WO2023090751A1 - Plaque d'acier à haute résistance et à haute limite d'élasticité ayant une excellente résistance aux chocs après formage à froid et son procédé de fabrication - Google Patents

Plaque d'acier à haute résistance et à haute limite d'élasticité ayant une excellente résistance aux chocs après formage à froid et son procédé de fabrication Download PDF

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WO2023090751A1
WO2023090751A1 PCT/KR2022/017540 KR2022017540W WO2023090751A1 WO 2023090751 A1 WO2023090751 A1 WO 2023090751A1 KR 2022017540 W KR2022017540 W KR 2022017540W WO 2023090751 A1 WO2023090751 A1 WO 2023090751A1
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steel sheet
less
temperature
thickness
steel
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PCT/KR2022/017540
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Korean (ko)
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김성일
나현택
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주식회사 포스코
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    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21C—MANUFACTURE OF METAL SHEETS, WIRE, RODS, TUBES OR PROFILES, OTHERWISE THAN BY ROLLING; AUXILIARY OPERATIONS USED IN CONNECTION WITH METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL
    • B21C47/00—Winding-up, coiling or winding-off metal wire, metal band or other flexible metal material characterised by features relevant to metal processing only
    • B21C47/02—Winding-up or coiling
    • 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
    • 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
    • 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
    • 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
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • 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 high-strength steel and a method for manufacturing the same, and more particularly, to a high-strength steel having excellent impact resistance after cold forming and a high yield ratio, and a method for manufacturing the same.
  • steel with a thickness of 8 mm or more is mainly applied due to the characteristics of the vehicle, and high-strength hot-rolled steel with a tensile strength of 600 MPa is used for members and hot-rolled steel with a tensile strength of 440 MPa is used for wheel disks.
  • high-strength steel sheets with a tensile strength of 700 MPa or more and a tensile strength of 590 MPa or more are applied to members and wheel disks, and the thickness of the steel sheet is reduced or the design of the parts is changed.
  • the wheel has been conventionally manufactured by a press forming process, but recently, it is a trend to be manufactured by spinning and flow forming. Since such a forming process gives a greater amount of deformation to the hot-rolled steel sheet, a hot-rolled steel sheet having a higher elongation is required, and the molded part is required to secure durability and impact resistance during use.
  • the durability of the part should be equal to or higher than that of the existing one, but when forming the part, fine cracks occur on the shear surface, etc. It is difficult to apply, such as cancellation.
  • Patent Documents 1 and 2 the conventional high-strength steel is wound at high temperature after undergoing normal austenitic hot rolling to form ferrite as a base structure and fine precipitates, and as in Patent Document 3, coarse To prevent the formation of pearlite, the winding temperature is cooled to the temperature at which the bainite base structure is formed, and then the winding technology is applied.
  • Patent Document 4 a technique of refining austenite crystal grains by reducing the size to 40% or more in the non-recrystallization region during hot rolling using Ti, Nb, etc. has also been proposed.
  • Patent Document 5 a technique for improving the uniformity of the microstructure between the thick surface layer portion and the deep layer portion of the steel sheet and suppressing the formation of coarse carbides has been proposed, and as in Patent Document 6, pearlite and MA phases that adversely affect durability (Martensite and Austenite) and a technique for suppressing the formation of martensite at the same time have been proposed.
  • Patent Documents 1 to 4 do not consider the occurrence of cracks on the shear surface and its surroundings during shear forming of high-strength thick material materials, and in the case of thick material materials having a thickness of 8 mm or more, it is difficult to secure cooling rate conditions and large pressure made up of conditions.
  • precipitate-forming elements such as Ti, Nb, and V to secure strength while miniaturizing the crystal grains of the material material
  • ferrite grows excessively and the yield strength decreases. There is a problem of reduced and coarse pearlite formation.
  • Patent Literatures 5 and 6 are inventions for post-material materials.
  • Patent Document 5 provides a technology for manufacturing to suppress the formation of MA phase and martensite by making the crystal grain shape in the deep-thickness portion (1/4t to 1/2t) have equiaxed and fine crystal grains in order to improve the durability of thick high-strength steel. suggested.
  • Patent Document 6 proposes a technique for manufacturing by dividing a hot-rolled coil into three parts in the longitudinal direction through a relational expression derived for a specific component, cooling the head, mid, and tail under constant cooling rate conditions to different cooling end temperatures, respectively, and then winding it up. It became. These technologies are technologies that uniformly manufacture microstructures by controlling the cooling rate after hot rolling through a relational expression derived for a specific component in consideration of the quality of the shear surface of the part. There is an effective aspect to improve the durability of the applied commercial vehicle wheel, but the impact resistance after molding has not been considered. In addition, it is difficult to uniformly control the cooling of the steel sheet after hot rolling over the entire width, and in the case of a thick hot-rolled steel sheet having a thickness of 8 mm or more, it is difficult to control the actual cooling rate.
  • the cooling process after hot rolling usually takes place within tens of seconds at a ROT (Run-Out Table) with a length of 100 to 120 m. difficult. Therefore, in the prior art, the hot-rolled steel sheet having a thickness of 8 mm or more is difficult to obtain an effect of inhibiting the formation of coarse carbides and is insufficient to secure a high level of impact resistance.
  • Patent Document 1 Korean Patent Publication No. 10-2010-0029138 (published on March 15, 2010)
  • Patent Document 2 Japanese Unexamined Patent Publication No. 2007-262487 (published on October 11, 2007)
  • Patent Document 3 Korean Patent Registration No. 10-1528084 (published on June 10, 2015)
  • Patent Document 4 Japanese Unexamined Patent Publication No. 9-143570 (published on June 3, 1997)
  • Patent Document 5 Korean Patent Publication No. 10-2020-0062422 (published on June 4, 2020)
  • Patent Document 6 Korean Patent Publication No. 10-2021-0068808 (published on June 10, 2021)
  • it is intended to provide a high-strength steel having excellent impact resistance and a high yield ratio after cold forming, and a manufacturing method thereof.
  • C 0.05 to 0.15%
  • Si 0.01 to 0.5%
  • Mn 1.0 to 2.0%
  • Al 0.01 to 0.1%
  • Cr 0.001 to 1.0%
  • P 0.001 to 0.001% 0.05%
  • S 0.001 to 0.01%
  • N 0.001 to 0.01%
  • Ti 0.03 to 0.08%
  • Nb 0.01 to 0.05%
  • the microstructure of the surface layer with a thickness of 50 ⁇ m from the surface contains, by area%, 95% or more of equiaxed ferrite and 3% or less of pearlite, and contains bainitic ferrite, bainite, MA (Martensite-Austenite constituent) phase and martensite one or more of them in a total of 5% or less,
  • the microstructure of the center in the range of 1/4 to 3/4 in thickness is 80 to 95% of bainitic ferrite, 10% or less of bainite, 3% or less of pearlite, MA (Martensite-Austenite constituent) phase and It is possible to provide a steel sheet containing 5 to 10% of one or two martensite in total and containing equiaxed ferrite as the remainder.
  • the steel sheet may have a thickness of 8 to 25 mm.
  • the steel sheet may have tensile strength of 590 MPa or more, elongation at break of 25% or more, yield ratio of 0.75 to 0.9, and impact toughness at -20° C. of 70 J or more after cold forming.
  • the steel sheet may have a ratio (E/YS) of impact toughness (E) at -20° C. after cold forming and yield strength (YS) before cold forming (E/YS) of 0.15 or more.
  • the steel sheet includes an edge portion corresponding to 30% of the area from both ends and a center portion of the central 40% area corresponding to the area excluding both edge portions based on the width direction, and the edge portion and the center portion have a tensile strength of
  • the difference is 10 MPa or less, the difference in breaking elongation is 8% or less, and the difference in impact toughness at -20 ° C. after cold forming may be 20 J or less.
  • C 0.05 to 0.15%
  • Si 0.01 to 0.5%
  • Mn 1.0 to 2.0%
  • Al 0.01 to 0.1%
  • Cr 0.001 to 1.0%
  • P 0.001 ⁇ 0.05%
  • S 0.001 ⁇ 0.01%
  • N 0.001 ⁇ 0.01%
  • Ti 0.03 ⁇ 0.08%
  • Nb 0.01 ⁇ 0.05%
  • the edge portion corresponding to 30% of the area at both ends based on the coil width direction has a temperature (TE) of 550 to 650 ° C.
  • the center 40 corresponding to the area excluding both edge portions in the width direction The central part of the % region is cooled to a temperature (TC) of 500 to 550 ° C,
  • the reheating temperature is 1100 ⁇ 1350 °C
  • the hot rolling temperature may be 800 ⁇ 1150 °C.
  • Air-cooling the wound coil to a temperature range of 200° C. or less may be further included.
  • the steel sheet may have a thickness of 8 to 25 mm.
  • a high-strength steel that can be applied as a steel material used for chassis members and wheels of medium-large commercial vehicles and a manufacturing method thereof.
  • Example 1 is a photograph showing the microstructure of Inventive Example 2 according to an embodiment of the present invention observed with a scanning electron microscope (x3,000).
  • Figure 2 is a photograph showing the microstructure of Comparative Example 2 according to an embodiment of the present invention observed with a scanning electron microscope (x3,000).
  • the present inventors studied the change in impact resistance after cold forming according to the characteristics of the microstructure of the steel sheet in order to solve the above-mentioned conventional problems and to secure excellent formability and impact resistance. Accordingly, it was confirmed that target physical properties can be secured by controlling the microstructure along the thickness and width direction of the steel sheet by optimizing the alloy composition and manufacturing conditions, and the present invention has been completed.
  • a hot-rolled steel sheet usually manufactured in the form of a coil coarse carbides and pearlite are likely to be formed when maintained in a high temperature region of about 500 to 700 ° C. for a long time.
  • the solid solution of carbon increases in the non-transformed phase, making it easy to form coarse carbides or pearlite.
  • the cooling rate of the central part of the width of the coil is slower than that of the edge part, so that such a structure is further developed. Therefore, in order to suppress the formation of such coarse carbides and pearlite at the center of the coil width, it is necessary to cool the wound coil to room temperature through forced cooling such as water cooling.
  • the present invention proposes a method capable of suppressing the formation of coarse carbides and pearlite without forcibly cooling the coil.
  • % indicating the content of each element is based on weight.
  • Carbon (C) is the most economical and effective element for strengthening steel, and when the added amount increases, the precipitation hardening effect or the bainite phase fraction increases, making it easy to secure strength.
  • the cooling rate at the center of the thickness decreases during cooling after hot rolling, so that coarse carbides or pearlite are easily formed when the carbon (C) content is high. Therefore, if the carbon (C) content is less than 0.05%, it is difficult to obtain a sufficient strengthening effect, and if the content exceeds 0.15%, there is a problem in that impact resistance is lowered due to the formation of pearlite or coarse carbides in the center of the thickness, and the weldability is also poor. There is a risk of A more preferred lower limit may be 0.055%, and a more preferred upper limit may be 0.12%.
  • Silicon (Si) is an element that is effective in deoxidizing molten steel and strengthening steel by solid solution, and is also effective in improving formability by delaying the formation of coarse carbides.
  • the content is less than 0.01%, the effect of delaying solid solution strengthening and carbide formation cannot be maximized, and if the content exceeds 0.5%, red scale is formed on the surface of the steel sheet during hot rolling, resulting in very poor quality.
  • Manganese (Mn), like Si, is an element effective for solid solution strengthening of steel, and may facilitate the formation of bainite during cooling after hot rolling by increasing hardenability of steel.
  • the content is less than 1.0%, the above effects due to addition cannot be obtained, and if it exceeds 2.0%, the hardenability is greatly increased, so that martensitic phase transformation is likely to occur and segregation at the center of the thickness is greatly developed during slab casting in the casting process.
  • Aluminum (Al) is a component mainly added for deoxidation, and if the content thereof is less than 0.01%, the effect of addition may be insufficient. On the other hand, when the content exceeds 0.1%, AlN is formed by combining with nitrogen, and thus corner cracks easily occur in the slab during continuous casting, and defects due to inclusion formation easily occur. More preferably 0.015% or more, more preferably 0.06% or less.
  • Chromium (Cr) strengthens steel by solid solution similarly to Mn, and serves to help bainite formation by delaying ferrite phase transformation during cooling.
  • the content is less than 0.001%, the effect of addition cannot be obtained, and if it exceeds 1.0%, ferrite transformation is excessively delayed, and elongation may be inferior due to excessive martensite formation.
  • excessive addition of Cr greatly develops segregation at the center of the thickness, and makes the microstructure in the thickness direction non-uniform, resulting in inferior impact resistance.
  • a more preferred lower limit may be 0.01%, and a more preferred upper limit may be 0.5%.
  • Phosphorus (P) has the effect of strengthening solid solution and accelerating ferrite transformation at the same time as Si.
  • a phosphorus (P) content of less than 0.001%, a lot of manufacturing costs are required, which is economically unfavorable, and is insufficient to obtain strength.
  • the content exceeds 0.05%, brittleness occurs due to grain boundary segregation, and it is easy to generate fine cracks during molding, which can significantly deteriorate impact resistance.
  • S is an impurity present in steel, and when its content exceeds 0.01%, it combines with Mn to form non-metallic inclusions. Accordingly, there is a problem in that it is easy to generate fine cracks during molding and significantly lowers the impact resistance. However, in order to manufacture the content to less than 0.001%, a lot of time is required during steelmaking operation, resulting in a decrease in productivity.
  • Nitrogen (N) is a representative solid-solution strengthening element together with C, and forms coarse precipitates with Ti and Al.
  • the solid solution strengthening effect of nitrogen (N) is superior to that of C, but as the amount of nitrogen (N) in steel increases, the toughness decreases significantly. Therefore, the upper limit of the content may be limited to 0.01%.
  • the upper limit of the content may be limited to 0.01%.
  • a lot of time is required during steelmaking operation, resulting in a decrease in productivity.
  • Titanium (Ti) is a typical precipitation hardening element and forms coarse TiN in steel due to its strong affinity with N. TiN has an effect of suppressing the growth of crystal grains during the heating process for hot rolling. In addition, titanium (Ti) remaining after reacting with N is dissolved in steel and combined with C to form TiC precipitates, which is a useful component for improving the strength of steel. If the content of titanium (Ti) is less than 0.03%, the above effect cannot be obtained, and if the content exceeds 0.08%, there is a problem of inferior crash resistance during molding due to generation of coarse TiN and coarsening of precipitates. More preferably, it may contain 0.04% or more, and more preferably, it may contain 0.075% or less.
  • Niobium (Nb) is a representative precipitation hardening element along with Ti, and is effective in improving the strength and impact toughness of steel by precipitating during hot rolling and refining grains by recrystallization delay. If the content of niobium (Nb) is less than 0.01%, the above effect cannot be obtained, and if the content exceeds 0.05%, formability is deteriorated due to the formation of elongated crystal grains and the formation of coarse complex precipitates due to excessive recrystallization delay during hot rolling. There is a problem with A more preferred lower limit may be 0.015%, and a more preferred upper limit may be 0.04%.
  • the steel of the present invention may include the remaining iron (Fe) and unavoidable impurities in addition to the above-described composition. Since unavoidable impurities may be unintentionally incorporated in the normal manufacturing process, they cannot be excluded. Since these impurities are known to anyone skilled in the steel manufacturing field, not all of them are specifically mentioned in this specification.
  • the sum of niobium (Nb) and titanium (Ti) may be 0.04 to 0.1%.
  • Niobium (Nb) and titanium (Ti) are precipitated as (Ti, Nb) (C, N)-based composite precipitates, and when precipitated during hot rolling, the crystal grain refinement effect due to recrystallization delay is greatly increased.
  • the formation of the complex precipitates is excessive, there is a problem in that the strength improvement effect is small and the formability is inferior because the coarse complex precipitates increase.
  • the sum of niobium (Nb) and titanium (Ti) is less than 0.04%, the effect of grain refinement and strength improvement may be small.
  • the sum exceeds 0.1%, formability is inferior and economically unfavorable.
  • a more preferred lower limit may be 0.045%, and a more preferred upper limit may be 0.09%.
  • % representing the fraction of the microstructure is based on the area unless otherwise specified.
  • the microstructure of the surface layer in the range of 50 ⁇ m in thickness from the surface is area%, containing 95% or more of equiaxed ferrite and 3% or less of pearlite, bainitic ferrite, bainite, MA ( It contains 5% or less of one or more of the martensite-austenite constituent) phase and martensite in total, and the microstructure in the center in the range of 1/4 to 3/4 in thickness is 80 to 95% of bainitic ferrite by area% , 10% or less of bainite, 3% or less of pearlite, 5 to 10% of one or two of MA (Martensite-Austenite constituent) phase and martensite in total, and equiaxed crystal ferrite as the remainder. there is.
  • the equiaxed crystal ferrite is less than 95% in the surface layer portion, ductility is insufficient during spinning and flow forming molding applied during commercial vehicle wheel manufacturing, and work hardening in the surface layer portion is severe, so that fine cracks may occur during molding.
  • cracks easily propagate along the interface with the matrix phase when brittle pearlite is formed at 3% or more, or at least one of bainitic ferrite, bainite, MA phase, and martensite with high hardness is included in an amount exceeding 5%.
  • equiaxed crystal ferrite may be 100% as the microstructure of the surface layer portion, and the sum of pearlite, bainitic ferrite, bainite, MA phase, and martensite may be 0%.
  • bainitic ferrite is less than 80% in the center, there is a problem that cracks easily occur on the shear surface during punching and shear molding during wheel manufacturing, and after molding, there is a problem that the impact resistance is also inferior.
  • bainitic ferrite which is a matrix structure, is formed in the center of the thickness of the rolled sheet, and a high concentration of residual C remains in the untransformed austenite, so it is easy to form pearlite. can At this time, when the pearlite is formed in excess of 3%, cracks are severely generated on the shear surface during the molding process, and the impact resistance after molding is also inferior.
  • the pearlite fraction is 3% or less, there is no cracking caused by molding such as shearing, and the impact resistance at low temperatures may be excellent.
  • carbides and nitrides having a diameter of 1 ⁇ m or more may be included as pearlite.
  • the MA phase or martensite when included in an amount of 5 to 10%, crack generation is not affected, and it may be advantageous to secure impact resistance and high strength after molding.
  • the MA phase has an advantage in securing high strength by forming dislocation density around it, and when formed with a base structure composed of ferrite and bainite, it can have excellent impact resistance even if the dislocation density increases after cold forming.
  • the MA phase or martensite when the MA phase or martensite is less than 5%, yield strength and tensile strength are insufficient, and when it is included in excess of 10%, there is a problem in that ductility is insufficient and formability is inferior.
  • the bainite content also exceeds 10%, there may be a problem of insufficient ductility.
  • bainite and pearlite may each be 0% as the microstructure of the center, and may inevitably include equiaxed ferrite in addition to bainitic ferrite, bainite, pearlite, MA phase, and martensite.
  • the area fraction of the microstructure can be analyzed using an optical microscope and a scanning electron microscope (SEM), and the area fraction of the image obtained from the image observed at 3,000 magnification at a location corresponding to the center of the thickness of the rolled section can be measured
  • Steel according to one aspect of the present invention can be produced by reheating, rolling, cooling and winding a steel slab satisfying the above-described alloy composition.
  • a steel slab satisfying the alloy composition of the present invention can be reheated to a temperature range of 1100 to 1350 ° C.
  • the reheating temperature is less than 1100 ° C., the formation of precipitates may be reduced in the process after hot rolling because the precipitates are not sufficiently re-dissolved, and coarse TiN may remain.
  • the temperature exceeds 1350 ° C., strength may be reduced due to abnormal grain growth of austenite grains.
  • the reheated steel slab may be hot-rolled in a temperature range of 800 to 1150 ° C.
  • the hot-rolling temperature exceeds 1150° C.
  • the temperature of the hot-rolled steel sheet increases, resulting in coarse grain size and poor surface quality of the hot-rolled steel sheet.
  • the temperature is less than 800 ° C, elongated crystal grains may develop due to recrystallization retardation, resulting in severe anisotropy and poor formability.
  • the hot-rolled steel sheet may be cooled and wound at an average cooling rate greater than or equal to the CR value defined in the following relational expression 1 within the range of 1 to 30 °C/s up to a temperature range of 500 to 650 °C.
  • the edge portion which is 30% from both ends in the coil width direction, has a temperature (TE) of 550 to 650 ° C
  • the central portion which is 40% of the center excluding both edge portions in the width direction, has a temperature (TC) of 500 to 550 ° C.
  • TE temperature
  • TC temperature difference between the edge portion and the central portion may be 50 to 150°C.
  • relational expression 1 was derived in order to induce an appropriate level of ferrite phase transformation, form a fine and uniform MA phase, and suppress excessive pearlite formation during cooling of the steel sheet.
  • the cooling rate is less than the CR value of relational expression 1
  • ferrite in the center of the thickness becomes coarse, pearlite is excessively formed, and cracks on the shear surface become severe, and impact resistance after molding may be inferior.
  • the cooling rate exceeds 30 °C / s, bainite, MA phase, and martensite are excessively formed, resulting in insufficient ductility and poor shear quality.
  • the cooling end temperature should be lowered during cooling after hot rolling, but there is a concern that it is difficult to secure the target elongation due to ferrite reduction due to excessive formation of bainite or excessive formation of MA phase and martensite. .
  • the cooling end temperature of the central portion in the width direction and the edge portion may be set differently.
  • the average temperature difference between the edge portion and the central portion may be 50 to 150°C. If the average temperature difference is less than 50 ° C., it may be difficult to obtain the above effect. On the other hand, if the temperature exceeds 150 °C, the above effect does not increase any more, but it may be difficult to control the temperature of each section of the coil.
  • a method of differently controlling the cooling end temperature of the edge portion and the center portion during winding is not particularly limited. For example, when cooling a hot-rolled steel sheet, the cooling water poured into the edge portion is blocked before reaching the steel sheet. Alternatively, a method of differently adjusting the amount of cooling water injected can be applied. Alternatively, both methods may be used in parallel.
  • both the relational expression 1 and the cooling end temperature condition it is preferable to satisfy both the relational expression 1 and the cooling end temperature condition in order to secure the desired strength, moldability and impact resistance.
  • bainitic ferrite is used as a base structure in the center of the thickness direction to have a uniform and fine microstructure, and coarse carbides or pearlite are reduced in the inner winding part and the center of the thickness of the coil with a slow cooling rate.
  • the non-uniform structure of the hot-rolled steel sheet can be eliminated.
  • non-uniform formation of the MA phase and formation of coarse martensite can be suppressed in the outer winding part and the edge part of the coil where the cooling rate is relatively fast.
  • the wound coil may be air-cooled to a temperature range of 200° C. or less.
  • the wound coil can be air-cooled to a temperature range of 200° C. or less.
  • Air cooling of the coil means cooling in the air at room temperature at a cooling rate of 0.001 to 10 °C/h. At this time, if the cooling rate exceeds 10 ° C / h, some of the non-transformed phase of the steel in the outer winding portion of the coil is easily transformed to the MA phase, and the shear formability and punching formability of the steel may be inferior.
  • a more preferred lower limit may be 0.01 °C/h, and a more preferred upper limit may be 1 °C/h.
  • the steel of the present invention thus prepared has a thickness of 8 to 25 mm, a tensile strength of 590 MPa or more, a breaking elongation of 25% or more, a yield ratio of 0.75 to 0.9, and an impact toughness at -20 ° C of 70 J after cold forming.
  • the ratio of the impact toughness after cold forming and the yield strength before cold forming is 0.15 or more, so that while having a high yield ratio, it is possible to have excellent impact toughness characteristics.
  • the steel includes an edge portion corresponding to 30% of the area from both ends and a center portion of the center 40% area corresponding to the area excluding both edge portions based on the width direction, and the edge portion and the center portion are tensile
  • the difference in strength is 10 MPa or less
  • the difference in elongation at break is 8% or less
  • the difference in impact toughness at -20 ° C after cold forming may be 20 J or less.
  • a steel slab having the alloy composition shown in Table 1 below was prepared as a hot-rolled steel sheet under the conditions shown in Table 2 below. At this time, the steel slab was reheated to a temperature of 1100 ⁇ 1350 °C and then hot rolled.
  • Table 2 shows the cooling rate applied during manufacturing and the CR value of relational expression 1, and the cooling end temperature is a temperature (TC) in the range of 40% in the center of the width direction and a temperature (TE) in the range of 30% in both edge parts in the width direction, respectively. indicated respectively.
  • TC temperature
  • TE temperature
  • the microstructure of the manufactured steel sheet was measured and described.
  • the microstructure was expressed by measuring the surface layer portion and the center portion in the thickness direction, respectively, and measuring the fractions of the edge portion and the center portion in the width direction, respectively.
  • the microstructure of the surface layer part was observed from the surface to the thickness of 50 ⁇ m, and the center part was observed for the 1/4 ⁇ 3/4t (25 ⁇ 75% section, t is the thickness (mm)) part from the steel plate surface in the thickness direction.
  • the microstructure of the portion corresponding to 30% of the edge portion was observed at both ends in the width direction, and the central portion was observed based on the portion corresponding to the central 40% excluding the edge portion.
  • the area fraction of the MA phase and martensite was measured using an optical microscope and an image analyzer after etching with the Lepera etching method, and was analyzed at 1,000 magnification.
  • the fractions of equiaxed ferrite (PF), bainitic ferrite (BF), bainite (B) and pearlite (P) were measured using a scanning electron microscope (SEM) at 3,000x and 5,000x magnifications.
  • SEM scanning electron microscope
  • PF is polygonal ferrite having an equiaxed crystal shape
  • BF may include ferrite observed in a low temperature region such as acicular ferrite and bainitic ferrite.
  • Table 4 the physical property values for each of the manufactured specimens are measured and shown for the central portion and the edge portion in the width direction.
  • Yield strength (YS), tensile strength (TS), yield ratio (YR), and breaking elongation (T-El) were evaluated by taking a JIS5 standard test piece in a direction perpendicular to the rolling direction and performing a tensile test.
  • the impact absorption energy (E) at -20 °C after cold forming was measured, and the ratio (E / YS) of impact absorption energy and yield strength at -20 °C after cold forming was shown.
  • the impact absorption energy was tested using a Charpy V-notch specimen prepared in accordance with the ASTM standard (ASTM A370), sampled in a direction perpendicular to the rolling direction.
  • FIG. 1 is a photograph showing the microstructure of a hot-rolled steel sheet of Inventive Example 2 according to one aspect of the present invention observed with a scanning electron microscope (x3,000).
  • Comparative Example 1 when the sum of Ti and Nb contents exceeded the scope of the present invention, impact resistance was poor due to the formation of coarse precipitates and TiN due to excessive precipitates in ferrite grains.
  • Comparative Example 2 is a case where the cooling rate did not meet the cooling rate standard proposed in relational expression 1, and as shown in FIG. 2, excessive pearlite was formed in the microstructure.
  • Comparative Example 4 is a case where the winding temperature of the edge portion in the width direction exceeds the range proposed in the present invention. Accordingly, the impact resistance was inferior due to the formation of excessive pearlite at the center of the thickness direction of the edge portion. This is because the temperature of the edge portion is high and the heat transfer of the winding coil proceeds slowly at the edge portion.
  • Comparative Example 5 is a case where the winding temperature of the central part in the width direction does not reach the range of the present invention, and bainite is excessively formed in the central part in the thickness direction of the central part, and pearlite, MA phase, martensite, etc. are at the level proposed in the present invention It was excessively formed and the elongation was inferior.
  • the edge portion satisfies the winding temperature range and the elongation rate and impact resistance were relatively good, but the ferrite in the surface layer portion fell short of the range proposed in the present invention. It is believed that this is because the temperature of the edge portion of the coil after winding also rapidly decreased due to the low cooling end temperature of the central portion in the width direction.
  • Comparative Example 7 is a case where the thickness of the steel is less than 8 mm, and since the cooling rate is excessively applied for a given steel component, ferrite is insufficient in the surface layer portion, and bainitic ferrite is reduced in the center in the thickness direction. At the same time, pearlite is excessively formed. . This is considered to be that the untransformed phase increased during the initial cooling process, and pearlite was formed in a region with a relatively high C content. As a result, it was not possible to secure a target level of elongation.
  • Comparative Example 9 is a case where the cooling rate exceeds the CR value of relational expression 1 and satisfies the range proposed in the present invention, but exceeds 30 ° C / s. As a result, the polygonal ferrite in the surface layer was insufficient, and bainite was excessively formed in the center in the thickness direction, so that the desired level of elongation could not be secured.

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Abstract

La présente invention concerne une plaque d'acier à haute résistance et son procédé de fabrication et, plus particulièrement, une plaque d'acier à haute résistance ayant une excellente résistance aux chocs après formage à froid et ayant une limite d'élasticité élevée, et son procédé de fabrication.
PCT/KR2022/017540 2021-11-17 2022-11-09 Plaque d'acier à haute résistance et à haute limite d'élasticité ayant une excellente résistance aux chocs après formage à froid et son procédé de fabrication WO2023090751A1 (fr)

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KR1020210158366A KR20230072050A (ko) 2021-11-17 2021-11-17 냉간 성형 후 내충격성이 우수한 고항복비형 고강도강 및 그 제조방법
KR10-2021-0158366 2021-11-17

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Publication number Priority date Publication date Assignee Title
JPH09143570A (ja) 1995-11-17 1997-06-03 Kawasaki Steel Corp 極微細組織を有する高張力熱延鋼板の製造方法
JP2007262487A (ja) 2006-03-28 2007-10-11 Nippon Steel Corp 伸びフランジ性に優れた高強度鋼板
KR20100029138A (ko) 2007-07-31 2010-03-15 제이에프이 스틸 가부시키가이샤 고강도 강판
KR101304859B1 (ko) * 2009-12-04 2013-09-05 주식회사 포스코 표면균열 저항성이 우수한 초고강도 라인파이프용 강판 및 그 제조방법
KR101333854B1 (ko) * 2009-01-30 2013-11-27 제이에프이 스틸 가부시키가이샤 저온 인성이 우수한 후육 고장력 열연 강판 및 그 제조 방법
US20140205855A1 (en) * 2011-07-29 2014-07-24 Nippon Steel & Sumitomo Metal Corporation High-strength steel sheet excellent in impact resistance and manufacturing method thereof, and high-strength galvanized steel sheet and manufacturing method thereof
KR101528084B1 (ko) 2010-09-17 2015-06-10 제이에프이 스틸 가부시키가이샤 타발 가공성이 우수한 고강도 열연 강판 및 그 제조 방법
JP2016060955A (ja) * 2014-09-19 2016-04-25 株式会社神戸製鋼所 熱延鋼板及びその製造方法
KR20190077829A (ko) * 2017-12-26 2019-07-04 주식회사 포스코 고강도 고인성 열연강판 및 그 제조방법
KR20200062422A (ko) 2018-11-26 2020-06-04 주식회사 포스코 내구성이 우수한 고강도 강재 및 이의 제조방법
KR20210068808A (ko) 2019-12-02 2021-06-10 주식회사 포스코 내구성이 우수한 후물 복합조직강 및 그 제조방법

Patent Citations (11)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPH09143570A (ja) 1995-11-17 1997-06-03 Kawasaki Steel Corp 極微細組織を有する高張力熱延鋼板の製造方法
JP2007262487A (ja) 2006-03-28 2007-10-11 Nippon Steel Corp 伸びフランジ性に優れた高強度鋼板
KR20100029138A (ko) 2007-07-31 2010-03-15 제이에프이 스틸 가부시키가이샤 고강도 강판
KR101333854B1 (ko) * 2009-01-30 2013-11-27 제이에프이 스틸 가부시키가이샤 저온 인성이 우수한 후육 고장력 열연 강판 및 그 제조 방법
KR101304859B1 (ko) * 2009-12-04 2013-09-05 주식회사 포스코 표면균열 저항성이 우수한 초고강도 라인파이프용 강판 및 그 제조방법
KR101528084B1 (ko) 2010-09-17 2015-06-10 제이에프이 스틸 가부시키가이샤 타발 가공성이 우수한 고강도 열연 강판 및 그 제조 방법
US20140205855A1 (en) * 2011-07-29 2014-07-24 Nippon Steel & Sumitomo Metal Corporation High-strength steel sheet excellent in impact resistance and manufacturing method thereof, and high-strength galvanized steel sheet and manufacturing method thereof
JP2016060955A (ja) * 2014-09-19 2016-04-25 株式会社神戸製鋼所 熱延鋼板及びその製造方法
KR20190077829A (ko) * 2017-12-26 2019-07-04 주식회사 포스코 고강도 고인성 열연강판 및 그 제조방법
KR20200062422A (ko) 2018-11-26 2020-06-04 주식회사 포스코 내구성이 우수한 고강도 강재 및 이의 제조방법
KR20210068808A (ko) 2019-12-02 2021-06-10 주식회사 포스코 내구성이 우수한 후물 복합조직강 및 그 제조방법

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