite by area fraction and bainite in the remainder.

4. The weld heat-affected zone according to claim 1, wherein the average grain size of the original austenite in the weld heat-affected zone is 200 μm or less.

5. The weld heat-affected zone according to claim 1, wherein the weld heat-affected zone has a Charpy impact absorption energy of 40 J or greater at -40°C and a CTOD characteristic of 0.1 mm or greater at -35°C.

6. A method for manufacturing a weld heat-affected zone, the method comprising: A steel product is prepared, wherein the steel product comprises, by weight percent: 0.05% to 0.20% C, 0.05% to 0.50% Si, 0.5% to 2.0% Mn, 0.005% to 0.10% Al, 0.010% or less P, 0.0050% or less S, 0.001% to 0.070% Nb, 0.001% to 0.30% V, 0.001% to 0.030% Ti, 0.01%... The steel contains 1% to 1.0% Cr, 0.01% to 1.0% Mo, 0.01% to 0.60% Cu, 1.0% to 4.0% Ni, 0.001% to 0.005% B, 0.0005% to 0.0040% Ca, 0.002% or less O, and the remainder Fe and other impurities. The steel satisfies a Ceq in the range of 0.50 to 0.70 as defined by Relation 1, and has a maximum size of 50 μm or less of CaO-Al2O3-X (X being one or more of Ti, Mg, and S) complex oxide inclusions observed in the upper portion, and 50 or fewer inclusions per 1 μm. 2 The number of inclusions per unit area, in the upper portion of the region t / 4 (where t is the thickness in mm) from the surface layer in the thickness direction; Welding the steel; and The welded steel is then subjected to post-weld heat treatment: [Relation 1] Wherein [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] are the contents (by weight%) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, and 0 is used when the component is not intentionally included.

7. The method for manufacturing a weld heat-affected zone according to claim 6, wherein the welding thermal regime comprises heating to a maximum temperature of 1250°C to 1400°C at a heating rate of up to 50°C / second, holding at the target temperature for 5 to 10 seconds, performing a first cooling to 500°C at a cooling rate of 10°C / second to 20°C / second, and then performing a second cooling to 200°C at a cooling rate of 5°C / second to 15°C / second.

8. The method for manufacturing a weld heat-affected zone according to claim 6, wherein the welding is FCAW welding performed under welding conditions of a preheating temperature and interpass temperature of 150°C to 250°C with a heat input of 1.0 kJ / mm to 2.0 kJ / mm, or SAW welding performed under welding conditions of a preheating temperature and interpass temperature of 150°C to 250°C with a heat input of 2.0 kJ / mm to 4.0 kJ / mm.

9. The method for manufacturing a weld heat-affected zone according to claim 6, wherein the post-weld heat treatment is performed by holding at 595°C to 625°C for 180 to 540 minutes.

10. The method for manufacturing a weld heat-affected zone according to claim 6, The steel comprises 50% or more (including 100%) tempered martensite by area fraction, with the remainder consisting of bainite and other impurities as a microstructure. The lath group size of the tempered martensite and the bainite is 17 μm or smaller.

11. The method for manufacturing a weld heat-affected zone according to claim 10, wherein when the tempered martensite content is greater than 80%, the cross-sectional Brinell hardness of the steel is 220 HB to 290 HB.

12. The method for manufacturing a weld heat-affected zone according to claim 10, wherein the precipitates observed in the cross-section of the steel are present in a density of 1 μm per unit area of ​​the steel. 2 Five or more VC or VCN precipitates with a diameter of 5 nm to 15 nm.

13. The method for manufacturing a weld heat-affected zone according to claim 10, wherein the average grain size of the original austenite in the steel is 30 μm or less.

14. The method for manufacturing a weld heat-affected zone according to claim 6, wherein the steel is manufactured by a manufacturing method comprising: Manufacturing steel billets; The steel billet is then reheated to a temperature in the range of 1050°C to 1300°C; The steel billet is hot-rolled at a fine hot-rolling temperature of 830°C to 1050°C to manufacture steel products; The steel is heated to a temperature ranging from 820°C to 950°C, held at that temperature for 10 to 40 minutes, and then cooled to 300°C or lower at a cooling rate of 3°C / sec to 80°C / sec, based on the temperature at t / 4 of the steel thickness; and After cooling, the steel is subjected to a tempering heat treatment in which it is heated to 550°C to 700°C and held at that temperature for 5 to 60 minutes.

15. The method for manufacturing a weld heat-affected zone according to claim 14, wherein the manufacturing of the steel billet comprises: After the second refining, metallic Ca wire is added to the molten steel, such that the amount of Ca added is 0.015 kg / ton to 0.15 kg / ton; as well as After adding Ca, perform cleaning bubbling for 5 to 40 minutes to ensure that the inert gas flow rate in the ladle is 10 to 50 liters per minute. The Ca wire is added at a rate of 100 m / min to 300 m / min, and the ladle has two inert gas blowing points.

16. The method for manufacturing a weld heat-affected zone according to claim 14, wherein when the tempering heat treatment and the post-weld heat treatment temperatures are the same, the LMP according to the following relationship 2 is 17.0 to 19.5: [Relationship 2] Where T k =Tempering and PWHT temperatures in Kelvin, and t =Tempering and PWHT time (hours).

17. The method for manufacturing a weld heat-affected zone according to claim 14, wherein when the tempering heat treatment and the post-weld heat treatment temperatures are different, the LMP according to the following relationship 3 is 17.0 to 19.5: [Relationship 3] Where T i The temperature of PWHT in Kelvin, t i PWHT time (hours), T: tempering temperature in Kelvin, and t eq Equivalent to the tempering time (in hours) converted to the PWHT temperature.

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