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 invention relates to a high strength steel having excellent brittle crack propagation resistance and resistance to brittle crack initiation at a welded part and a method of manufacturing the same.
• the microstructure of the ultra-thin material is coarse because sufficient deformation is not achieved due to the decrease in the total reduction ratio during the manufacture of the ultra-thin material, resulting in a low temperature which has the greatest influence on the grain size. Physical properties are reduced.
• the toughness of the welded heat affected zone (HAZ) becomes very weak as the microstructure of the heat affected zone (HAZ) consists of low-temperature transformation phases having high strength such as bainite.
• the island-like martensite generated from the unmodified austenite during the formation of low temperature transformation phase is the Since it is a nucleation site, it is very difficult to improve the brittle cracking resistance of high strength steels.
• the microstructure of the weld heat affected zone is refined by using TiN, or ferrite is formed on the weld heat affected zone by using oxide metallurgy.
• oxide metallurgy is partly helpful in improving the impact toughness through the microstructure of the tissue, but it does not have a significant effect in reducing the fraction of the phase martensite which has a major influence on the resistance to brittle crack initiation resistance.
• the resistance to brittle crack initiation of the base material can be secured by transforming the martensite phase into another phase through tempering, but in the case of a welding heat affected zone where the effect of tempering disappears due to thermal history It is impossible to apply this.
• One aspect of the present invention is to provide a high-strength steel excellent in brittle crack propagation resistance and resistance to weld brittle crack initiation, the object thereof.
• Another aspect of the present invention is to provide a method of manufacturing a high strength steel excellent in brittle crack propagation resistance and resistance to weld brittle crack initiation, an object thereof.
• C 0.05-0.09%, Mn: 1.5-2.2%, Ni: 0.3-1.2%, Nb: 0.005-0.04%, Ti: 0.005-0.04%, Cu: 0.1 -0.8%, Si: 0.05-0.3%, Al: 0.005-0.05%, P: 100 ppm or less, S: 40 ppm or less, and remaining Fe and other unavoidable impurities;
• MA martensite
• the microstructure of the surface part of the area of 2 mm or less directly below the surface is area%, and is composed of at least 20% of ferrite and at least one of the remaining bainite and martensite;
• a high-strength steel having an area% of the weld heat affected zone formed during welding, having excellent brittle crack propagation resistance and resistance to weld brittle crack initiation, including island martensite of 5% or less.
• the content of Cu and Ni may be set such that the Cu / Ni weight ratio is 0.8 or less, preferably 0.6 or less.
• the steel may preferably have a yield strength of 460 MPa or more.
• the steel may preferably have a Charpy wavefront transition temperature at a steel thickness of 1 / 2t (t: sheet thickness) in the steel thickness direction at ⁇ 40 ° C. or less.
• C 0.05 to 0.09%, Mn: 1.5 to 2.2%, Ni: 0.3 to 1.2%, Nb: 0.005 to 0.04%, Ti: 0.005 to 0.04%, Cu: 0.1-0.8%, Si: 0.05-0.3%, Al: 0.005-0.05%, Reheating the slab containing P: 100 ppm or less, S: 40 ppm or less, remaining Fe and other unavoidable impurities to 1000 to 1100 ° C., followed by rough rolling at a temperature of 1100 to 900 ° C .; Obtaining a steel sheet by finishing rolling the rough rolled bar at a temperature between Ar 3 + 60 ° C. and Ar 3 ° C. based on the central temperature; And it provides a brittle crack propagation resistance and welded brittle crack initiation resistance excellent resistance method comprising the step of cooling the steel sheet to a temperature below 500 °C.
• the reduction rate per pass is preferably 5% or more and the total cumulative reduction rate is 40% or more.
• the strain rate is preferably 2 / sec or less.
• the thickness center grain size of the bar after rough rolling and before finishing rolling may be 150 ⁇ m or less, preferably 100 ⁇ m or less, and more preferably 80 ⁇ m or less.
• the rolling reduction ratio during the finish rolling may be set so that the ratio of slab thickness (mm) / thickness of the steel sheet thickness (mm) after finishing rolling is 3.5 or more, preferably 4 or more.
• the cumulative reduction rate during the finish rolling is preferably maintained at 40% or more, and the reduction rate per pass excluding final shape even rolling is preferably maintained at 4% or more.
• Final shape even rolling refers to the process of rolling with low rolling reduction to ensure flatness of the plate
• the steel sheet may be cooled at a central cooling rate of 2 ° C./s or more.
• Cooling of the steel sheet can be carried out at an average cooling rate of 3 ⁇ 300 °C / s.
• the inventors of the present invention conducted studies and experiments to improve the yield strength and the brittle crack propagation resistance and the weld brittle crack initiation resistance of thick steel, and proposed the present invention based on the results.
• the present invention improves the yield strength, the brittle crack propagation resistance and the weld brittle crack initiation resistance of the thick steel by controlling the steel composition, structure and manufacturing conditions of the steel.
• the main concept of the present invention is as follows.
• Steel composition is appropriately controlled to improve strength through improving hardenability.
• Mn, Ni and Cu content is optimized with the carbon content to improve the hardenability.
• the microstructure is secured to the center of the thick steel.
• the composition is appropriately controlled to control the fraction of the martensite phase.
• the C, Si and Nb content that affects the production of phase martensite is optimized.
• the structure of the steel in order to improve strength and brittle crack propagation resistance.
• region was controlled in the thickness direction of steel materials.
• the resistance to brittle crack propagation is improved by excluding the microstructure that promotes the formation of cracks.
• the rough rolling conditions can be controlled in order to refine the structure of the steel.
• the microstructure is secured at the center by controlling the rolling reduction condition during rough rolling. This also promotes the production of acicular ferrite and granular bainite.
• finishing rolling conditions are controlled.
• finishing rolling temperature and rolling conditions to produce a large amount of strain bands in the austenite during the finish rolling to secure a large amount of ferrite nucleus site (site) to ensure a fine structure to the center of the steel. This also promotes the production of acicular ferrite and granular bainite.
• High strength steels having excellent resistance to brittle crack propagation and resistance to brittle crack initiation of welds which are an aspect of the present invention, are by weight%, C: 0.05 to 0.09%, Mn: 1.5 to 2.2%, Ni: 0.3 to 1.2%, and Nb: 0.005 to 0.04%, Ti: 0.005-0.04%, Cu: 0.1-0.8%, Si: 0.05-0.3%, Al: 0.005-0.05%, P: 100 ppm or less, S: 40 ppm or less and the remaining Fe and other unavoidable impurities; Central microstructure with area%, mixed phase of more than 70% of acicular ferrite and granular bainite, less than 20% of upper bainite, and the rest, ferrite, pearlite, The equivalent diameter of the effective crystal grains consisting of at least one selected from the group consisting of phase martensite (MA) and having a high boundary angle of 15 degrees or more measured by the upper bainite EBSD method is 15 ⁇ m (micrometer) or less
• the microstructure of the surface part of the area of 2 mm or less directly below the surface is area%, and is composed of at least 20% of ferrite and at least one of the remaining bainite and martensite;
• the weld heat affected zone formed at the time of welding includes an area martensite of 5% or less.
• C is the most important element for securing basic strength, it needs to be contained in steel within an appropriate range, and in order to obtain such an addition effect, it is preferable to add C 0.05% or more.
• the content of C is preferably limited to 0.05 ⁇ 0.09%.
• the content of C is more preferably 0.055% to 0.08%, and even more preferably 0.06% to 0.075%.
• Mn is a useful element that improves the strength by solid solution strengthening and improves the hardenability to produce a low temperature transformation phase.
• it is possible to generate a low temperature transformation phase even at a slow cooling rate due to the improvement of the hardenability, it is a major element for securing the strength of the core of the ultra-thick material.
• the Mn content is preferably limited to 1.5 to 2.2%.
• the content of Mn is more preferably 1.6 to 2.0%, and even more preferably 1.65 to 1.95%.
• Ni is an important element for facilitating cross slip of dislocations at low temperatures, improving impact toughness, improving hardenability, and improving strength, and 0.3% or more is preferably added to obtain such effects.
• the Ni is added more than 1.2%, the hardenability is excessively increased to form low-temperature transformation phase to reduce toughness, and due to the high cost of Ni compared to other hardenable elements, the manufacturing cost may also increase, so the upper limit of the Ni content is 1.2. It is preferable to limit to%.
• Ni is 0.4 to 1.0%, More preferably, it is limited to 0.45 to 0.9%.
• Nb precipitates in the form of NbC or NbCN to improve the base material strength.
• Nb dissolved in reheating at a high temperature precipitates very finely in the form of NbC during rolling, thereby suppressing recrystallization of austenite, thereby miniaturizing the structure.
• Nb is preferably added in an amount of 0.005% or more.
• Nb promotes the generation of phase martensite in the weld heat affected zone, thereby lowering the brittle crack initiation resistance, and may cause brittle cracks at the edges of the steel.
• the upper limit of the Nb content is preferably limited to 0.04%.
• the content of Nb is more preferably 0.01 to 0.035%, and even more preferably 0.015 to 0.03%.
• Ti is a component that precipitates with TiN upon reheating and inhibits the growth of crystal grains of the base metal and the weld heat affected zone to greatly improve low-temperature toughness. To obtain such an additive effect, Ti is preferably added at least 0.005%.
• the Ti content is preferably limited to 0.005 to 0.04%.
• More preferable content of Ti is 0.008 to 0.03%, More preferably, it is limited to 0.01 to 0.02%.
• Si is a substitutional element
• the strength of steel is improved through solid solution strengthening, and since it has a strong deoxidation effect, it is preferable to add 0.05% or more since it is an essential element for clean steel production.
• coarse phase martensite (MA) phase may be generated to lower brittle crack propagation and weld brittle crack initiation resistance, so the upper limit of the Si content is preferably limited to 0.3%.
• More preferable content of Si is 0.1 to 0.25%, still more preferably limited to 0.1 to 0.2%.
• Cu is a major element for improving hardenability and solid solution strength and improving strength of steel and is a major element for increasing yield strength through generation of epsilon Cu precipitates when tempering is applied, so it is preferably added at least 0.1%. However, when a large amount is added, the slab may be cracked due to hot shortness in the steelmaking process, so the upper limit of the Cu content is preferably limited to 0.8%.
• More preferable content of Cu is 0.2 to 0.6%, still more preferably limited to 0.25 to 0.5%.
• the content of Cu and Ni may be set such that the Cu / Ni weight ratio is 0.8 or less, preferably 0.6 or less. Even more preferably, it is limited to 0.5 or less.
• the surface quality may be further improved.
• Al is a component that acts as a deoxidizer, and when added in excess, it may form inclusions and lower the toughness. Therefore, the content is preferably limited to 0.005 to 0.05%.
• P, S is an element that causes brittleness or forms coarse inclusions at grain boundaries, and is preferably limited to P: 100 ppm or less and S: 40 ppm or less in order to improve brittle crack propagation resistance.
• the remaining component of the present invention is iron (Fe).
• the steel of the present invention has a central microstructure of area%, 70% or more of acicular ferrite and granular bainite, and 20% or less of upper bainite, and the rest , At least one member selected from the group consisting of ferrite, pearlite, phase martensite (MA), and the equivalent diameter of the effective crystal grain having a high boundary angle of 15 degrees or more measured by the EBSD method of the upper bainite is 15 ⁇ m.
• the microstructure of the area of 2 mm or less directly below the surface is the area%, and is composed of 20% or more of ferrite and one or more kinds of remaining bainite and martensite;
• the weld heat affected zone formed at the time of welding includes area martensite of 5% or less.
• the fraction of the mixed phase of the acicular ferrite and granular bainite of the central microstructure is less than 70%, it may be difficult to secure sufficient yield strength, for example, yield strength of 460 MPa or more. Can be difficult to secure.
• the fraction of the mixed phase of the eccentric ferrite and granular bainage is 75% or more, even more preferably 80% or more.
• the proportion of the acicular ferrite is preferably 20 to 70%.
• the fraction of the acicular ferrite exceeds 70% it may be difficult to secure sufficient yield strength due to the decrease in strength, for example, it may be difficult to secure a yield strength of 460MPa or more, less than 20% In this case, impact toughness may be lowered due to high strength.
• the fraction of the more preferable cyclic ferrite is 30-50%, More preferably, it is limited to 30-40%.
• the fraction of granular bainite is preferably 10 to 60%.
• the fraction of granular bainite is greater than 60%, impact toughness may decrease due to high strength, and if it is less than 10%, it may be difficult to secure sufficient yield strength due to the decrease in strength. For example, it may be difficult to secure a yield strength of 460 MPa or more.
• the fraction of more preferable granular bainite is 20 to 50%, even more preferably limited to 30 to 50%.
• the upper bainite at the center is preferably 20% or less.
• the more preferable fraction of upper bainite is 15% or less, and even more preferably 10% or less.
• the equivalent diameter of the effective grain having a high boundary angle of 15 degrees or more measured by the EBSD method of the upper center bainite exceeds 15 ⁇ m (micrometer), cracks are easily induced in spite of the low fraction of upper bainite. Since there is a problem, it is preferable that the circular equivalent diameter of the effective grains of the upper bainite in the center is 15 ⁇ m (micrometer) or less.
• brittle crack propagation resistance can be improved by effectively preventing crack propagation at the surface during brittle crack propagation.
• the fraction of more preferable ferrite is 30% or more, More preferably, it is limited to 40% or more.
• the ferrite in the central and surface microstructure refers to polygonal ferrite or elongatged polygonal ferrite.
• the in-phase martensite of the welded heat affected zone of the steel When the in-phase martensite of the welded heat affected zone of the steel is more than 5%, it acts as a starting point of cracking and lowers the brittle crack initiation resistance, so that the fraction of the in-phase martensite of the welded heat affected zone is 5% or less.
• Welding heat input during the welding may be 0.5 ⁇ 10kJ / mm.
• the welding method in the welding is not particularly limited, and examples thereof include FCAW (Flux Cored Arc Welding) and SAW (Submerged Arc Welding).
• the steel may preferably have a yield strength of 460 MPa or more.
• the steel may preferably have a Charpy wavefront transition temperature at a steel thickness of 1 / 2t (t: sheet thickness) in the steel thickness direction at ⁇ 40 ° C. or less.
• the steel material may have a thickness of 50 mm or more, and preferably may have a thickness of 50 to 100 mm.
• Another aspect of the present invention is a method of manufacturing a high strength steel having excellent resistance to brittle crack propagation and resistance to brittle crack initiation at a welded part in weight%, C: 0.05 to 0.09%, Mn: 1.5 to 2.2%, Ni: 0.3 to 1.2%, and Nb: 0.005 ⁇ 0.04%, Ti: 0.005 ⁇ 0.04%, Cu: 0.1 ⁇ 0.8%, Si: 0.05 ⁇ 0.3%, Al: 0.005 ⁇ 0.05%, P: 100ppm or less, S: 40ppm or less, remaining Fe and other unavoidable impurities Reheating the slab to 1000 to 1100 ° C.
• the slab reheating temperature is preferably at least 1000 ° C, in order to solidify the carbonitrides of Ti and / or Nb formed during casting.
• the upper limit of the reheating temperature is preferably 1100 ° C.
• the rough rolling temperature is preferably limited to 1100 ⁇ 900 °C.
• More preferable crude rolling temperature is 1050-950 degreeC.
• the rolling reduction rate per pass is preferably 5% or more and the total cumulative reduction rate is 40% or more for the last three passes during rough rolling.
• More preferred rolling reduction per pass is 7-20%
• More preferred total cumulative reduction is at least 45%.
• the recrystallized structure causes grain growth due to the high temperature, but during the last three passes, the grain growth rate is slowed down as the bar is air-cooled in the rolling atmosphere.
• the reduction rate of the pass has the greatest influence on the particle size of the final microstructure.
• the total cumulative reduction rate during rough rolling is preferably set to 40% or more in order to refine the structure of the central portion.
• the strain rate is preferably 2 / sec or less.
• the rough rolled bar is finish rolled at Ar 3 (ferrite transformation start temperature) + 60 ° C. to Ar 3 ° C. to obtain a steel sheet.
• More preferable cumulative reduction rate is 40 to 80%
• More preferable rolling reduction per pass is 4.5% or more.
• Finish rolling temperature is Ar 3
• coarse ferrite is produced before rolling and elongated during rolling, thereby lowering the impact toughness, and when finish rolling at Ar 3 + 60 ° C. or higher, it is not effective for miniaturization of particle size.
• the reduction rate of the unrecrystallized region during finishing rolling It is desirable to limit it to 40 to 80%.
• the thickness center grain size of the bar after the rough rolling and before the finish rolling may be 150 ⁇ m or less, preferably 100 ⁇ m or less, and more preferably 80 ⁇ m or less.
• the thickness center grain size of the bar after the rough rolling and the finish rolling may be controlled according to rough rolling conditions.
• the final microstructure is refined by the austenite grain refining, thereby improving the low-temperature impact toughness.
• the rolling reduction ratio during the finish rolling may be set so that the ratio of slab thickness (mm) / thickness of the steel sheet thickness (mm) after finishing rolling is 3.5 or more, preferably 4 or more.
• the steel sheet may have a thickness of 50 mm or more, preferably 50 to 100 mm.
• the steel sheet After finish rolling, the steel sheet is cooled to 500 ° C or lower.
• cooling end temperature exceeds 500 °C it may be difficult to secure a sufficient yield strength because the microstructure is not formed properly, for example, it may be difficult to secure a yield strength of 460MPa or more.
• Preferred cooling end temperature is 400 °C or less.
• the cooling of the steel sheet may be performed at a central cooling rate of 2 ° C./s or more, and when the central cooling rate of the steel sheet is less than 2 ° C./s, the microstructure may not be properly formed, and sufficient yield strength may be difficult to secure. For example, it may be difficult to secure a yield strength of 460 MPa or more.
• the steel sheet may be cooled at an average cooling rate of 3 to 300 ° C / s.
• the roughly rolled bar had a thickness of 192 mm, and the grain size of the core before rough rolling after rough rolling was 66 to 82 ⁇ m as shown in Table 2 below.
• the rolling reduction rate of the last three passes during the rough rolling was made within 7.9 ⁇ 14.1%, the deformation rate during rolling was carried out in the range of 1.22 ⁇ 1.68 / s.
• finish rolling was carried out at 50% cumulative reduction rate at a reduction ratio of 4.2 to 5.6% per pass at a temperature of the difference between the finish rolling temperature and the Ar 3 temperature shown in Table 2 to have the thickness of Table 3 below.
• finish rolling was carried out at 50% cumulative reduction rate at a reduction ratio of 4.2 to 5.6% per pass at a temperature of the difference between the finish rolling temperature and the Ar 3 temperature shown in Table 2 to have the thickness of Table 3 below.
• the Kca value of Table 4 is the value obtained by performing the ESSO test on the steel sheet, CTOD value was subjected to FCAW (1.0kJ / mm) welding to conduct a tissue analysis and CTOD evaluation for the weld heat affected zone The result is.
• Example No. Steel grade Grain size in the center before rough rolling after rough rolling ( ⁇ m) Average rolling reduction of the last 3 passes during rough rolling (%) Average strain rate of the last 3 passes during rough rolling (/ s) Average rolling reduction per pass during finishing rolling (%) Finish rolling temperature-Ar 3 temperature (°C) Central cooling rate (°C / sec) Cooling end temperature (°C) Inventive Example 1 Inventive Steel 1 78 11.3 1.61 4.2 35 4.1 324 Inventive Example 2 Inventive Steel 2 66 9.6 1.35 4.5 43 4.4 285 Inventive Example 3 Invention Steel 3 79 10.3 1.44 5.1 29 4.3 296 Inventive Example 4 Inventive Steel 4 75 8.9 1.46 5.3 41 3.8 335 Inventive Example 5 Inventive Steel 5 69 13.2 1.67 4.8 23 3.9 342 Inventive Example 6 Inventive Steel 6 73 14.1 1.32 4.2 15 4.3 312 Comparative Example 1 Inventive Steel 7 79 12.2 1.22 4.7 93 4.8 256 Comparative Example 2 Compar
• the finish rolling temperature-Ar3 temperature difference at the time of finish rolling proposed in the present invention is controlled to 60 ° C. or more, and the rolling is carried out at a high temperature, so that sufficient reduction to the center part is achieved. It can be seen that as the cooling is started at a high temperature, no more than 20% of ferrite is formed on the surface, so that the Kca value measured at -10 ° C does not exceed the 6000 required for general shipbuilding steels.
• the content of C is higher than the upper limit of the C content of the present invention, and Kca measured at -10 ° C due to the formation of a large amount of coarse upper bainite in the center during rough rolling. It can be seen that the value has a value of 6000 or less, a large amount of phase-like martensite (MA) structure is also generated in the weld heat affected zone, the CTOD value is 0.25mm or less.
• MA phase-like martensite
• the Si content is higher than the upper limit of the Si content of the present invention.
• a large amount of Si is added, a large amount of MA structure is generated in the weld heat affected zone, and the CTOD value is 0.25 mm or less. It can be seen that.
• the Mn content has a higher value than the upper limit of the Mn content of the present invention. Since the high baenite has a large amount of upper bainite in the center, the Kca value also has a value of 6000 or less at -10 ° C. Able to know. In addition, due to the high Ceq value, the CTOD value is 0.25 or less, although there are few MA phases in the weld heat affected zone.
• the Ni content is higher than the upper limit of the Ni content of the present invention, and a large amount of upper bainite was generated in the center with high curing ability, so that the Kca value was also lower than 6000 at -10 ° C. It can be seen that it has. However, due to the high Ni content, the CTOD value is excellent.
• Inventive Example 7 it has a component exceeding the Cu / Ni ratio suggested in one preferred aspect of the present invention, and although it is excellent in other physical properties, it is known that star cracks are generated on the surface and thus there is an abnormality in the surface quality. Can be.
• AF + GB of the central microstructure has 70% or more, the fraction of bainite in the upper part of the center is 20% or less, and in the upper part of the center. It can be seen that the circular equivalent diameter of the effective crystal grain having a high boundary angle of bainite of 15 degrees or more is 15 ⁇ m or less, and the MA phase fraction of the weld heat affected zone is less than 5%.
• Inventive Examples 1 to 6 yield strength 460MPa or more, Kca value satisfies the value of 6000 or more at -10 °C, CTOD value can also be seen that excellent value of 0.25mm or more.

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• 2015
  • 2015-12-04 KR KR1020150172689A patent/KR101726082B1/ko active IP Right Grant
• 2016
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