Steel, welding heat affected zone and manufacturing method thereof

Abstract

The present invention relates to a steel applied to a liquefied gas storage tank and the like, a welding heat affected zone, and a method for manufacturing same.

Description

Steel, weld heat affected zone and manufacturing method thereof

[0001] The present invention relates to steel materials, welded heat affected zones, and methods for manufacturing the same, which are applicable to liquefied gas storage tanks and the like.

[0002] With the recent strengthening of environmental regulations, demand for liquefied gas (LNG), an environmentally friendly fuel, is increasing. This has led to an increase in the construction of LNG carriers and propulsion vessels that use LNG as fuel. Among these liquefied gases, LNG and Lethylene / Ethane have very low liquefaction temperatures, so high-Ni steel is generally used. However, LPG, ammonia, and CO2 have liquefaction temperatures of around -60°C. To reduce the construction costs of tanks and vessels that accommodate these gases, carbon steel with enhanced strength and low-temperature toughness is increasingly being used.

[0003] Existing liquefied gas transport tanks primarily utilize the large Type A tank design, which does not require high pressure, and therefore do not require high strength or thicker materials. However, recent CO2 tanks and ammonia fuel tanks, which require high pressure during liquefaction, are being designed as Type C tanks, which necessitates high strength and thicker steel materials. Furthermore, since the IGC code mandates post-weld heat treatment (PWHT) when constructing Type C tanks, additional post-PWHT property guarantees are required for steel.

[0004] The above PWHT is a process that maintains the material at a high temperature for a certain period of time to relieve residual stress after welding. During the PWHT process, the grain size and precipitate size of the low-temperature structure increase, which deteriorates the material properties such as strength and toughness, making quality assurance difficult. In particular, for steels for CO2 tanks, the classification society has recently required the quality of CTOD (Crack Tip Opening Displacement) in the heat affected zone (HAZ) after PWHT as a structural compatibility evaluation item to improve the stability of tank design.

[0005] Meanwhile, in the case of liquefied CO2, the liquefaction temperature varies depending on the pressure, and accordingly, it can be broadly divided into low-pressure tanks and medium-pressure tanks.

[0006] For low-pressure tanks with typical operating pressures of 5 to 10 bar, CO2 liquefaction requires temperatures down to -60°C. Therefore, storage tank steel with a tensile strength (TS) of 600 to 650 MPa is used, ensuring low-temperature impact toughness below -60°C. Maintaining the low temperature after liquefaction requires separate cooling equipment, which increases initial installation and maintenance costs.

[0007] For medium-pressure tanks operating at 18-20 bar, the liquefaction temperature can rise to -35°C, making it easier to store and transport liquefied CO2. However, due to the high operating pressure, the steel used for medium-pressure tanks requires ultra-high-strength steel with a strength of 780 MPa or higher, rather than the conventional TS 600-650 MPa. To achieve this high strength, steel containing high alloy components is manufactured through quenching-tempering (QT) heat treatment, making it difficult to ensure weld quality.

[0008] In Patent Document 1, a slab containing alloy components of 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 in wt% and the remainder Fe and impurities is heated to 1100 to 1200°C, rolled at 800 to 900°C with a cumulative reduction ratio of 30% or more, air-cooled, quenched at 910 to 930°C, and then rolled at 650 to 680°C. It is described that by performing tempering heat treatment, steel can be manufactured that guarantees low-temperature impact toughness down to -40 degrees.

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

[0010] Meanwhile, Patent Document 2 describes a method for manufacturing P690QL2 steel for LCO2 storage tanks, which contains C: 0.09 to 0.12%, Mn: 0.9 to 1.3%, Ti: 0.015% or less, S: 0.002% or less, Nb: 0.06 or less, V: 0.08 or less, Ni: 1.5 to 2.5%, Cr: 0.2 to 0.4%, Mo: 0.3 to 0.4%, Cu: 0.3% or less, Al: 0.03 to 0.05%, Ca: 0.0005 to 0.005%, N: 0.005% or less, H: 0.0003% or less, and the remainder being Fe and impurities, and has a thickness of 50 mm or less and can guarantee low-temperature impact toughness. The above patent document 2 describes that when a material is heated to 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 secured after PWHT.

[0011] However, since the CTOD guarantee temperature and standard of the HAZ section do not reach the level required by the present invention, a practical method that can be used as steel for medium-pressure LCO2 storage tanks is not presented.

[0012] Patent Document 3 contains 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 in weight%, and the remainder is composed of Fe and impurities, and the content ratio of Mn and Ni, Mn / Ni, is 0.80 or less, the carbon equivalent Ceq based on WES is 0.430 to 0.900%, the tensile strength is 780 MPa or more and 930 MPa or less, the yield strength is 630 MPa or more and 750 MPa or less, and the yield ratio is 85% or less. It is described that a steel plate can be manufactured that is characterized in that the HAZ impact toughness evaluated at 0℃ is 100J or more, the plate thickness is 40mm or more and 120mm or less, and the Vickers hardness is 225 points or more at a location 1 / 4 of the plate thickness from the surface, and when the average value of the smaller 20% of the Vickers hardness is Hvmin and the average value of the larger 20% is Hvmax, Hvmin / Hvmax is 0.85 or less. The strength satisfies the range required for steel for medium-pressure LCO2, but it may not be suitable because the impact toughness is not guaranteed at around -40℃, the temperature at which CO2 liquefies. In addition, when the Ceq and hardness are excessively high, the martensite fraction becomes excessively high during quenching, making it difficult to secure impact toughness and HAZ CTOD characteristics.

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

[0014] Therefore, the above-described conventional technology does not provide a practical method for manufacturing ultra-high strength steel for medium-pressure LCO2 having a yield strength of 690 MPa or more and a tensile strength of 770 to 940 MPa after PWHT (Post Weld Heat Treatment), low-temperature impact toughness at -40°C or lower, and CTOD quality of 0.1 mm or more.

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

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

[0017] (Patent Document 3) PCT Publication No. WO2021-255858-023761 (Published on December 23, 2021)

[0018] One aspect of the present invention relates to a weld heat affected zone having excellent strength and low-temperature impact toughness as well as excellent low-temperature CTODD quality after post-weld heat treatment (PWHT), and a method for manufacturing the same.

[0019] In addition, another aspect of the present invention relates to a steel having excellent strength and low-temperature impact toughness after post-weld heat treatment (PWHT) and a method for manufacturing the same.

[0020] The objectives of the present invention are not limited to the above-described matters. Additional objectives of the present invention are described throughout the specification, and those skilled in the art will have no difficulty understanding the additional objectives of the present invention from the contents described in the specification.

[0021] One embodiment of the present invention comprises, in wt%, 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 unavoidable impurities. Including,

[0022] Ceq defined in the following relational expression 1 satisfies the range of 0.50 to 0.70,

[0023] The maximum size of the CaO-Al2O3-X (X is one or more of Ti, Mg, and S) composite oxidizing inclusions observed in the upper layer, which is the area from the surface to the surface layer in the thickness direction to t / 4 (t is the thickness, unit mm), is 50㎛ or less, and the unit area is 1㎛. 2 The number of parties is 50 or less,

[0024] The grain boundary occupancy ratio of cementite existing at the grain boundary is 10% or less, and the average size of cementite is 1㎛ or less,

[0025] The cross-sectional hardness relates to the weld heat affected zone, including those having a hardness of 200 to 300 HB.

[0026] [Relationship 1]

[0027] Ceq = [C] + [Mn] / 6 + ([Cr] + [Mo] + [V]) / 5 + ([Ni] + [Cu]) / 15

[0028] In the above relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] represent the contents (weight%) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, respectively, and 0 is substituted when these components are not intentionally added.

[0029] The above welding heat affected zone may be a coarse grained heat affected zone (CGHAZ).

[0030] The microstructure of the above-mentioned weld heat-affected zone may include tempered martensite with an area fraction of 50% or more, and the remainder may include bainite.

[0031] The average grain size of the old austenite in the above-mentioned weld heat-affected zone may be 200㎛ or less.

[0032] The above-mentioned weld heat affected zone may have a Charpy impact absorption energy of 40 J or more at -40°C and a CTOD characteristic of 0.1 mm or more at -35°C.

[0033]

[0034] Another aspect of the present invention satisfies the above composition and Ceq,

[0035] The maximum size of the CaO-Al2O3-X (X is one or more of Ti, Mg, and S) composite oxidized inclusions observed in the upper layer, which is the area from the surface to the surface layer in the thickness direction to t / 4 (t is the thickness, unit mm), is 50㎛ or less, and the unit area is 1㎛. 2 A step for preparing steel having a number of 50 or less;

[0036] A step of welding the above steel material; and

[0037] Step of post-welding heat treatment of the above welded steel

[0038] It relates to a method for manufacturing a heat affected zone of welding including:

[0039] The thermal cycle of the above welding may be to increase the temperature to a maximum temperature of 1250 to 1400°C at a maximum heating rate of 50°C / s, maintain the temperature at the target temperature for 5 to 10 seconds, perform a first cooling at a cooling rate of 10 to 20°C / s to 500°C, and then perform a second cooling at a cooling rate of 5 to 15°C / s to 200°C.

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

[0041] The above post-welding heat treatment can be performed by maintaining the temperature at 595°C to 625°C for 180 to 540 minutes (min).

[0042] The above steel may have a microstructure in which tempered martensite comprises an area fraction of 50% or more (including 100%), the remainder comprising bainite and other inevitable structures, and the packet size of the tempered martensite and bainite may be 17 ㎛ or less.

[0043] The above steel may have a steel cross-section Brinell hardness of 220 to 290 HB when tempered martensite exceeds 80%.

[0044] Among the precipitates observed in the cross-section of the steel, the above steel has VC or VCN precipitates with a diameter of 5 to 15 nm and a diameter of 1 ㎛. 2 There may be more than 5 per party.

[0045] The above steel may have an average austenite grain size of 30㎛ or less.

[0046] The above steel is,

[0047] Steps for manufacturing steel slabs;

[0048] A step of reheating the above steel slab to a temperature range of 1050 to 1300°C;

[0049] A step of manufacturing steel by hot rolling the above steel slab at a finishing hot rolling temperature of 830 to 1050°C;

[0050] A step of heating the above steel to a temperature range of 820 to 950°C and maintaining it for 10 to 40 minutes, and then cooling it to a temperature of 300°C or lower at a cooling rate of 3°C / s to 80°C / s based on the steel thickness t / 4 temperature; and

[0051] After the above cooling, it can be manufactured by a manufacturing method including a tempering heat treatment step of heating the steel to 550 to 700°C and maintaining it for 5 to 60 minutes.

[0052] The steps for manufacturing the above steel slab are:

[0053] A step of adding Metal Ca Wire to the molten steel after secondary refining so that the amount of Ca added is 0.015 to 0.15 kg / ton; and

[0054] After the above Ca injection, a step of clean bubbling for 5 to 40 minutes is included so that the amount of inert gas injected into the ladle becomes 10 to 50 liter / min.

[0055] The above Ca Wire injection speed is 100 to 300 meters / min, and there may be two inert gas injection points in the ladle.

[0056] When the temperatures of the above tempering heat treatment and post-welding heat treatment are the same, the LMP according to the following relational expression 2 may be 17.0 to 19.5.

[0057] [Relationship 2]

[0058] LMP = T k (Log t + 20)

[0059] Here, T k = Kelvin temperature of tempering and PWHT, t = tempering and PWHT time (hour)

[0060] When the temperatures of the above tempering heat treatment and post-welding heat treatment are different, the LMP according to the following relational expression 3 may be 17.0 to 19.5.

[0061] [Relationship 3]

[0062] T(logt eq +20) = T i (logt i +20)

[0063] Total LMP = T i [log(t i +t eq)+20]

[0064] Here, T i : Kelvin temperature of PWHT, t i : PWHT time (hour), T: Tempering temperature in Kelvin, t eq : Equivalent tempering time (hour) when converted to PWHT temperature

[0065]

[0066] Another aspect of the present invention satisfies the above composition and Ceq,

[0067] The microstructure contains tempered martensite with an area fraction of more than 50%.

[0068] The packet size of the above microstructure includes 17㎛ or less,

[0069] The maximum size of the composite oxidation inclusions of CaO-Al2O3-X (X is one or more of Ti, Mg, and S) observed in the upper layer, which is the area from the surface of the steel to the surface layer in the thickness direction to t / 4 (t is the thickness, unit mm), is 50㎛ or less, and the unit area is 1㎛. 2 This is for steel containing less than 50 sugars.

[0070] [Relationship 1]

[0071] Ceq = [C] + [Mn] / 6 + ([Cr] + [Mo] + [V]) / 5 + ([Ni] + [Cu]) / 15

[0072] In the above relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] represent the contents (weight %) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, respectively, and 0 is substituted when these components are not intentionally added.

[0073]

[0074] Another aspect of the present invention is a composite oxidized inclusion having a maximum size of 50 ㎛ or less of CaO-Al2O3-X (X is at least one of Ti, Mg, and S) observed in the upper layer, which is a region from the surface of the steel slab to the surface layer in the thickness direction of t / 4 (t is thickness, unit mm), and satisfying the above composition and Ceq, and having a unit area of ​​1 ㎛. 2 A step for manufacturing a steel slab having a number of 50 or less;

[0075] A step of reheating the above steel slab to a temperature range of 1050 to 1300°C;

[0076] A step of manufacturing steel by hot rolling the above steel slab at a finishing hot rolling temperature of 830 to 1050°C;

[0077] A step of heating the above steel to a temperature range of 820 to 950°C and maintaining it for 10 to 40 minutes, and then cooling it to a temperature of 300°C or lower at a cooling rate of 3°C / s to 80°C / s based on the steel thickness t / 4 temperature; and

[0078] After the above cooling, a tempering heat treatment step is performed in which the steel is heated to 550 to 700°C and maintained for 5 to 60 minutes.

[0079] It relates to a method for manufacturing steel including .

[0080] According to one aspect of the present invention, a weld heat affected zone having excellent low-temperature impact toughness and low-temperature CTOD characteristics after post-weld heat treatment (PWHT) and a method for manufacturing the same can be provided.

[0081] According to another aspect of the present invention, a steel material having excellent strength and low-temperature impact toughness and a method for manufacturing the same can be provided, in which the deterioration of physical properties is minimized even after post-weld heat treatment (PWHT).

[0082] The various advantageous and beneficial advantages and effects of the present invention are not limited to the above-described contents, and will be more easily understood in the course of explaining specific embodiments of the present invention.

[0083] The terminology used herein is for the purpose of describing the present invention and is not intended to limit the present invention. Furthermore, the singular forms used herein also include the plural forms, unless the context clearly dictates otherwise.

[0084] The meaning of "comprising" as used in the specification is to specify a configuration and not to exclude the presence or addition of other configurations.

[0085] Unless otherwise defined, all terms, including technical and scientific terms, used herein have the same meaning as commonly understood by a person of ordinary skill in the art to which this invention pertains. Terms defined in the dictionary are to be interpreted to have meanings consistent with the relevant technical literature and the present disclosure.

[0086] Preferred embodiments of the present invention are described below. These embodiments may be modified in various ways, and the scope of the present invention should not be construed as being limited to the embodiments described below. These embodiments are provided to further illustrate the present invention to those skilled in the art.

[0087] First, the composition of the steel material, which is an example of the present invention, will be described in detail.

[0088] In weight %, C: 0.05 to 0.20%, Si: 0.05 to 0.50%, Mn: 0.5 to 2.0%, Al: 0.005 to 0.10%, P: 0.010% or less, S: 0.0050% or less, Nb: 0.001 to 0.070%, V: 0.001 to 0.30%, Ti: 0.001 to 0.030%, Cr: 0.01 to 1.0%, Mo: 0.01 to 1.0%, Cu: 0.01 to 0.60%, Ni: 1.0 to 4.0%, B: 0.001 to 0.005%, Ca: 0.0005 to 0.0040%, O: 0.002% or less, including the remainder of Fe and other unavoidable impurities.

[0089] Ceq, defined by the following relational expression 1, can be 0.50 to 0.70.

[0090] [Relationship 1]

[0091] Ceq = [C] + [Mn] / 6 + ([Cr] + [Mo] + [V]) / 5 + ([Ni] + [Cu]) / 15

[0092] In the above relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] represent the contents (weight%) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, respectively, and 0 is substituted when these components are not intentionally added.

[0093] The above steel composition may be the same as the composition of the heat affected zone (HAZ) of the weld as well as the base material of the steel slab, the steel before welding, the steel after welding and the steel after welding heat treatment.

[0094] Hereinafter, the above composition will be described in detail. Hereinafter, unless otherwise specified, the % and ppm described in relation to the alloy composition are based on weight.

[0095] Carbon (C): 0.05~0.20%

[0096] Carbon (C) is the most important element for securing basic strength, so it needs to be contained in steel within an appropriate range. To obtain this additive effect, 0.05% or more of carbon (C) can be added. Preferably, 0.10% or more of carbon (C) can be included. On the other hand, if the carbon (C) content exceeds a certain level, the fraction of martensite becomes excessively high, the hardness of the heat affected zone (HAZ) increases, and the fraction of grain boundary cementite increases, so that the impact propagation resistance decreases in the impact test or CTOD test, so that the toughness value required in the present invention cannot be secured. Therefore, the present invention can be performed with a carbon (C) content of 0.20%. A more preferable upper limit of the carbon (C) content may be 0.18%.

[0097] Silicon (Si): 0.05~0.50%

[0098] Silicon (Si) is a substitutional element that enhances the strength of steel through solid solution strengthening and has a strong deoxidation effect, making it an essential element for manufacturing clean steel. Therefore, silicon (Si) may be contained in an amount of 0.05% or more, and preferably 0.20% or more. On the other hand, if silicon (Si) is contained in a large amount, it may form a martensite-austenite (MA) phase and excessively increase the matrix strength of the martensite phase, thereby deteriorating low-temperature impact toughness. Therefore, the upper limit of the content may be contained as 0.50%. A more preferable upper limit of the silicon (Si) content may be 0.40%.

[0099] Manganese (Mn): 0.5~2.0%

[0100] Manganese (Mn) is a useful element that improves strength through solid solution strengthening and enhances hardenability by forming low-temperature transformation phases. Therefore, to secure a yield strength of 690 MPa or higher, it is desirable to include manganese (Mn) of 0.5% or more. A more desirable manganese (Mn) content may be 0.7% or more. On the other hand, manganese (Mn) can form non-metallic inclusions, MnS, which are elongated together with sulfur (S), which reduces toughness and can act as an impact initiation point, which can be a factor that drastically reduces the low-temperature impact toughness of steel. Therefore, it is desirable to manage the manganese (Mn) content to be 2.0% or less, and a more desirable manganese (Mn) content may be 1.5% or less.

[0101] Aluminum (Al): 0.005~0.10%

[0102] Aluminum (Al), along with silicon (Si), is one of the powerful deoxidizing agents in the steelmaking process. To achieve this effect, it is desirable to include it at 0.005% or more. A more desirable lower limit of the aluminum (Al) content may be 0.01%. On the other hand, if the aluminum (Al) content is excessive, the fraction of Al2O3 among the oxidizing inclusions generated as a result of deoxidation increases excessively, making the size of the inclusions coarse and making it difficult to remove the inclusions during refining, which may become a factor in lowering the low-temperature impact toughness. Therefore, it is desirable to manage the aluminum (Al) content at 0.10% or less. A more desirable aluminum (Al) content may be 0.07% or less.

[0103] Phosphorus (P): 0.010% or less (including 0%) and Sulfur (S): 0.0050% or less (including 0%)

[0104] Phosphorus (P) and sulfur (S) are elements that cause embrittlement at grain boundaries or form coarse inclusions, thereby causing embrittlement. Therefore, it is preferable not to include phosphorus (P) and sulfur (S). To improve brittle crack propagation resistance, it is preferable to include phosphorus (P) at 0.010% or less and sulfur (S) at 0.0050% or less.

[0105] Niobium (Nb): 0.001~0.070%

[0106] Niobium (Nb) is an element that improves the strength of the base metal by precipitating in the form of NbC or NbCN. In addition, niobium (Nb) dissolved during high-temperature reheating precipitates very finely in the form of NbC during rolling, which inhibits recrystallization of austenite, thereby having the effect of refining the structure. Therefore, it is preferable that niobium (Nb) be contained in an amount of 0.001% or more, and a more preferable niobium (Nb) content may be 0.005% or more. On the other hand, if niobium (Nb) is contained excessively, undissolved niobium (Nb) is generated in the form of TiNb(C,N), which becomes a factor that lowers low-temperature impact toughness. Therefore, the upper limit of the niobium (Nb) content is preferably 0.070%. A more preferable upper limit of the niobium (Nb) content may be 0.065% or less.

[0107] Vanadium (V): 0.001–0.30%

[0108] Vanadium (V) is almost entirely re-dissolved during reheating, so its strengthening effect through precipitation or solid solution during subsequent rolling is minimal. However, it has the effect of improving strength by precipitating very fine carbonitrides during tempering heat treatment or PWHT. To fully achieve this effect, it is necessary to add vanadium (V) of 0.001% or more. A more desirable lower limit of the vanadium (V) content may be 0.01%. On the other hand, if the content is excessive, the high hardenability may excessively increase the surface hardness of the slab, which may cause surface cracks during flange processing, and the manufacturing cost may also increase sharply, which is not commercially advantageous. Therefore, the vanadium (V) content may be limited to 0.3% or less. A more desirable vanadium (V) content may be 0.25% or less.

[0109] Titanium (Ti): 0.001~0.030%

[0110] Titanium (Ti) is a component that significantly improves low-temperature toughness by precipitating as TiN during reheating and inhibiting the growth of old austenite grains at high temperatures. To achieve this effect, it is desirable to contain 0.001% or more of titanium (Ti). On the other hand, if titanium (Ti) is contained excessively, low-temperature toughness may be reduced due to clogging of the casting nozzle or crystallization at the center. In addition, titanium (Ti) combines with nitrogen (N) to form coarse TiN precipitates at the center of the thickness, which reduces the elongation of the product, thereby reducing the uniform elongation during the forging process and causing surface cracks. Therefore, the titanium (Ti) content may be 0.030% or less. The upper limit of the desirable titanium (Ti) content may be 0.025% or less, and the more desirable titanium (Ti) content may be 0.018% or less.

[0111] Chromium (Cr): 0.01~1.0%

[0112] Chromium (Cr) is an element that increases yield resistance 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 rate of spheroidization of cementite. For this effect, 0.01% or more of chromium (Cr) may be included. On the other hand, if the chromium (Cr) content is excessive, M 23 As the size and fraction of Cr-rich coarse carbides such as C6 increase, the impact toughness of the steel decreases, and the solubility of niobium (Nb) in the steel and the fraction of fine precipitates such as NbC decrease, so that the strength of the product may decrease. Therefore, the upper limit of the chromium (Cr) content in the present invention may be 1.0%. The upper limit of the preferable chromium (Cr) content may be 0.8%.

[0113] Molybdenum (Mo): 0.01~1.0%

[0114] Molybdenum (Mo) is an element that increases grain boundary strength and has a significant solid-solution strengthening effect within ferrite, effectively contributing to the strength and ductility of products. Furthermore, molybdenum (Mo) prevents the deterioration of toughness caused by grain boundary segregation of impurity elements such as phosphorus (P). For this effect, 0.01% or more of molybdenum (Mo) may be added. However, molybdenum (Mo) is an expensive element, and excessive addition can significantly increase manufacturing costs. Therefore, the upper limit of the molybdenum (Mo) content is preferably 1.0%.

[0115] Copper (Cu): 0.01~0.60%

[0116] Copper (Cu) is an advantageous element in the present invention because it can significantly improve the strength of the matrix through solid solution strengthening within ferrite, and also has the effect of inhibiting corrosion in a wet hydrogen sulfide atmosphere. For this effect, 0.01% or more of copper (Cu) may be included. A more preferable copper (Cu) content may be 0.03% or more. However, if the copper (Cu) content is excessive, the possibility of causing star cracks on the surface of the steel sheet increases, and since copper (Cu) is an expensive element, there may be a problem of significantly increasing the manufacturing cost. Therefore, the upper limit of the copper (Cu) content may be preferably 0.60%, and a more preferable upper limit of the copper (Cu) content may be 0.35%.

[0117] Nickel (Ni): 1.0~4.0%

[0118] Nickel (Ni) is an element that effectively contributes to improving impact toughness by increasing stacking faults at low temperatures and facilitating cross-slip of dislocations, and to improving hardenability and strength. For this effect, nickel (Ni) of 1.0% or more may be included. The preferred nickel (Ni) content may be 0.10% or more. On the other hand, if nickel (Ni) is added excessively, it may increase the manufacturing cost due to its high cost. Therefore, the upper limit of nickel (Ni) content is preferably 4.0%, and a more preferred upper limit of nickel (Ni) content may be 3.0%.

[0119] Calcium (Ca): 0.0005~0.0040%

[0120] When calcium (Ca) is added after deoxidation by aluminum (Al), it combines with sulfur (S) that forms MnS inclusions, thereby suppressing the formation of MnS, and at the same time, by forming spherical CaS, it has the effect of suppressing the occurrence of cracks due to hydrogen-induced cracking. In order to sufficiently form sulfur (S) contained as an impurity into CaS, it is preferable to contain 0.0005% or more of calcium (Ca). However, if the amount added is excessive, the calcium (Ca) remaining after forming CaS combines with oxygen (O) to produce coarse oxidative inclusions, which may be elongated and destroyed during rolling, thereby lowering the lamellar tearing properties. Therefore, the upper limit of the calcium (Ca) content is preferably 0.0040%.

[0121] Boron (B): 0.001~0.005%

[0122] The above B is an effective element for improving the hardenability of hot-rolled steel sheets by delaying the transformation of austenite into ferrite during cooling transformation. To achieve this effect, it is preferable to add 0.001% or more. However, if the content exceeds 0.005%, there is a problem that the strength increases excessively and the impact toughness decreases as a result. In addition, boron oxide may be formed, which may deteriorate the surface quality of the steel sheet.

[0123] Oxygen (O): 0.002% or less

[0124] The above O combines with Ca or Al in the molten steel to exist in the form of CaO or Al2O3 or CaO-Al2O3 complex oxidative inclusions. The oxidative inclusions generated at high temperatures are crushed during the rough rolling stage of the slab and are elongated in the rolling direction, and remain undissolved not only during the rolling / heat treatment of the base metal but also in the HAZ during welding. Since the ends of the crushed oxidative inclusions can act as initiating points of fracture when evaluating impact and CTOD due to the notch effect, it is recommended to remove them through flotation in the secondary refining process as much as possible, and therefore the oxygen concentration in the molten steel is preferably 20 ppm or less. The present invention proposes a secondary refining method capable of extremely controlling oxidative inclusions, and after applying this process, the total oxygen content can be more preferably 15 ppm or less, and most preferably 10 ppm or less.

[0125] In addition to the aforementioned components, the remaining iron and other unavoidable impurities may be included. However, since unintended impurities from raw materials or the surrounding environment may inevitably be mixed in during the normal manufacturing process, they cannot be completely excluded. Since these impurities are readily apparent to anyone skilled in the art, their full contents are not specifically mentioned in this specification.

[0126] In addition, the addition of additional effective ingredients other than the aforementioned ingredients is not completely ruled out.

[0127] The above composition may have a Ceq of 0.50 to 0.70, as defined by the following equation 1.

[0128] [Relationship 1]

[0129] Ceq = [C] + [Mn] / 6 + ([Cr] + [Mo] + [V]) / 5 + ([Ni] + [Cu]) / 15

[0130] In the above relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] represent the contents (weight%) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, respectively, and 0 is substituted when these components are not intentionally added.

[0131] When Ceq by the above relationship 1 is less than 0.50, the martensite fraction decreases, so the high yield strength value required in the present invention cannot be secured, and the bainite fraction increases excessively, so that the occupation ratio of grain boundary cementite in the heat affected zone (HAZ), especially in the coarse grained heat affected zone (CGHAZ), after post-weld heat treatment (PWHT) increases, so that the appropriate impact toughness and HAZ CTOD value required in the present invention cannot be secured. In addition, when the Ceq value exceeds 0.70, the hardness becomes excessively high and cleavage fracture may be induced, so that the crack propagation resistance decreases, making it difficult to secure appropriate impact toughness and CTOD characteristics in the heat affected zone.

[0132] The steel material of the present invention has a maximum size of CaO-Al2O3-X (X is at least one of Ti, Mg, and S) composite oxidation inclusions observed in the upper layer, which is the area from the surface to t / 4 (t: thickness) in the thickness direction from the surface of the steel material based on the cross-section, of 50㎛ or less, and a unit area of ​​1㎛. 2 The number of parties may be less than 50.

[0133] The above complex oxidation inclusions can be observed not only in steel slabs, but also in the base metal and heat affected zone (HAZ) of steel before welding, steel after welding, and steel after welding heat treatment.

[0134] The above complex oxidation inclusions, when formed during the steel slab manufacturing stage, do not change in shape and fraction during the subsequent product manufacturing stage, and may be crushed and elongated in the rolling direction during rough rolling. The crushed oxidation inclusions may act as a place where stress is concentrated and may act as a starting point when a fracture occurs. Therefore, if the amount or size of the oxidation inclusions formed during the steel slab manufacturing stage is large, the amount of the crushed oxidation inclusions after rough rolling increases, which becomes the main cause of the deterioration of impact toughness and CTOD quality. Therefore, the size of the complex oxidation inclusions in the upper layer of the product, which is the area from the surface to t / 4, exceeds 50㎛ or 1㎛. 2 If the number of sugars exceeds 50, it may be difficult to secure low-temperature impact toughness at -40℃ and HAZ CTOD quality at -35℃ or lower based on the parent material.

[0135] The steel of the present invention has a tempered martensite or tempered martensite-bainite mixed structure as its base structure. Specifically, the steel may contain tempered martensite at an area fraction of 50% or more, and the remainder may contain bainite and structures that inevitably occur.

[0136] If the matrix structure is a tempered martensite-bainite mixed structure, the area fraction of tempered martensite can be 50 to 80%.

[0137] If the base structure is composed of a single structure of more than 80% tempered martensite, the hardness of the cross-section of the parent material of the steel after PWHT can be 220 to 290 HB based on the Brinell standard.

[0138] The above-mentioned inevitably occurring organization may be martensite (MA), etc.

[0139] The packet size of the low-temperature structure of the above-mentioned tempered martensite or bainite may be 17㎛ or less.

[0140] The average austenite grain size of the above steel may be 30㎛ or less.

[0141] Fine VC or VCN precipitates with a diameter of 5 to 15 nm observed on the cross-section of the above steel are 1 ㎛ 2 It may contain more than 5 precipitates.

[0142] The above microstructural characteristics can be the same in the steel before welding, the base material after welding, and the base material of the steel after welding heat treatment.

[0143] The average prior austenite grain size of the above steel may be 30㎛ or less. Post-weld heat treatment (PWHT) is typically performed at a high temperature of 595℃ or higher to prevent weld cracking through stress relief after welding. At this time, dislocation density decreases and cross-sectional hardness may drop, so the soundness of the weld can be improved. However, during the process of maintaining it at high temperature for a long time, cementite formed at the bainite lath interface decomposes, causing long-range diffusion of carbon into the austenite grain boundary. Therefore, when the grain size increases, the grain boundary area that can be occupied by cementite decreases, so segregation increases, which may deteriorate the impact toughness. Therefore, it is necessary to control the prior austenite grain size under appropriate low-temperature pressure. Accordingly, the austenite particle size is preferably 30 ㎛ or less, and more preferably 25 ㎛ or less.

[0144] As the temperature and time of the above PWHT increase, coarse cementite becomes spheroidized at the prior austenite grain boundary, and some of it is formed as film-like cementite. When such cementite is formed at the grain boundary, the intergranular strength becomes lower than the shear strength, which ultimately becomes the main cause of intergranular fracture. Therefore, in order to prevent intergranular fracture after PWHT, an appropriate proportion of martensite, which only causes short-range diffusion, must be included. However, if the proportion of martensite is too high, the cross-sectional hardness becomes excessively high, and even if intergranular fracture occurs, the impact absorption energy value will be low, which will ultimately cause a decrease in toughness. Therefore, it is desirable that the steel microstructure be composed of a mixed structure of tempered martensite and bainite, and at this time, the fraction of tempered martensite is preferably 50 to 80% in area fraction.

[0145] However, if the base structure of the above steel microstructure is tempered martensite, i.e., if martensite exceeds 80% (including 100%), the tempering temperature must be further increased to maintain the cross-sectional hardness of the base material at 200 to 290 HB. If the cross-sectional hardness is out of the range, it is difficult to secure the low-temperature impact toughness and CTOD value required in the present invention.

[0146] The packet size of the low-temperature structure of the above-mentioned tempered martensite or bainite may be 17㎛ or less. The packet size is a group unit of laths having the same orientation relationship in a low-temperature transformation structure such as bainite or martensite, and can be understood as a domain boundary that can have resistance (a characteristic in which the crack bends at the boundary when the crack propagates) when a physical crack propagates. As an example of measuring the above-mentioned packet size, it can be measured as a grain boundary that can be confirmed when domains within an orientation relationship of 15 degrees or less are set to the same orientation using an EBSD (Electro-back-scattered diffraction) device. Since the above-mentioned bainite or martensite grows only within the grains of the prior austenite, it is smaller than the grain size of the prior austenite. When the above packet size exceeds 17㎛, it is difficult to secure sufficient impact toughness because it is difficult to secure sufficient crack propagation resistance, so it is preferable that it be 17㎛ or less.

[0147] Meanwhile, the strength is reduced during tempering and PWHT because the dislocation density of the steel is lowered. If fine precipitates are used to appropriately compensate for this, the strength deterioration phenomenon during the heat treatment process can be reduced. Normally, precipitates such as NbC and NbCN are widely used in normalizing materials or TMCP (Thermo-Mechanical Control Process) materials. However, the temperature at which these precipitates are formed is 800-900℃, and since they are formed by deformation, they are not suitable for QT (Quenching-Tempering) heat treatment steels. Therefore, it is easy to utilize vanadium carbides / carbonitrides such as VC and VCN, which are formed in the temperature range of 550-700℃. If the diameter of the above VC or VCN fine precipitates is less than 5 nm, the strength improvement effect may be minimal because the size is too small, and if it exceeds 15 nm, the strength improvement effect is also reduced because the size is too coarse. Therefore, the size of the fine precipitates should be 5 to 15 nm. Meanwhile, if the density of VC or VCN is 1 ㎛ per unit area, 2 If the number is less than 5, the aforementioned precipitation strengthening effect is also reduced, so 5 or more is preferable. The formation of the above precipitates is determined by the chemical potential and temperature determined by the composition of the steel. The above 5-15 nm sized precipitates may be formed in large quantities, but if the number increases, the precipitates aggregate with each other and increase in size, so the 5-15 nm sized precipitates are not formed indefinitely. Therefore, even if the upper limit of the number of the above fine precipitates is not limited, it is not necessarily unclear to those skilled in the art.

[0148] Meanwhile, the microstructure of the heat affected zone (HAZ), which is an example of the present invention, will be described in detail. The HAZ may be a HAZ after post-weld heat treatment (PWHT), and in particular, may be a coarse grained heat affected zone (CGHAZ).

[0149] The above heat affected zone (HAZ) of the weld may have a grain boundary occupancy ratio of cementite existing at the grain boundary of 10% or less. The grain boundary occupancy ratio is

[0150] ) can be understood as the length ratio of film-like cementite occupied in the grain boundary, and in three dimensions, it becomes the area ratio. As an example of measuring the above occupancy ratio, the specimen can be observed using SEM, etc. after etching with nital, and the grain boundary where general grain boundary and film-like cementite are formed can be distinguished and measured.

[0151] Meanwhile, the average size of the cementite may be 1㎛ or less.

[0152] After post-weld heat treatment (PWHT), the HAZ, especially the CGHAZ, microstructure is preferably tempered martensite with an area fraction of 50% or more. In addition to the tempered martensite, the remainder may be bainite and other structures. If the area fraction of tempered martensite is less than 50%, the occupancy ratio of film-like cementite at the grain boundaries increases after PWHT, and the grains coarsen, making it difficult to secure an appropriate CTOD value of 0.1 mm or more at -35℃ or lower. For the same reason, the grain boundary occupancy ratio of cementite present at the grain boundaries after PWHT is preferably 10% or less, and the average size of cementite present at the grain boundaries is preferably 1 ㎛ or less.

[0153] The average grain size of the old austenite in the heat-affected zone (HAZ) may be 200㎛ or less. The HAZ microstructure is determined by the microstructure of the base material and the influence of welding conditions such as the welding heat input, preheating temperature, and lamination temperature. In particular, since the CGHAZ is a region where grains become significantly coarser compared to the base material, the material may be degraded compared to the base material. The average grain size of the old austenite in the CGHAZ is preferably 200㎛ or less. If it exceeds this, the hardenability increases excessively, increasing the cross-sectional hardness. Therefore, the average grain size of the CGHAZ is preferably 200㎛ or less. There may be no difference in the average grain size of the old austenite before and after PWHT. Since the grains must be reversely transformed to the austenite single-phase region for the grains to change in size, the temperature must be increased to 800℃ or higher. However, since the PWHT temperature is around 500-600℃, there may be no change in the AGS itself.

[0154] The above heat affected zone (HAZ) of the weld may have a grain boundary occupancy ratio of cementite existing at the grain boundary of 10% or less. The grain boundary occupancy ratio is

[0155] ) can be understood as the length ratio of film-like cementite occupied in the grain boundary, and in three dimensions, it becomes the area ratio. As an example of measuring the above occupancy ratio, the specimen can be observed using SEM, etc. after etching with nital, and the grain boundary where general grain boundary and film-like cementite are formed can be distinguished and measured.

[0156] The steel of the present invention can secure the strength and low-temperature impact toughness of the parent material after PWHT. Specifically, it has a yield strength of 690 MPa or more, a tensile strength of 770 to 940 MPa, and when a Charpy V-Notch Test is performed at -40°C or lower, the low-temperature impact toughness absorption energy of the parent material can be 50 J or more.

[0157] Meanwhile, the heat-affected zone after PWHT, especially the CGHAZ, can secure excellent low-temperature impact toughness and CTOD properties. Specifically, the low-temperature impact toughness absorption energy can be 40 J or more in a Charpy V-Notch Test performed at -40°C or lower, and the CTOD properties can be 0.1 mm or more at -35°C or lower.

[0158] Below, an example of a method for manufacturing steel according to the present invention is described in detail.

[0159] The above manufacturing method can manufacture steel slabs that satisfy the alloy composition and Ceq range of equation 1 described above, and perform heating, hot rolling, reheating, quenching, and tempering on the steel slabs to manufacture steel materials. This can be called the so-called RQT method, as it performs the reheating-quenching-tempering process.

[0160] Welding and post-weld heat treatment (PWHT) can be performed on steel manufactured in this way.

[0161] Below, each process is described in detail.

[0162] Steel slab manufacturing

[0163] The alloy composition and Ceq defined in the relational expression 1 described above satisfy the range of 0.50 to 0.70, and the maximum size of the CaO-Al2O3-X (X is at least one of Ti, Mg, and S) composite oxidation inclusions observed in the upper layer, which is the region from the surface to t / 4 (t: thickness, unit mm) in the thickness direction from the steel slab surface, is 50㎛ or less, and the unit area is 1㎛. 2 Manufactures steel slabs with a number of 50 or less.

[0164] As an example of a method for manufacturing the above steel slab, a calcium (Ca) raw material may be added to the secondary refined molten steel and a clean bubbling process may be performed. As a specific example, after the secondary refining, a metal Ca wire may be added to the molten steel, and the amount of Ca may be added in an amount of 0.015 to 0.15 kg / ton. The metal Ca wire may be composed of a steel material wrapping a Ca alloy, and the thickness of the steel material may be 1.2 to 1.4 mm, and the insertion speed of the Ca wire may be 100 to 300 meters / min.

[0165] After the above Ca injection, clean bubbling can be performed so that the inert gas injection amount in the ladle becomes 10 to 50 liters / min. There may be two inert gas injection points in the ladle, and the clean bubbling time can be 5 to 40 min.

[0166] In the present invention, there is no particular limitation on the process before secondary refining, and any conventional method can be applied. In the present invention, there is no particular limitation on the process after Ca injection and clean bubbling process, and the molten steel can be cooled under conventional conditions to manufacture a slab. In addition, as the amount of Al2O3 in the molten steel increases, the creation and coarsening of inclusions containing both Ca and Al progress, which increases the number of crushed inclusions during rolling, making it impossible to secure the low-temperature impact toughness and CTOD quality of the HAZ zone. Therefore, the total amount of inclusions in the molten steel before Ca injection can be limited to 2 to 5 ppm.

[0167] Hereinafter, the Ca injection process and clean bubbling process are described in detail.

[0168] (1) Calcium (Ca) injection

[0169] When manufacturing a steel slab, after secondary refining, metal Ca wire can be injected into the molten steel so that the amount of Ca injected is 0.015 to 0.15 kg / ton. The metal Ca wire is composed of steel wrapped around a Ca alloy, and the thickness of the steel can be 1.2 to 1.4 mm, and the injection speed of the Ca wire can be 100 to 300 meters / min.

[0170] When Ca is injected, if the thickness of the steel covering the Ca alloy of the Metal Ca Wire is less than 1.2 mm, the Ca melts at the top of the ladle, reducing the effect of the ferrostatic pressure, resulting in a poor Ca yield and an increase in the injected amount. On the other hand, if the thickness exceeds 1.4 mm, the Ca wire may contact the bottom of the ladle, causing the ladle's refractory to be melted, making it difficult to ensure operational stability.

[0171] When injecting Ca into the molten steel, the injection speed of the wire into the molten steel must be controlled in order to secure the Ca yield along with the thickness of the Metal Ca Wire. If the wire injection speed is less than 100 meter / min, the Ca will melt at the top of the ladle, reducing the effect of the ferrostatic pressure, resulting in a poor Ca yield and an increase in the injection amount. On the other hand, if the wire injection speed exceeds 300 meter / min, the Ca wire will contact the bottom of the ladle, causing a problem of melting of the ladle's refractory material, which makes it difficult to secure the stability of the operation. In addition, the preferable injection speed may be 120 to 300 meter / min, and more preferably, it may be 140 to 180 meter / min.

[0172] If the amount of Ca input is too small, MnS is generated in the center during solidification, which makes it difficult to secure low-temperature impact toughness due to segregation defects. If the amount is excessive, there is a problem in that it reacts with the Al2O3 of the refractory, accelerating the dissolution of the refractory, making it difficult to secure operational stability. Therefore, considering the above-mentioned problems, the amount of Ca input can be 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.

[0173] (2) Clean bubbling

[0174] After the above Ca injection, clean bubbling can be performed for 5 to 40 minutes so that the amount of inert gas injected into the ladle becomes 10 to 50 liters / min. There may be two inert gas injection points in the ladle.

[0175] If the amount of inert gas injected into the ladle is too small, the amount of Al2O3 clusters attached to and removed by the inert gas and complex inclusions containing both Ca and Al is small, resulting in poor cleanliness and the inability to secure impact toughness in the HAZ zone. On the other hand, if the amount is excessive, the stirring force becomes strong, causing slag to be mixed in at the same time as the molten steel surface to become bare, resulting in poor cleanliness. Therefore, the inert gas injection amount may be 10 to 50 liter / min, preferably 15 to 40 liter / min, and more preferably 20 to 30 liter / min.

[0176] Meanwhile, if there is only one inert gas injection point in the ladle, an uneven area exists in the molten steel, resulting in poor removal of Al2O3Cluster and complex inclusions containing both Ca and Al. In addition, if there are three or more, the stirring force becomes strong due to the occurrence of overlapping areas when the gas is injected, causing slag to be mixed in at the same time as the molten steel surface to become bare, resulting in poor cleanliness.

[0177] If the time for applying clean bubbling is too short even when the amount of inert gas injected into the ladle is limited, the amount of Al2O3Clusters and complex inclusions containing both Ca and Al that are attached to and removed by the inert gas will decrease, resulting in poor cleanliness and making it impossible to secure toughness. On the other hand, if the time is excessive, the temperature drop within the molten steel will increase, and a temperature gradient will occur within the ladle, which may result in poor cleanliness. The clean bubbling time may be 5 to 40 min, preferably 7 to 17 min, and more preferably 10 to 15 min.

[0178] steel slab heating

[0179] The above steel slab can be heated to a temperature range of 1050 to 1300°C.

[0180] The above heating is preferably performed at a slab heating temperature of 1150°C or higher to maximize the austenite grain size by re-dissolving the carbonitride of Ti or Nb or the coarse TiNb(C,N) crystallites formed during casting and heating and maintaining the austenite above the recrystallization temperature after sizing rolling. On the other hand, if the temperature is excessively high, problems may occur due to oxide scale at high temperatures, and the low-temperature impact toughness may deteriorate due to excessive grain growth, so the upper limit can be 1300°C.

[0181] hot rolling

[0182] The above heated steel slab can be hot rolled at a finishing hot rolling temperature of 830 to 1050°C.

[0183] The above 830℃ corresponds to the range above the recrystallization temperature. The composite inclusions created during the refining process must accommodate the deformation caused by rolling as the strength of the steel sheet increases as the rolling temperature decreases. As a result, they may be crushed or segmented into small inclusions, or inclusions such as MnS may be elongated. Since such crushed or segmented inclusions or elongated inclusions directly cause the initiation and propagation of low-temperature impact toughness, it is preferable to perform the finishing hot rolling at 830℃ or higher, which is higher than the recrystallization temperature at which work hardening does not occur. However, if the temperature exceeds 1050℃, austenite grain growth continues to occur even after the end of rolling, which may deteriorate the impact toughness due to the increase in the ductile-brittle transition temperature (DBTT).

[0184] The thickness of the steel sheet after final hot rolling can be 5 to 100 mm, more preferably 5 to 70 mm, and even more preferably 5 to 50 mm.

[0185] The steel sheet manufactured through the above hot rolling process can be cooled to room temperature in the air. In principle, air cooling is performed after the hot rolling process is completed. However, accelerated cooling to an appropriate temperature may be performed after the rolling process is completed to suppress precipitate growth and shorten process time.

[0186] In the present invention, when manufacturing a slab, if the dissolved hydrogen content in the molten steel is 1.3 ppm or more, a step of multi-stage cooling from a temperature of 200°C or higher based on the surface temperature to room temperature after hot rolling before reheating heat treatment may be further included. When cooling in multi-stage stages, by releasing dissolved hydrogen in the steel, internal micro-cracks caused by hydrogen can be more effectively suppressed, ultimately improving the low-temperature impact toughness of the base metal and the weld heat-affected zone.

[0187] Reheating

[0188] After the above hot rolling, reheating can be performed by heating the steel plate to a temperature of 820 to 950°C and maintaining it for 10 to 40 minutes.

[0189] If the heating temperature is lower than 820℃ or the holding time is lower than 10 minutes during the above reheating, the carbides generated during cooling after rolling or the impurity elements segregated at grain boundaries may not be smoothly re-dissolved and may remain after cooling, which may be a major cause of the deterioration of low-temperature impact toughness. On the other hand, if the heating temperature exceeds 950℃ or the holding time exceeds 40 minutes, the old austenite grain size may grow excessively due to grain growth, which may be a major cause of the deterioration of toughness due to the increase in DBTT. Therefore, the reheating temperature is preferably 820 to 950℃, and more preferably 850 to 890℃.

[0190] Quenching & Tempering

[0191] Quenching and tempering can be performed on the above reheated steel plate.

[0192] The cooling rate during the above quenching can be said to be one of the most important process factors that determine the microstructure of the base material, and in the present invention, cooling can be performed at a cooling rate of 3°C / s to 80°C / s based on the temperature at the point of the steel thickness t / 4 (t: thickness, unit: mm) after reheating to 300°C or lower. If the cooling rate is terminated at 300°C or higher, the microstructure of the base material may include martensite (MA, Martensite-Austenite Constituent) due to incomplete cooling transformation, which does not dissolve in the subsequent tempering and PWHT processes, causing deterioration of the low-temperature impact toughness. Therefore, the cooling termination temperature is preferably 300°C or lower.

[0193] Meanwhile, when the cooling rate is less than 3℃ / s, the microstructure of the base material is composed of a mixed structure of ferrite and bainite or a single-phase structure of bainite, rather than a mixed structure of martensite and bainite or a single-phase structure of martensite required in the present invention. When the bainite structure becomes the main structure, the strength before PWHT and low-temperature impact toughness can be secured, but it is not easy to secure low-temperature impact toughness and CTOD quality due to the formation of grain boundary cementite after the PWHT described above. On the other hand, when the cooling rate exceeds 80℃ / s, the microstructure is composed of 100% martensite, and the cross-sectional hardness becomes so high that it cannot be controlled by the tempering temperature alone, and there may be a risk of plate fracture due to the difference in cooling rate between the surface and the center.

[0194] 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.

[0195] The above quenched steel plate can be subjected to tempering heat treatment by heating it to a temperature of 550 to 700°C and maintaining it for 5 to 60 minutes. If the tempering temperature is lower than 550°C or the maintaining time is lower than 5 minutes, dislocation recovery does not occur properly during quenching, so the strength of the base material becomes excessively high, making it difficult to secure sufficient low-temperature impact toughness. In addition, if the temperature exceeds 700°C or the maintaining time exceeds 60 minutes, the grain boundaries (cementite) in the heat-affected zone of the weld may become coarsened, making it difficult to guarantee the toughness after PWHT.

[0196] Meanwhile, in addition to the above RQT, it can also be manufactured by a method of direct quenching and tempering after hot rolling (so-called DQT).

[0197] Welding can be performed on the above-manufactured steel material, and post-weld heat treatment can be performed.

[0198] welding

[0199] The steel of the present invention can be used by performing post-weld heat treatment (PWHT) after welding, and the steel of the present invention can secure strength and low-temperature impact toughness for the base material after the post-weld heat treatment, and can secure low-temperature impact toughness and low-temperature CTOD characteristics of the heat-affected zone (HAZ).

[0200] When welding steel, there are three parts: the unaffected base metal, the molten part (weld metal) that is melted and joined, and the heat-affected zone (HAZ), which is the part of the base metal affected by the welding heat. The molten part and the heat-affected zone can be collectively referred to as the weld zone. Meanwhile, the heat-affected zone can be classified into the coarse-grain HAZ (CGHAZ), subcritical HAZ (SCHAZ), intercritical HAZ (ICHAZ), and fine-grain HAZ (FGHAZ) depending on the maximum heating temperature and histological characteristics. Among them, the CGHAZ is the part where the old austenite grain size is the largest in the heat-affected structure (HAZ) during welding, and is maintained at a high temperature of 1000℃ or higher and then cooled. When the grain size is large, the ductile-brittle transition temperature increases, and since the hardenability increases, the brittle fracture resistance decreases when a hard structure is created. Therefore, the characteristics of the CGHAZ are important because it is the part of the HAZ where the toughness values ​​such as low-temperature impact toughness and CTOD can be the most deteriorated.

[0201] The above welding can be performed by any method that can be performed in the technical field to which the present invention pertains, and is not necessarily limited. In the present invention, the thermal cycle of the welding can be performed by heating up to a maximum temperature of 1250 to 1400°C at a heating rate of up to 50°C / s, maintaining the target temperature for 5 to 10 seconds, performing a first cooling at a cooling rate of 10 to 20°C / s to 500°C, and then performing a second cooling at a cooling rate of 5 to 15°C / s to 200°C. The thermal cycle related to the above maximum temperature and cooling rate can be set as described above based on the thermal history measured at the CGHAZ location with a thermocouple when performing actual welding.

[0202] The above heat history of welding is implemented as identically as possible through the heat history temperature measurement of the heat affected zone (HAZ) during actual FCAW, SAW, etc. welding. As a practical example of the above heat history, in the case of the FCAW (Flux Cored Arc Weld) method, the heat input can be 1.0 to 2.0 kJ / mm, and in the case of the SAW (Submerged Arc Weld) method, the heat input can be 2.5 to 4.0 kJ / mm. In this case, the preheating temperature and interpass temperature can both be 150 to 250℃.

[0203] Post-weld heat treatment (PWHT)

[0204] The above post-weld heat treatment (PWHT) can be performed by any method that can be performed in the technical field to which the present invention pertains, and is not particularly limited. For example, the PWHT can be performed by maintaining the temperature at 595°C to 625°C for 180 to 540 minutes (min) in accordance with ASME Sec. VIII. Meanwhile, it is preferable that the combined LMP in the tempering heat treatment and PWHT heat treatment according to the following equations 2 and 3 be 17.0 to 19.5.

[0205] 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.

[0206] [Relationship 2]

[0207] LMP = T k (Log t + 20)

[0208] Here, T k = Kelvin temperature of tempering and PWHT, t = tempering and PWHT time (hour)

[0209] [Relationship 3]

[0210] T(logt eq +20) = T i (logt i+20)

[0211] Total LMP = T i [log(t i +t eq )+20]

[0212] Here, T i : Kelvin temperature of PWHT, t i : PWHT time (hour), T: Tempering temperature in Kelvin, t eq : Equivalent tempering time (hour) when converted to PWHT temperature

[0213] If the LMP is less than 17.0, as previously explained, the dislocation density does not decrease, which may result in strength overrun failure, or the cross-sectional hardness may be too high, resulting in reduced impact toughness. On the other hand, if the LMP exceeds 19.5, the grain boundary occupancy ratio of cementite after PWHT is too high, making it difficult to adequately secure impact toughness and CTOD quality. Therefore, the LMP is preferably between 17.0 and 19.5, and more preferably between 18.0 and 19.3.

[0214] Hereinafter, embodiments of the present invention will be described. It should be apparent to those skilled in the art that various modifications to the following embodiments may be made without departing from the scope of the present invention. The following embodiments are intended to facilitate understanding of the present invention, and the scope of the present invention should not be limited to the following embodiments, but should be determined not only by the claims set forth below but also by their equivalents.

[0215] (Example 1)

[0216] A steel slab was manufactured using molten steel having the alloy composition of Table 1 above (the content unit of the component elements in Table 1 is weight%, but the unit of P, S, and Ca is ppm. And the remaining components are Fe and unavoidable impurities) and Ceq of the above-mentioned relational expression 1 under the conditions of Table 2, and heating and hot rolling were performed to manufacture a hot-rolled steel sheet having a thickness of 50 mm.

[0217] The hot-rolled steel sheet manufactured above was reheated and quenched under the conditions disclosed in Table 3, and then tempered at 610°C for 20 minutes.

[0218] Meanwhile, in order to confirm the excellent strength and low-temperature impact toughness of the steel even after post-weld heat treatment, a heat treatment process simulating post-weld heat treatment was performed. In this heat treatment process, the hot-rolled steel sheet subjected to tempering was heat treated at 595°C for 180 minutes for 3 cycles.

[0219] The microstructure and physical properties of the parent material were measured for each specimen manufactured as described above, and the results are shown in Tables 4 and 5 below.

[0220] Here, the austenite grain size and the fraction of each phase were measured using electron back-scattered diffraction (EBSD) on the collected specimens. Precipitates and inclusions were measured using transmission electron microscopy (TEM). Precipitates and inclusions were confirmed through diffraction patterns and EDX mapping.

[0221] Yield / tensile strength was evaluated through room-temperature tensile tests, with a 0.2% offset applied to yield strength. Impact toughness for each specimen was determined using the average of absorbed energy values ​​measured three times at the corresponding temperature using a Charpy V-notch test. Brinell hardness was tested according to ISO 6506 Part 1.

[0222]

[0223]

[0224]

[0225]

[0226]

[0227] As can be seen from Tables 1 to 5 above, in the case of specimen numbers 1 to 5 and 17 to 20 that satisfy the alloy composition and manufacturing conditions proposed by the present invention, even after excellent post-weld heat treatment, excellent quality can be secured, such as strength of the base material and excellent low-temperature impact toughness at -40°C.

[0228] In contrast, specimens No. 6 to 16 satisfy the alloy composition proposed by the present invention but do not satisfy the manufacturing conditions, making it difficult to secure the characteristics of the base material proposed by the present invention. In addition, specimens No. 21 to 25 satisfy the steel manufacturing conditions proposed by the present invention but do not satisfy the alloy composition, making it difficult to secure the characteristics of the base material proposed by the present invention.

[0229] (Example 2)

[0230] For steel manufactured under the conditions of Table 3 above, welding and post-weld heat treatment were performed to confirm the low-temperature toughness and CTOD characteristics of the CGHAZ among the weld heat-affected zone (HAZ).

[0231] To this end, specifically, welding simulations were performed on steels manufactured according to Tables 1 to 3 under the welding heat history conditions of Table 6, and a post-weld heat treatment process was performed at 595°C for 180 minutes in 3 cycles. In Table 6, LMP was derived by the aforementioned relational expression 3.

[0232] The microstructure and physical properties of CGHAZ were measured in each specimen manufactured as described above, and the results are shown in Tables 7 and 8 below.

[0233] Here, the phase fraction was measured by collecting the specimen and using EBSD (Electron Back Scattered Diffraction).

[0234] The impact toughness of each specimen was determined by the average of absorbed energy values ​​measured three times at the corresponding temperature using the Charpy V-notch Test. Brinell hardness was tested according to ISO 6506 Part 1, and CTOD testing was performed according to BS5762 (British Standard).

[0235] In addition, the cementite grain boundary occupancy ratio was measured using transmission electron microscopy (TEM). The cementite grain boundary occupancy ratio was measured as the ratio of the total grain boundary length occupied by film-like cementite.

[0236]

[0237]

[0238]

[0239] As can be seen from Tables 6 to 8 above, in the case of specimen numbers 1 to 5 and 11 to 16 that satisfy the alloy composition and manufacturing conditions proposed in the present invention, it can be seen that not only excellent low-temperature impact toughness but also good CTOD quality can be secured in CGHAZ.

[0240] In contrast, specimens No. 6 to 10 satisfy the alloy composition proposed by the present invention but do not satisfy the conditions for manufacturing steel, and in particular, the size of the complex oxide inclusions is coarse, which can be seen to deteriorate the properties of the CGHAZ. On the other hand, specimens No. 17 to 20 satisfy the alloy composition proposed by the present invention and the conditions for manufacturing steel, but do not satisfy the welding and post-weld heat treatment conditions, and thus fail to secure the properties of the CGHAZ.

[0241] It can be seen that specimens No. 21 to 25 satisfy the manufacturing conditions proposed by the present invention but do not satisfy the alloy composition, and thus do not satisfy the appropriate quality required by the present invention.

[0242] As described above, the detailed description of the present invention has described preferred embodiments of the present invention. However, it will be apparent to those skilled in the art that various modifications may be made without departing from the scope of the present invention. Therefore, the scope of the rights of the present invention should not be limited to the described embodiments, but should be determined not only by the claims described below but also by equivalents thereof.

Claims

1. In weight%, C: 0.05 to 0.20%, Si: 0.05 to 0.50%, Mn: 0.5 to 2.0%, Al: 0.005 to 0.10%, P: 0.010% or less, S: 0.0050% or less, Nb: 0.001 to 0.070%, V: 0.001 to 0.30%, Ti: 0.001 to 0.030%, Cr: 0.01 to 1.0%, Mo: 0.01 to 1.0%, Cu: 0.01 to 0.60%, Ni: 1.0 to 4.0%, B: 0.001 to 0.005%, Ca: 0.0005 to 0.0040%, O: 0.002% or less, including the remainder of Fe and other unavoidable impurities. Ceq defined in the following relational expression 1 satisfies the range of 0.50 to 0.70, The maximum size of the CaO-Al2O3-X (X is at least one of Ti, Mg, and S) composite oxidizing inclusions observed in the upper layer, which is the area from the surface to the surface layer in the thickness direction to t / 4 (t is the thickness, unit: mm), is 50㎛ or less, and the unit area is 1㎛. 2 The number of sugars is 50 or less, The grain boundary occupancy ratio of cementite existing in the grain boundary is 10% or less, and the average size of cementite is 1㎛ or less. Welded heat affected zone including cross-sectional hardness of 200 to 300 HB. [Relationship 1] Ceq = [C] + [Mn] / 6 + ([Cr] + [Mo] + [V]) / 5 + ([Ni] + [Cu]) / 15 In the above relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] represent the contents (in weight%) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, respectively, and if these components are not intentionally added, 0 is substituted.

2. In claim 1, The above weld heat affected zone is a coarse grained heat affected zone (CGHAZ).

3. In claim 1, The microstructure of the above weld heat affected zone is a weld heat affected zone in which tempered martensite has an area fraction of 50% or more and the remainder contains bainite.

4. In claim 1, A weld heat affected zone including an average grain size of austenite in the weld heat affected zone of 200㎛ or less.

5. In claim 1, The above welding heat affected zone is a welding heat affected zone having a Charpy impact absorption energy of 40J or more at -40°C and a CTOD characteristic of 0.1mm or more at -35°C.

6. In weight%, C: 0.05 to 0.20%, Si: 0.05 to 0.50%, Mn: 0.5 to 2.0%, Al: 0.005 to 0.10%, P: 0.010% or less, S: 0.0050% or less, Nb: 0.001 to 0.070%, V: 0.001 to 0.30%, Ti: 0.001 to 0.030%, Cr: 0.01 to 1.0%, Mo: 0.01 to 1.0%, Cu: 0.01 to 0.60%, Ni: 1.0 to 4.0%, B: 0.001 to 0.005%, Ca: 0.0005 to 0.0040%, O: 0.002% or less, containing the remainder of Fe and other unavoidable impurities, and the following relationship: The maximum size of the composite oxidizing inclusions of CaO-Al2O3-X (X is at least one of Ti, Mg, and S) observed in the upper layer, which is the area from the surface to the surface layer to t / 4 (t is the thickness, unit: mm) in the thickness direction, is 50㎛ or less, and the unit area is 1㎛. 2 A step for preparing steel having a number of sugars of 50 or less; a step of welding the above steel; and Step of performing post-welding heat treatment on the above welded steel A method for manufacturing a heat affected zone including a welding part. [Relationship 1] Ceq = [C] + [Mn] / 6 + ([Cr] + [Mo] + [V]) / 5 + ([Ni] + [Cu]) / 15 In the above relational expression 1, [C], [Mn], [Cr], [Mo], [V], [Ni] and [Cu] represent the contents (in weight%) of C, Mn, Cr, Mo, V, Ni and Cu contained in the steel, respectively, and if these components are not intentionally added, 0 is substituted.

7. In claim 6, The above welding thermal cycle is a method for manufacturing a heat-affected zone, which includes heating up to a maximum temperature of 1250 to 1400°C at a heating rate of up to 50°C / s, maintaining the temperature at the target temperature for 5 to 10 seconds, performing a first cooling at a cooling rate of 10 to 20°C / s to 500°C, and then performing a second cooling at a cooling rate of 5 to 15°C / s to 200°C.

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

9. In claim 6, The above post-weld heat treatment is a method for manufacturing a heat-affected zone by maintaining the temperature at 595°C to 625°C for 180 to 540 minutes.

10. In claim 6, The above steel material, As a microstructure, tempered martensite contains 50% or more (including 100%) in area fraction, and the remainder contains bainite and other inevitable structures. A method for manufacturing a heat-affected zone of a weld, comprising: a packet size of the above tempered martensite and bainite being 17 ㎛ or less.

11. In claim 10, A method for manufacturing a welded heat affected zone, wherein the steel has a steel cross-section Brinell hardness of 220 to 290 HB when the steel contains more than 80% tempered martensite.

12. In claim 10, Among the precipitates observed in the steel cross section, the above steel has VC or VCN precipitates with a diameter of 5 to 15 nm and a diameter of 1 ㎛. 2 A method for manufacturing a heat affected zone of a weld comprising five or more compounds.

13. In claim 10, A method for manufacturing a heat-affected zone of a weld, wherein the steel includes an average austenite grain size of 30㎛ or less.

14. In claim 6, The above steel material, Steps for manufacturing a steel slab; A step of reheating the above steel slab to a temperature range of 1050 to 1300℃; A step of manufacturing steel by hot rolling the above steel slab at a finishing hot rolling temperature of 830 to 1050℃; A step of heating the above steel to a temperature range of 820 to 950°C and maintaining it for 10 to 40 minutes, and then cooling it to a temperature of 300°C or lower at a cooling rate of 3°C / s to 80°C / s based on the steel thickness t / 4 temperature; and A method for manufacturing a welded heat affected zone, the method comprising a tempering heat treatment step of heating the steel to 550 to 700°C and maintaining it for 5 to 60 minutes after the above cooling.

15. In claim 14, The steps for manufacturing the above steel slab are: A step of adding Metal Ca Wire to the molten steel after secondary refining so that the amount of Ca added is 0.015 to 0.15 kg / ton; and After the above Ca injection, a step of clean bubbling is included for 5 to 40 minutes so that the amount of inert gas injected into the ladle becomes 10 to 50 liter / min. A method for manufacturing a heat-affected zone of a weld, wherein the feeding speed of the Ca Wire is 100 to 300 meters / min and the number of inert gas injection points in the ladle is two.

16. In claim 14, A method for manufacturing a heat-affected zone of a weld, wherein the LMP according to the following relational expression 2 is 17.0 to 19.5 when the temperatures of the above tempering heat treatment and the post-weld heat treatment are the same. [Relationship 2] LMP = T k (Log t + 20) Here, T k = Kelvin temperature of tempering and PWHT, t = time of tempering and PWHT (hour) 17. In claim 14, A method for manufacturing a heat-affected zone of a weld, wherein the LMP according to the following relational expression 3 is 17.0 to 19.5 when the temperatures of the above tempering heat treatment and the post-weld heat treatment are different. [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 (hour) converted to PWHT temperature