Classifications

• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0226—Hot rolling
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
  • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
  • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/002—Bainite
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/005—Ferrite

Definitions

• the present disclosure relates to a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance, and to a method of manufacturing the same.
• microstructures of thick steel plates may be coarse, so that low temperature properties on which grain sizes have the most significant effect may be degraded.
• Such technologies may contribute to refining a structure of a surface portion, but may not solve a problem in which impact toughness is degraded due to coarsening of structures other than the surface portion. Thus, such technologies may not be fundamental countermeasures to brittle crack arrestability.
• the microstructure in a HAZ includes low temperature transformation ferrite having high strength, such as bainite, there is a limitation in which HAZ properties, in detail, toughness, is significantly reduced.
• martensite-austenite may be transformed to have a different phase through tempering, or the like, to secure physical properties.
• HAZ in which an effect of tempering disappears due to thermal history, it is impossible to apply brittle crack initiation resistance.
• An aspect of the present disclosure may provide a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance.
• Another aspect of the present disclosure may provide a method of manufacturing a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance.
• a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises, by wt %, carbon (C): 0.05% to 0.09%, manganese (Mn): 1.5% to 2.2%, nickel (Ni): 0.3% to 1.2%, niobium (Nb): 0.005% to 0.04%, titanium (Ti): 0.005% to 0.04%, copper (Cu): 0.1% to 0.8%, silicon (Si): 0.05% to 0.3%, aluminum (Al): 0.005% to 0.05%, phosphorus (P): 100 ppm or less, sulfur (S): 40 ppm or less, iron (Fe) as a residual component thereof, and inevitable impurities, wherein a microstructure of a central portion includes, by area %, a mixed phase of acicular ferrite and granular bainite in an amount of 70% or greater, upper bainite in an amount of 20% or less,
• Contents of Cu and Ni may be set such that a weight ratio of Cu to Ni may be 0.8 or less, and in more detail, 0.6 or less.
• the high-strength steel material may have yield strength of 460 MPa or greater.
• the high-strength steel material may have a Charpy fracture transition temperature of ⁇ 40° C. or lower in a 1 ⁇ 2t position in a steel material thickness direction, where t is a steel sheet thickness.
• a method of manufacturing a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises rough rolling a slab at a temperature of 900° C. to 1100° C. after reheating the slab at 1000° C.
• a reduction ratio per pass of three final passes during the rough rolling a slab may be 5% or greater, and a total cumulative reduction ratio may be 40% or greater.
• a strain rate of three final passes during the rough rolling a slab may be 2/sec or lower.
• a grain size of a central portion in a bar thickness direction before finish rolling after the rough rolling a slab may be 150 ⁇ m or less, in detail, 100 ⁇ m or less, and more specifically, 80 ⁇ m or less.
• a reduction ratio during the finish rolling may be set such that a ratio of a slab thickness (mm) to a steel sheet thickness (mm) after the finish rolling may be 3.5 or greater, and in more detail, 4 or greater.
• a cumulative reduction ratio during the finish rolling may be maintained to be 40% or greater, while the reduction ratio per pass, not including skin pass rolling, may be maintained to be 4% or greater.
• Skin pass rolling refers to a process of rolling a sheet at a relatively low reduction ratio in order to secure flatness of the sheet.
• the cooling the steel sheet may be performed at a cooling rate of the central portion of 2° C./s or higher.
• the cooling the steel sheet may be performed at an average cooling rate of 3° C./s to 300° C./s.
• a high-strength steel material having a relatively high level of yield strength, as well as excellent brittle crack arrestability and welding zone brittle crack initiation resistance.
• the inventors of the present disclosure conducted research and experiments to improve yield strength, brittle crack arrestability, and welding zone brittle crack initiation resistance of a thick steel material and proposed the present disclosure based on results thereof.
• a steel composition, a structure, and manufacturing conditions of a steel material may be controlled, thereby improving yield strength, brittle crack arrestability, and welding zone brittle crack initiation resistance of the thick steel material.
• a main concept of an exemplary embodiment is as follows.
• the steel composition is appropriately controlled to improve strength through solid solution strengthening.
• contents of manganese (Mn), nickel (Ni), copper (Cu), and silicon (Si) are optimized for solid solution strengthening.
• the steel composition is appropriately controlled to improve strength by increasing hardenability.
• the contents of Mn, Ni, and Cu, as well as a carbon (C) content are optimized to increase hardenability.
• a fine structure is secured in a central portion of the thick steel material even at a relatively slow cooling rate.
• a composition is appropriately controlled to control a fraction of martensite-austenite.
• contents of C, Si, and niobium (Nb), affecting generation of martensite-austenite, are optimized.
• the steel composition may be optimized, thereby securing excellent brittle crack initiation resistance even in a heat affected zone (HAZ).
• HAZ heat affected zone
• a structure of the steel material may be controlled to improve strength and brittle crack arrestability.
• a structure of the central portion and a surface layer region is controlled in a direction of a steel material thickness.
• a microstructure may be controlled, thereby securing strength required in the steel material, while the microstructure facilitating generation of a crack may be excluded, thereby improving brittle crack arrestability.
• rough rolling conditions may be controlled to refine the structure of the steel material.
• the fine structure is secured in the central portion by controlling a rolling condition during rough rolling. Using a process described above, the generation of acicular ferrite and granular bainite is facilitated.
• a finish rolling condition is controlled to further refine the structure of the steel material.
• a finish rolling temperature and rolling conditions may be controlled to generate a relatively large amount of strain bands in austenite during finish rolling and secure a large number of ferrite nucleation sites, thereby securing a fine structure in the central portion of the steel material. As such, generation of acicular ferrite and granular bainite is facilitated.
• the high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises, by wt %, carbon (C): 0.05% to 0.09%, manganese (Mn): 1.5% to 2.2%, nickel (Ni): 0.3% to 1.2%, niobium (Nb): 0.005% to 0.04%, titanium (Ti): 0.005% to 0.04%, copper (Cu): 0.1% to 0.8%, silicon (Si): 0.05% to 0.3%, aluminum (Al): 0.005% to 0.05%, phosphorus (P): 100 ppm or less, sulfur (S): 40 ppm or less, iron (Fe) as a residual component thereof, and inevitable impurities, wherein a microstructure of a central portion includes, by area %, a mixed phase of acicular ferrite and granular bainite in an amount of 70% or greater, upper bainite in an amount of 20% or less, and
• C Since C is the most significant element used in securing basic strength, C is required to be contained in steel within an appropriate range. In order to obtain an effect of addition, C may be added in an amount of 0.05% or greater.
• the C content may be limited to 0.055% to 0.08%, and more specifically, to 0.06% to 0.075%.
• Mn is a useful element improving strength through solid solution strengthening and increasing hardenability to generate low temperature transformation ferrite.
• Mn since Mn may generate low temperature transformation ferrite even at a relatively low cooling rate due to improved hardenability, Mn is a main element to secure strength of a central portion of a thick steel plate.
• Mn may be added in an amount of 1.5% or greater.
• the Mn content may be limited to 1.5% to 2.2%.
• the Mn content may be limited to 1.6% to 2.0%, and more specifically, to 1.65% to 1.95%.
• Ni is a significant element used in improving impact toughness by facilitating a dislocation cross slip at a relatively low temperature and increasing strength by improving hardenability.
• Ni may be added in an amount of 0.3% or greater.
• hardenability is excessively increased to generate low temperature transformation ferrite, thereby degrading toughness, and a manufacturing cost may be increased due to a relatively high cost of Ni, as compared with other hardenability elements.
• an upper limit value of the Ni content maybe limited to 1.2%.
• the Ni content may be limited to 0.4% to 1.0%, and more specifically, to 0.45% to 0.9%.
• Nb is educed to have a form of NbC or NbCN to improve strength of a base material.
• Nb solidified when being reheated at a relatively high temperature is significantly finely educed to have the form of NbC during rolling to suppress recrystallization of austenite, thereby having an effect of refining a structure.
• Nb may be added in an amount of 0.005% or greater.
• generation of martensite-austenite in the HAZ may be facilitated to degrade brittle crack initiation resistance and cause a brittle crack in an edge of the steel material.
• an upper limit value of an Nb content maybe limited to 0.04%.
• the Nb content may be limited to 0.01% to 0.035%, and more specifically, to 0.015% to 0.03%.
• Ti is a component educed to be TiN when being reheated and inhibiting growth of the base material and a grain in the HAZ to greatly improve low temperature toughness.
• Ti may be added in an amount of 0.005% or greater.
• a Ti content may be limited to 0.005% to 0.04%.
• the Ti content may be limited to 0.008% to 0.03%, and more specifically, to 0.01% to 0.02%.
• Si is a substitutional element improving strength of the steel material through solid solution strengthening and having a strong deoxidation effect, so that Si may be an element essential in manufacturing clean steel.
• Si may be added in an amount of 0.05% or greater.
• a coarse martensite-austenite phase may be formed to degrade brittle crack arrestability and welding zone brittle crack initiation resistance.
• an upper limit value of an Si content may be limited to 0.3%.
• the Si content may be limited to 0.1% to 0.25%, and more specifically, to 0.1% to 0.2%.
• Cu is a main element used in improving hardenability and causing solid solution strengthening to enhance strength of the steel material.
• Cu is a main element used in increasing yield strength through the generation of an epsilon Cu precipitate when tempering is applied.
• Cu may be added in an amount of 0.1% or greater.
• an upper limit value of a Cu content may be limited to 0.8%.
• the Cu content may be limited to 0.2% to 0.6%, and more specifically, to 0.25% to 0.5%.
• Contents of Cu and Ni may be set such that the weight ratio of Cu to Ni may be 0.8 or less, and in more detail, 0.6 or less. More specifically, the weight ratio of Cu to Ni may be limited to 0.5 or less.
• Al is a component functioning as a deoxidizer.
• an inclusion may be formed to degrade toughness.
• an Al content may be limited to 0.005% to 0.05%.
• P and S are elements causing brittleness in a grain boundary or forming a coarse inclusion to cause brittleness.
• a P content may be limited to 100 ppm or less, while an S content may be limited to 40 ppm or less.
• a residual component of an exemplary embodiment is iron (Fe).
• a microstructure of a central portion includes, by area %, a mixed phase of acicular ferrite and granular bainite in an amount of 70% or greater, upper bainite in an amount of 20% or less, and one or more selected from a group consisting of ferrite, pearlite, and martensite-austenite (MA), as residual components; a circle-equivalent diameter of an effective grain of the upper bainite having a high angle grain boundary of 15° or greater measured using an electron backscatter diffraction (EBSD) method being 15 ⁇ m or less; a microstructure in a region at a depth of 2 mm or less, directly below a surface, includes, by area %, ferrite in an amount of 20% or greater and one or more of bainite and martensite as residual components; and a heat affected zone (HAZ) formed during welding includes, by area %, martensite-austenite (MA) in an amount of 5% or less.
• EBSD electron backscatter dif
• the fraction of the mixed phase of acicular ferrite and granular bainite may be 75% or greater, and more specifically, may be limited to 80% or greater.
• a fraction of acicular ferrite may be 20% to 70%.
• the fraction of acicular ferrite may be limited to 30% to 50%, and more specifically, to 30% to 40%.
• a fraction of granular bainite may be 10% to 60%.
• the fraction of granular bainite may be limited to 20% to 50%, and more specifically, to 30% to 50%.
• a microcrack may be generated in a front end of a crack during brittle crack propagation, thereby degrading brittle crack arrestability.
• the fraction of upper bainite in the central portion may be 20% or less.
• the fraction of upper bainite may be limited to 15% or less, and more specifically, to 10% or less.
• the circle-equivalent diameter of the effective grain of upper bainite in the central portion having a high angle grain boundary of 15° or greater measured using an EBSD method exceeds 15 ⁇ m, there is a problem in which a crack may be easily generated despite a relatively low fraction of upper bainite.
• the circle-equivalent diameter of the effective grain of upper bainite in the central portion may be 15 ⁇ m or less.
• the surface portion microstructure in the region at a depth of 2 mm or less, directly below the surface includes ferrite in an amount of 20% or greater, crack propagation may be effectively prevented on the surface during brittle crack propagation, thereby improving brittle crack arrestability.
• the fraction of ferrite may be limited to 30% or greater, and more specifically, to 40% or greater.
• Ferrite in the microstructure in the central portion and the surface portion refers to polygonal ferrite or elongated polygonal ferrite.
• the fraction of martensite-austenite in the HAZ may be 5% or less.
• Welding heat input during welding may be 0.5 kJ/mm to 10 kJ/mm.
• a welding method during welding is not specifically limited and may include, for example, flux cored arc welding (FCAW), submerged arc welding (SAW), and the like.
• FCAW flux cored arc welding
• SAW submerged arc welding
• the steel material may have yield strength of 460 MPa or greater.
• the steel material may have a Charpy fracture transition temperature of ⁇ 40° C. or lower in a 1 ⁇ 2t position in a steel material thickness direction, where t is a steel sheet thickness.
• the steel material have a thickness of 50 mm or greater, and in detail, a thickness of 50 mm to 100 mm.
• the method of manufacturing a high-strength steel material having excellent brittle crack arrestability and welding zone brittle crack initiation resistance comprises rough rolling a slab at a temperature of 900° C. to 1100° C. after reheating the slab at 1000° C.
• a slab is reheated before rough rolling.
• a reheating temperature of the slab may be 1000° C. or higher so that a carbonitride of Ti and/or Nb, formed during casting, may be solidified.
• an upper limit value of the reheating temperature may be 1100° C.
• a reheated slab is rough rolled.
• a rough rolling temperature may be a temperature Tnr at which recrystallization of austenite is halted, or higher. Due to rolling, a cast structure, such as a dendrite formed during casting, may be destroyed, and an effect of reducing a size of austenite may also be obtained. In order to obtain the effect, the rough rolling temperature may be limited to 900° C. to 1100° C.
• the rough rolling temperature may be 950° C. to 1050° C.
• a reduction ratio per pass of three final passes during rough rolling may be 5% or greater, and a total cumulative reduction ratio may be 40% or greater.
• the reduction ratio per pass may be 7% to 20%.
• the total cumulative reduction ratio may be 45% or greater.
• the reduction ratio per pass of the three final passes may be limited to 5% or greater.
• the total cumulative reduction ratio during rough rolling may be set to be 40% or greater.
• a strain rate of the three final passes during rough rolling may be 2/sec or lower.
• the strain rate may be limited to 2/sec or lower, thereby refining the grain size of the central portion.
• generation of acicular ferrite and granular bainite may be facilitated.
• a rough rolled bar may be finish rolled at a temperature of Ar 3 (a ferrite transformation initiation temperature)+60° C. to Ar 3 ° C. to obtain a steel sheet so that a further refined microstructure may be obtained.
• Ar 3 a ferrite transformation initiation temperature
• a relatively large amount of strain bands may be generated in austenite to secure a relatively large number of ferrite nucleation sites, thereby obtaining an effect of securing a fine structure in the central portion of a steel material.
• a cumulative reduction ratio during finish rolling may be maintained to be 40% or greater.
• the reduction ratio per pass, not including skin pass rolling, may be maintained to be 4% or greater.
• the cumulative reduction ratio may be 40% to 80%.
• the reduction ratio per pass may be 4.5% or greater.
• finish rolling temperature In a case in which a finish rolling temperature is reduced to Ar 3 or lower, coarse ferrite is generated before rolling and is elongated during rolling, thereby reducing impact toughness.
• finish rolling In a case in which finish rolling is performed at a temperature of Ar 3 +60° C. or higher, the grain size is not effectively refined, so that the finish rolling temperature during finish rolling may be set to be a temperature of Ar 3 +60° C. to Ar 3 ° C.,
• a reduction ratio in an unrecrystallized region may be limited to 40% to 80% during finish rolling.
• the grain size of the central portion of the bar in a thickness direction after rough rolling before finish rolling may be 150 ⁇ m or less, in detail, 100 ⁇ m or less, and more specifically, 80 ⁇ m or less.
• the grain size of the central portion of the bar in a thickness direction after rough rolling before finish rolling may be controlled depending on a rough rolling condition, or the like.
• the reduction ratio during finish rolling may be set such that a ratio of a slab thickness (mm) to a steel sheet thickness (mm) after finish rolling may be 3.5 or greater, and in detail, 4 or greater.
• an advantage of improving toughness of the central portion may be added by increasing yield strength/tensile strength, improving low temperature toughness, and decreasing the grain size of the central portion in the thickness direction through refinement of the final microstructure.
• the steel sheet may have a thickness of 50 mm or greater, and in detail, 50 mm to 100 mm.
• the steel sheet is cooled to a temperature of 500° C., or lower, after finish rolling.
• a microstructure may not be properly formed, so that sufficient yield strength may be difficult to secure. For example, yield strength of 460 MPa or greater may be difficult to secure.
• a generation amount of acicular ferrite and granular bainite may be reduced, and strength thereof may be reduced due to an auto-tempering effect.
• the cooling end temperature maybe 400° C. or lower.
• the steel sheet may be cooled at a cooling rate of the central portion of 2° C./s or higher.
• the cooling rate of the central portion of the steel sheet is lower than 2° C./s, the microstructure may not be properly formed, so that it may be difficult to secure sufficient yield strength. For example, yield strength of 460 MPa or greater may be difficult.
• the steel sheet may be cooled at an average cooling rate of 3° C./s to 300° C./s.
• a thickness of a bar having been rough rolled was 192 mm, while a grain size of a central portion after rough rolling before finish rolling, as illustrated in Table 2, was 66 ⁇ m to 82 ⁇ m.
• a reduction ratio of three final passes during rough rolling was within a range of 7.9% to 14.1%.
• a strain rate during rolling was within a range of 1.22/s to 1.68/s.
• finish rolling was performed at the reduction ratio per pass of 4.2% to 5.6% and at the cumulative reduction ratio of 50% at a temperature equal to a difference between a finish rolling temperature and an Ar a temperature, illustrated in Table 2 below to obtain a steel sheet having a thickness illustrated in Table 3 below, and then the steel sheet was cooled to a temperature of 241° C. to 378° C. at a cooling rate of the central portion of 3.8° C./sec to 5.0° C./sec.
• the Kca value in Table 4 below is a value evaluated by performing an ESSO test on the steel sheet.
• the CTOD value was a result in which a FCAW (1.0 kJ/mm) welding process is performed to carry out structure analysis and a CTOD test on the HAZ.
• the difference between the finish rolling temperature during finish rolling and the Ar 3 temperature was controlled to be 60° C. or higher. Rolling was performed at a relatively high temperature, so that sufficient reduction was not applied to the central portion. In addition, cooling was started at a relatively high temperature, so that ferrite of 20% or greater was not generated in a surface portion. Thus, it can be confirmed that the Kca value measured at a temperature of ⁇ 10° C. may not exceed 6000 required in a steel material for shipbuilding of the related art.
• a C content had a value higher than an upper limit value of a C content of an exemplary embodiment. It can be confirmed that a relatively large amount of coarse upper bainite was generated in the central portion during rough rolling, so the Kca value measured at a temperature of ⁇ 10° C. was 6000 or less. It can be confirmed that a relatively large amount of martensite-austenite (MA) was also generated in the HAZ, so the CTOD value was 0.25 mm or less.
• MA martensite-austenite
• a Si content had a value higher than an upper limit value of a Si content of an exemplary embodiment. It can be confirmed that a relatively large amount of Si was added to generate a relatively large amount of an MA structure in the HAZ, so the CTOD value is 0.25 mm or less.
• a Mn content has a value higher than an upper limit value of a Mn content of an exemplary embodiment. It can be confirmed that due to having a relatively high level of hardenability, a relatively large amount of upper bainite is formed in the central portion, thereby allowing the Kca value to be 6000 or less at a temperature of ⁇ 10° C. In addition, it can be confirmed that due to a relatively high carbon equivalent (Ceq) value, a relatively small amount of MA phase was present in the HAZ, but the CTOD value is 0.25 or less.
• an Ni content had a value higher than an upper limit value of an Ni content of an exemplary embodiment. It can be confirmed that due to a relatively high level of hardenability, a relatively large amount of upper bainite was generated in the central portion, thereby allowing the Kca value to be 6000 or less at a temperature of ⁇ 10° C. However, it can be confirmed that due to a relatively high Ni content, the CTOD value was relatively high.
• an Nb and Ti content has a value higher than an upper limit value of an Nb and Ti content of an exemplary embodiment. It can be confirmed that an entirety of other conditions satisfies a condition suggested in an exemplary embodiment, but due to a relatively high Nb and Ti content, a relatively large amount of the MA structure is generated in the HAZ, thereby allowing the CTOD value to be 0.25 mm or less.
• Inventive Example 7 includes a component exceeding a ratio of Cu to Ni suggested in an aspect of the present disclosure. It can be confirmed that despite having other, significantly excellent physical properties, a star crack was generated on a surface, thereby causing a default in surface quality.
• a C and Mn content has a value lower than a lower limit value of a C and Mn content of an exemplary embodiment. It can be confirmed that due to a relatively low level of hardenability, a fraction of AF+GB in the central portion is significantly low, and a relatively large amount of polygonal ferrite and a pearlite structure of 10% or greater are present, thereby allowing the Kca value to be 6000 or less at a temperature of ⁇ 10° C.
• AF+GB of a microstructure in the central portion was 70% or greater, a fraction of upper bainite in the central portion was 20% or less, a circle-equivalent diameter of an effective grain of upper bainite of the central portion having a high angle grain boundary of 15° or greater was 15 ⁇ m or less, and a fraction of the MA phase in the HAZ was less than 5%.

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• 2015
  • 2015-12-04 KR KR1020150172689A patent/KR101726082B1/ko active IP Right Grant
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