Steel plate and its manufacturing method

Abstract

The present invention relates to a high-strength steel plate and a manufacturing method thereof, and more particularly to a steel plate having high strength, excellent shear formability, and a low yield ratio, and a manufacturing method thereof.

Description

[Technical Field]

[0001] The present invention relates to a steel plate and a manufacturing method thereof, and more particularly to a high-strength steel plate with a low yield ratio and excellent shear formability, and a manufacturing method thereof. [Background technology]

[0002] Among conventional chassis parts, cross members and subframes, which require a relatively large amount of forming, have primarily used low-yield-ratio, high-strength hot-rolled steel sheets with a tensile strength of 540 to 590 MPa. Such low-yield-ratio hot-rolled steel sheets are ferrite-martensite dual-phase composite steels, and exhibit continuous yield behavior and low yield strength due to mobile dislocations introduced during martensitic transformation, resulting in excellent elongation and stretch formability.

[0003] In order to improve the elongation ratio and stretch formability, Patent Documents 1 to 3 utilize a method in which, after hot rolling, a Si-Mn or Mn-P-Cr alloy is used as the base material, and the temperature is maintained in the ferrite transformation region for several seconds, and then controlled to a temperature below the martensite transformation start temperature (Ms). Patent Document 4 similarly uses a Si-Mn-Cr or Si-Mn-Cr-Mo alloy, and the temperature is maintained in the ferrite transformation region for several seconds, and then coiled at a temperature above the martensite transformation start temperature. Patent Document 5 proposes a technique for suppressing the formation of coarse carbonitrides when Ti, Nb, V, etc. are added in order to obtain improved high-strength steel.

[0004] However, alloying elements such as Si, Al, Mn, Cr, and Mo, which are primarily used to produce high-strength hot-rolled steel sheets with low yield ratios and even higher strength ferrite-martensite dual-phase steels, cause severe segregation in the cast slab, resulting in the formation of cracks and defects during forming, degrading fatigue and impact resistance. Furthermore, excessive addition of these alloying elements increases hot deformation resistance, and when Ti, Nb, V, and W are added together, the rapid change in deformation resistance associated with dynamic deformation-induced precipitation during hot rolling leads to deterioration in the shape quality of the rolled sheet and a non-uniform microstructure, ultimately resulting in poor physical properties of the final part.

[0005] Furthermore, although high-strength hot-rolled steel sheets with low yield ratios utilize mobile dislocations formed at the boundaries between soft and hard phases in the microstructure, the above-mentioned technologies propose manufacturing methods that only consider the alloying elements and the ratios of each element, and therefore do not provide a solution to the problem of poor shear formability. Shear formability is the first step in the part forming process, and high-strength steels have a problem of cracks occurring during shear forming if the microstructure and elemental uniformity across the thickness are poor. These cracks can lead to serious cracks during part forming or adversely affect durability during use. [Prior art documents] [Patent documents]

[0006] [Patent Document 1] Japanese Patent Publication No. 1995-278731 [Patent Document 2] Japanese Patent Publication No. 1997-241790 [Patent Document 3] Japanese Patent Publication No. 1994-049591 [Patent Document 4] U.S. Patent Publication No. 4,502,897 [Patent Document 5] Korean Patent Registration No. 1543838 Summary of the Invention [Problem to be solved by the invention]

[0007] One embodiment of the present invention is to provide a steel sheet and a method for manufacturing the same.

[0008] An embodiment of the present invention provides a high-strength steel sheet with a low yield ratio and excellent shear formability, and a method for manufacturing the same.

[0009] The object of the present invention is not limited to the above-mentioned contents, and a person skilled in the art will have no difficulty in understanding further object of the present invention from the general contents of this specification. [Means for solving the problem]

[0010] According to one embodiment of the present invention, the alloy contains, in weight percent, C: 0.030 to 0.150%, Si: 0.01 to 1.00%, Mn: 1.00 to 2.50%, Al: 0.01 to 0.80%, Cr: 0.005 to 0.500%, Mo: 0.005 to 0.300%, P: 0.001 to 0.050%, S: 0.001 to 0.010%, N: 0.001 to 0.010%, the balance being Fe and unavoidable impurities, The value of X defined by the following relational expression 1 is 0.010 to 0.200, The T value defined by the following relational expression 2 is 1.500 to 4.200, The microstructure contains, by area percentage, 30 to 70% of a hard phase including bainite and martensite, 30 to 70% of a soft phase including ferrite, and 3% or less of pearlite. The average dislocation density of the hard phase is 2.0 to 3.0 × 10 14 m -2 and the average dislocation density of the soft phase is 0.50 to 2.00 × 10 14 m -2 It is possible to provide a steel sheet having the above formula: [Equation 1] X=([Nb] / 93+A / 48+[V] / 51) / ([C] / 12+[N] / 14) A=[Ti]-3.42[N]-1.5[S] (In the formula, [Nb], [V], [C], [N], [Ti], and [S] are the weight percentages of each element.) [Equation 2] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of each element.)

[0011] The steel plate may further contain, by weight percent, one or more selected from Nb: 0.005 to 0.03%, Ti: 0.005 to 0.120%, V: 0.005 to 0.200%, and B: 0.0003 to 0.0030%.

[0012] The steel plate may have a tensile strength of 780 MPa or more and a yield ratio of 0.70 to 0.85.

[0013] When the above steel sheets are punched with a punching clearance of 5 to 20%, microcracks of 0.1 mm or more in length occur on the shear surface at a rate of 10 / cm. 2 and the maximum crack length can be 1 mm or less.

[0014] According to one embodiment of the present invention, a method for manufacturing a steel slab includes, by weight, C: 0.030 to 0.150%, Si: 0.01 to 1.00%, Mn: 1.00 to 2.50%, Al: 0.01 to 0.80%, Cr: 0.005 to 0.500%, Mo: 0.005 to 0.300%, P: 0.001 to 0.050%, S: 0.001 to 0.010%, N: 0.001 to 0.010%, the balance being Fe and unavoidable impurities, and wherein an X value defined by the following relational expression 1 is 0.010 to 0.200, and a T value defined by the following relational expression 2 is 1.500 to 4.200; hot rolling the reheated steel slab; a step of primarily cooling the steel sheet produced in the hot rolling step to a temperature range of 430 to 600°C at an average cooling rate of 50 to 100°C / s; air-cooling the primarily cooled steel plate for 4.0 to 10.0 seconds; and Secondary cooling of the air-cooled steel sheet to a temperature range of 50 to 200°C at an average cooling rate of 10 to 100°C / s and coiling; During the primary cooling, edge portions corresponding to 30% regions from both ends to the other end in the width direction of the steel sheet are cooled to a surface temperature (TE) in the range of 500 to 600°C, and a central portion of a central 40% region corresponding to the region excluding both edge portions is cooled to a surface temperature (TC) in the range of 430 to 500°C. [Equation 1] X=([Nb] / 93+A / 48+[V] / 51) / ([C] / 12+[N] / 14) A=[Ti]-3.42[N]-1.5[S] (In the formula, [Nb], [V], [C], [N], [Ti], and [S] are the weight percentages of each element.) [Equation 2] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of each element.)

[0015] The steel slab may further contain, by weight percent, one or more selected from Nb: 0.005 to 0.030%, Ti: 0.005 to 0.120%, V: 0.005 to 0.200%, and B: 0.0003 to 0.0030%.

[0016] The reheating step is carried out in a temperature range of 1150 to 1350°C, The hot rolling step can be carried out at a finish rolling temperature of 850 to 1150°C.

[0017] The average temperature of the steel sheet after the air-cooling step may be 550 to 650°C. [Effects of the Invention]

[0018] According to one embodiment of the present invention, a steel sheet and a manufacturing method thereof can be provided.

[0019] According to one embodiment of the present invention, a high-strength steel sheet with a low yield ratio and excellent shear formability and a manufacturing method thereof can be provided.

[0020] According to one embodiment of the present invention, it is possible to provide a steel sheet that can be used for cross members and subframes, which are automobile chassis parts that require a relatively large amount of forming, and a manufacturing method thereof. [Brief explanation of the drawings]

[0021] [Figure 1] 10% and 20% punching clearances, respectively, for the invention example and the comparative example, showing the relationship between the number of cracks generated by each size on the shear surface. DETAILED DESCRIPTION OF THE INVENTION

[0022] Hereinafter, preferred embodiments of the present invention will be described. The embodiments of the present invention can be modified in various forms, and the scope of the present invention should not be construed as being limited to the embodiments described below. The embodiments are provided to further explain the present invention to those of ordinary skill in the art to which the present invention pertains.

[0023] The present invention overcomes the above-mentioned conventional problems by researching a method for preventing fractures at the cross sections of sheared and punched parts and preventing fatigue failure during product use. To this end, research was conducted to find a method for suppressing the formation of microcracks and facilitating product molding.

[0024] Therefore, the inventors of the present invention investigated the shear formability and microcrack formation on the shear surface depending on the characteristics of the components and microstructure of steels with various alloy compositions and microstructures. From the results, they found that excellent shear formability is achieved when the soft and hard phases constituting the microstructure each satisfy a specific range of dislocation density, rather than the composition ratio of the microstructure, and thus completed the present invention.

[0025] The present invention will be described in detail below.

[0026] The steel composition of the present invention will be described in detail below.

[0027] In the present invention, unless otherwise specified, the percentage representing the content of each element is based on weight.

[0028] A steel sheet according to one embodiment of the present invention may contain, in weight percent, C: 0.030 to 0.150%, Si: 0.01 to 1.00%, Mn: 1.00 to 2.50%, Al: 0.01 to 0.80%, Cr: 0.005 to 0.500%, Mo: 0.005 to 0.300%, P: 0.001 to 0.050%, S: 0.001 to 0.010%, N: 0.001 to 0.010%, the balance being Fe and unavoidable impurities.

[0029] Carbon (C):0.030~0.150% Carbon (C) is the most economical and effective element for strengthening steel, significantly affecting the hardness and dislocation density of each component. Increasing its content increases hardening ability and the fraction of hard phases such as bainite and martensite in the microstructure, thereby increasing dislocation density and tensile strength. It also forms fine precipitates with Ti and Nb, which have a high affinity for carbon (C), resulting in finer grains and enhanced precipitation strengthening, resulting in increased yield strength and tensile strength. If the carbon (C) content is less than 0.030%, it may be difficult to achieve sufficient strengthening. According to one embodiment of the present invention, carbon may be present in an amount of 0.050% or more to stably ensure a higher level of strength. On the other hand, if the carbon content exceeds 0.150%, the fraction of each phase, including bainite and martensite, increases, and the hardness of the phases also increases, resulting in an excessive increase in strength, which can cause problems such as reduced elongation and formability, and the weldability may also deteriorate. In the present invention, in order to ensure more stable formability, carbon (C) can be contained in an amount of 0.120% or less.

[0030] Silicon (Si): 0.01 to 1.00% Silicon (Si) has the effect of deoxidizing molten steel and solid-solution strengthening, and it also delays the formation of coarse carbides, which is advantageous for improving formability. In the present invention, silicon (Si) can be contained in an amount of 0.01% or more to achieve the above-mentioned effects. According to one embodiment of the present invention, silicon (Si) can be contained in an amount of 0.10% or more. However, if the silicon content exceeds 1.00%, red scale due to silicon (Si) may form on the steel sheet surface during hot rolling, significantly deteriorating the surface quality of the steel sheet and potentially reducing ductility and weldability. In one embodiment of the present invention, silicon can be contained in an amount of 0.90% or less.

[0031] Manganese (Mn): 1.00-2.50% Like Si, manganese (Mn) is an effective element for solid-solution strengthening of steel. It increases the hardenability of steel and facilitates the formation of hard phases, bainite and martensite, during cooling after hot rolling. However, if its content is less than 1.00%, the above effects cannot be achieved. According to one embodiment of the present invention, manganese (Mn) can be contained in an amount of 1.40% or more. On the other hand, if its content exceeds 2.50%, the hardenability increases significantly, and the fraction and hardness of each phase, including bainite and martensite, increase. This can lead to problems such as excessive strength increase and reduced formability. Furthermore, during slab casting in the continuous casting process, segregation can develop significantly in the center of the thickness, which can lead to uneven formation of a microstructure in the thickness direction during cooling after hot rolling, resulting in poor stretch flangeability. In particular, it can be difficult to produce a uniform microstructure throughout the entire length and width of the hot-rolled sheet during cooling. In one embodiment of the present invention, the upper limit of the manganese (Mn) content can be limited to 2.30%.

[0032] Aluminum (Al): 0.01 to 0.80% Aluminum (Al) is a component added primarily for deoxidation and has the effect of promoting ferrite transformation. If its content is less than 0.01%, the above-mentioned effect of addition may be insufficient. According to one embodiment of the present invention, aluminum (Al) may be contained in an amount of 0.02% or more. On the other hand, if its content exceeds 0.80%, it bonds with N to form AlN, which is likely to cause corner cracks in the slab during continuous casting and defects due to the formation of inclusions. According to one embodiment of the present invention, the upper limit of the aluminum (Al) content may be limited to 0.50%.

[0033] Chromium (Cr): 0.005 to 0.500% Chromium (Cr) strengthens steel through solid solution and retards the ferrite phase transformation during cooling, promoting the formation of bainite. However, if the chromium (Cr) content is less than 0.005%, the above effects cannot be achieved. To ensure these effects more effectively, 0.100% or more of Cr may be included in the present invention. On the other hand, if the Cr content exceeds 0.500%, the ferrite transformation is excessively delayed, resulting in a deterioration in elongation due to the formation of martensite. Furthermore, similar to Mn, large segregations may develop in the center of the steel, resulting in a non-uniform microstructure in the thickness direction and a deterioration in stretch flangeability. In one embodiment of the present invention, the upper limit of Cr may be limited to 0.300%.

[0034] Molybdenum (Mo): 0.005 to 0.300% Molybdenum (Mo) increases the hardenability of steel and facilitates the formation of bainite. However, if its content is less than 0.005%, the above-mentioned effects of addition cannot be obtained. In one embodiment of the present invention, its content can be 0.050% or more. On the other hand, if its content exceeds 0.300%, an excessive increase in hardenability may cause martensite to form, which may rapidly deteriorate formability. This is also economically disadvantageous and may also cause disadvantages in weldability. In one embodiment of the present invention, its upper limit can be limited to 0.200%.

[0035] Phosphorus (P): 0.001 to 0.050% Like Si, phosphorus (P) simultaneously promotes solid solution strengthening and ferrite transformation. However, manufacturing with a content of less than 0.001% is economically disadvantageous due to the high manufacturing costs and is insufficient to obtain sufficient strength, so the lower limit can be set to 0.001%. On the other hand, if the phosphorus (P) content exceeds 0.050%, embrittlement occurs due to grain boundary segregation, which makes it easy for fine cracks to form during forming, and this can significantly deteriorate ductility, stretch flangeability, and impact resistance.

[0036] Sulfur (S): 0.001 to 0.010% Sulfur (S) is an impurity present in steel. If its content exceeds 0.010%, it combines with Mn and other elements to form nonmetallic inclusions, which can easily cause fine cracks during cutting of the steel and significantly reduce stretch flangeability and impact resistance. In the present invention, sulfur (S) can be contained in an amount of 0.005% or less. While the present invention does not particularly limit the lower limit of the sulfur (S) content, reducing the sulfur (S) content to less than 0.001% requires a long time during steelmaking operations, thereby reducing productivity. Taking this into consideration, the lower limit of the sulfur (S) content can be set to 0.001%.

[0037] Nitrogen (N): 0.001 to 0.010% Nitrogen (N), along with carbon, is a typical solid-solution strengthening element, forming coarse precipitates along with titanium and aluminum. While nitrogen (N) generally has a stronger solid-solution strengthening effect than carbon, the toughness of steel decreases significantly as the amount of nitrogen (N) in the steel increases. Furthermore, reducing the nitrogen (N) content to less than 0.001% requires a long time during steelmaking operations, reducing productivity. Therefore, the lower limit can be set to 0.001%.

[0038] The steel material of the present invention may contain the remaining iron (Fe) and inevitable impurities in addition to the above-mentioned composition. The inevitable impurities cannot be excluded because they may be unintentionally mixed in during the normal manufacturing process. Since such impurities are known to any engineer in the field of normal steel manufacturing, the details of all of them will not be specifically mentioned in this specification.

[0039] The steel sheet according to one embodiment of the present invention may further contain one or more selected from Nb: 0.005 to 0.030%, Ti: 0.005 to 0.120%, V: 0.005 to 0.200%, and B: 0.0003 to 0.0030%.

[0040] Niobium (Nb): 0.005 to 0.030% Niobium (Nb), along with Ti and V, is a typical precipitation strengthening element. It precipitates during hot rolling and has the effect of refining crystal grains by delaying recrystallization, thereby effectively improving the strength and impact toughness of steel. If the niobium (Nb) content is less than 0.005%, the above effect cannot be obtained. On the other hand, if the niobium (Nb) content exceeds 0.030%, there is a problem that stretch flangeability is deteriorated due to the formation of elongated crystal grains and coarse complex precipitates associated with excessive delay in recrystallization during hot rolling.

[0041] Titanium (Ti): 0.005 to 0.120% Titanium (Ti), along with Nb and V, is a typical precipitation strengthening element. Its strong affinity with N allows it to form coarse TiN in steel. TiN has the effect of suppressing grain growth during the heating process for hot rolling. Furthermore, titanium (Ti) remaining after reacting with nitrogen dissolves in the steel and bonds with carbon to form TiC precipitates, making it a useful component for improving the strength of steel. If the titanium (Ti) content is less than 0.005%, the above effects cannot be achieved. On the other hand, if the titanium (Ti) content exceeds 0.120%, the generation of coarse TiN and the coarsening of the precipitates can lead to problems such as poor stretch flangeability during forming. In one embodiment of the present invention, the upper limit may be 0.105%. In another embodiment of the present invention, the upper limit may be 0.100%.

[0042] Vanadium (V): 0.005 to 0.200% Vanadium (V), along with Nb and Ti, is a typical precipitation strengthening element. It hardly precipitates during hot rolling, but forms precipitates after coiling, improving the strength of steel. Therefore, it is effective in further improving strength without increasing deformation resistance and rolling load due to delayed recrystallization during hot rolling. To achieve this effect, the vanadium (V) content in the present invention can be 0.005% or more. However, excessive vanadium (V) content can lead to problems such as poor stretch flangeability due to the formation of coarse precipitates, which is economically disadvantageous. Therefore, in the present invention, the upper limit can be limited to 0.200%. In one embodiment, the upper limit can be limited to 0.150%.

[0043] Boron (B): 0.0003 to 0.0030% When present in a solid solution state in steel, boron (B) segregates mainly at grain boundaries, stabilizing the grain boundaries and improving the brittleness of the steel. It also stabilizes solute N and suppresses the formation of coarse AlN nitrides. It also delays the ferrite phase transformation and effectively promotes the formation of hard phases, bainite and martensite. In the present invention, boron (B) may be contained in an amount of 0.0003% or more to ensure the above-mentioned effects. However, if the boron (B) content exceeds 0.0030%, the effect of the addition is not further increased, and there is a problem in that ductility is reduced and formability is deteriorated. According to one embodiment of the present invention, boron (B) may be contained in an amount of 0.0020% or less.

[0044] In a steel sheet according to an embodiment of the present invention, the X value defined by the following relational expression 1 may be 0.010 to 0.200, and the T value defined by the following relational expression 2 may be 1.500 to 4.200.

[0045] [Equation 1] X=([Nb] / 93+A / 48+[V] / 51) / ([C] / 12+[N] / 14) A=[Ti]-3.42[N]-1.5[S] (In the formula, [Nb], [V], [C], [N], [Ti], and [S] are the weight percentages of each element.) [Equation 2] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of each element.)

[0046] When the X value defined in Relational Formula 1 exceeds 0.200, the formation of precipitates increases, resulting in a slight increase in strength. However, the delayed recrystallization during hot rolling may lead to the formation of a microstructure elongated in the rolling direction, potentially reducing the elongation rate in the direction perpendicular to the rolling direction. Furthermore, during cooling of the hot-rolled steel sheet, the amount of solute C and N atoms in the untransformed phase may be insufficient, making it difficult to stably form a hard phase. This may result in brittle grain boundaries and poor shear surface quality. In one embodiment of the present invention, the X value may be 0.180 or less. On the other hand, if the X value is less than 0.010, grain growth during reheating is facilitated, recrystallization during hot rolling becomes nonuniform, resulting in the formation of locally coarse grains. This may result in an excessive amount of solute C and N, which may increase the hardness of the hard phase and ultimately result in poor elongation. In one embodiment of the present invention, the X value may be 0.030 or more. On the other hand, if the alloy element in Relational Formula 1 is not added, 0 may be substituted.

[0047] Relational Formula 2 factors the combination of alloying elements that can maintain the formation of the hard phases bainite, martensite, and MA phases in the microstructure at an appropriate level. The larger the T value defined by Relational Formula 2, the greater the formation of the hard phases bainite, martensite, and MA phases, and the greater the hardness of each hard phase. Therefore, in the present invention, the T value can be limited to 1.500 or more to achieve the desired strength. According to one embodiment of the present invention, the T value can be limited to 2.000 or more. While a higher T value is advantageous for ensuring strength, an excessive T value can reduce the ductility of the steel and unnecessarily increase the hardness difference between the soft and hard phases, potentially resulting in poor shear formability. Furthermore, there is a problem of increased variation in the quality of the hot-rolled steel sheet over its entire length and width. Therefore, in the present invention, the upper limit of the T value can be limited to 4.200. According to one embodiment of the present invention, the upper limit of the T value can be limited to 4.000.

[0048] The microstructure of the steel of the present invention will now be described in detail.

[0049] In the present invention, unless otherwise specified, the percentage representing the fraction of the microstructure is based on the area.

[0050] The inventors of the present invention have found that it is difficult to clearly distinguish, based on the area ratio of the microstructure alone, how excellent the inherent shear formability of a steel material is and whether stable results are obtained even with variations in punching clearance. In particular, they have confirmed that the dislocation density and physical characteristics vary greatly depending on the components that make up the steel.

[0051] Therefore, the inventors have confirmed through their research that the dislocation density in the microstructure (Geometrical Necessary Dislocation) is an important factor influencing the occurrence of microcracks, which is the quality of the shear surface of steel, and have proposed the present invention.

[0052] The microstructure of a steel sheet according to one embodiment of the present invention may contain, in area percentages, 30 to 70% of a hard phase including bainite and martensite, 30 to 70% of a soft phase including ferrite, and 3% or less of pearlite.

[0053] In the present invention, the microstructure is controlled to ensure a low yield ratio and stretch formability, and the area fractions of bainite and martensite are classified as hard phases, and ferrite as soft phases, and can be limited.

[0054] In the present invention, bainite may include upper bainite and lower bainite, and can be distinguished from a ferrite-based low-temperature transformation phase in that fine carbides are formed in a lath-shaped structure.

[0055] In the present invention, ferrite may include equiaxed ferrite and a ferrite-based low-temperature transformation phase. The ferrite-based low-temperature transformation phase may include acicular ferrite, bainitic ferrite, granular bainitic ferrite, etc., and may refer to ferrite having non-uniform grain boundaries, a high dislocation density within the grains, and a high structure density of low-angle grain boundaries within the grains compared to equiaxed ferrite.

[0056] The microstructure in the present invention can be observed in a cross section perpendicular to the rolling direction of the steel sheet, and can be analyzed at a point between 1 / 4 and 1 / 2t (t is the thickness of the steel sheet) in the thickness direction. The division and area fraction of the microstructure can be measured using electron backscattered diffraction (EBSD, (JEOL JSM-1001F)) at a magnification of 3000 to 5000.

[0057] If the area fraction of the hard phase exceeds 70%, the elongation rate will be significantly reduced, the proportion of fractures on the shear surface will increase, and the number of cracks with a length of 1 mm or more may increase significantly. Furthermore, the shear surface quality will be more dependent on changes in punching clearance, which may increase the occurrence of defects during actual part forming. Meanwhile, if the area fraction is less than 30%, it may be difficult to achieve the target strength. Meanwhile, according to one embodiment of the present invention, since martensite is a relatively hard phase compared to bainite, an increase in martensite may result in a decrease in ductility. Taking this into consideration, the upper limit of the area fraction of martensite may be limited to 60%. Meanwhile, according to one embodiment of the present invention, the martensite content in the hard phase may be 0%.

[0058] In the present invention, a small amount of MA phase may be observed, but since the effects on the physical properties proposed in the present invention and the cross-sectional quality of the punched portion are not particularly different, it is considered to be martensite. Furthermore, tempered martensite containing fine carbides has a lower dislocation density than martensite, which reduces the hardness value of the phase and helps improve the ductility of the entire steel. However, if the size of the formed carbides increases, it can have adverse effects such as causing brittleness, so it can be distinguished from regular martensite. However, in the present invention, since the dislocation density of the phase is measured and described together, there is no need to distinguish it from tempered martensite, so it is considered to be martensite.

[0059] The soft phase contributes to the ductility of the steel and the formation of fine precipitates, so the area fraction of the soft phase can be limited to a lower limit of 30% and, according to one embodiment of the present invention, to a higher limit of 70%.

[0060] Meanwhile, pearlite may be further included as a structure other than the hard and soft phases, but if the area fraction of pearlite exceeds 3%, it may become a brittle structure during shear forming of the steel, and cracks with a length of 1 mm or more may increase, so the upper limit of the area fraction of pearlite may be limited to 3%.

[0061] In the steel sheet according to one embodiment of the present invention, the average dislocation density (Geometrical Necessary Dislocation) of the hard phase is 2.0 to 3.0 × 10 14 m -2 The average dislocation density of the soft phase can be 0.50 to 2.00 × 10 14 m -2 It can be.

[0062] The average dislocation density (Geometrical Necessary Dislocation) can be calculated using kernel average misorientation (KAM) data after measuring a cross section parallel to the rolling direction at the 1 / 4 point in the thickness direction of the steel sheet using EBSD, and can be calculated using the following equation. For convenience, this calculation can be performed using OIM analysis software, which analyzes the EBSD measurement results. TM (EDAX) etc. can be used.

[0063] [formula] Dislocation density (GND, m -2 ) = 2θ / ub (In the formula, θ is the average misorientation (KAM values), u is the unit length (step size in the EBSD measurement), and b is the burgers vector.)

[0064] The average dislocation density of the hard phase is 2.0×10 14 m -2 If the value is less than 3.0×10, the strength may be significantly reduced. 14 m -2 If the shear strength exceeds 100%, the ductility may decrease and the quality of the shear surface may deteriorate.

[0065] The average dislocation density of the soft phase is 0.50×10 14 m -2 If the value is less than 2.00×10, the strength may not reach the level required by the present invention, and burrs may occur severely during shear forming.14 m -2 If the temperature exceeds this range, the yield strength increases and the elongation rate decreases.

[0066] The method for producing a steel sheet according to the present invention will be described in detail below.

[0067] A steel sheet according to one embodiment of the present invention can be produced by reheating a steel slab satisfying the alloy composition of the present invention, hot rolling it, primary cooling it, air cooling it, secondary cooling it, and coiling it.

[0068] reheating A steel slab satisfying the alloy composition of the present invention can be reheated to a temperature range of 1150 to 1350°C.

[0069] If the reheating temperature is less than 1150°C, the precipitates will not fully redissolve, the formation of precipitates will decrease in processes after hot rolling, coarse TiN will remain, and the steel slab will not be sufficiently heated, which may make it difficult to control the temperature of the steel sheet constant during hot rolling.On the other hand, if the temperature exceeds 1350°C, abnormal grain growth of austenite grains may occur, resulting in a decrease in strength.

[0070] hot rolling The reheated steel slab can be hot rolled at a finish rolling temperature of 850 to 1150°C.

[0071] If the finish rolling temperature during hot rolling exceeds 1150°C, the temperature of the hot-rolled steel sheet will become too high, causing the grain size to become coarse and potentially deteriorating the surface quality of the hot-rolled steel sheet.On the other hand, if the temperature is less than 850°C, excessive delay in recrystallization will cause the development of elongated grains, resulting in severe anisotropy and potentially poor formability.

[0072] Primary cooling The steel sheet produced in the hot rolling step may be primarily cooled to a temperature range of 430 to 600°C at an average cooling rate of 50 to 100°C / s. During the primary cooling, the edge portions corresponding to 30% of the area from both ends to the other end in the width direction of the steel sheet may be cooled to a surface temperature (TE) range of 500 to 600°C, and the center portion of the central 40% area corresponding to the area excluding both edge portions may be cooled to a surface temperature (TC) range of 430 to 500°C.

[0073] In the present invention, the edge portion is defined as an area corresponding to 30% of the width of the steel plate from both ends toward the other end or the center, i.e., an area corresponding to a total of 60% of the entire steel plate, and the central 40% area excluding the edge portion is defined as the center portion.

[0074] In the present invention, the soft phase in the microstructure of the steel is formed during primary cooling and air cooling, and the hard phase may be formed during secondary cooling and coiling. However, because the untransformed phase immediately before secondary cooling transforms into the hard phase after secondary cooling, the soft phase must be uniformly formed across the width of the steel sheet immediately after primary cooling and air cooling. Typically, during cooling of a steel sheet, heat transfer proceeds rapidly at the edge of the steel sheet, resulting in the formation of more hard phases than at the center. Therefore, in order to achieve a uniform cooling rate across the width of the steel sheet, the present invention attempts to control the cooling end temperatures of the center and edge portions differently. If the soft phase is uniformly formed across the width during primary cooling and air cooling as proposed in the present invention, the ratio of the hard phase formed during secondary cooling will be constant, thereby reducing dependency on cooling conditions after coiling. If the soft phase and hard phase are uniformly formed, the excellent material uniformity and shear formability desired in the present invention can be ensured.

[0075] If the cooling rate during primary cooling is less than 50°C / s, an excessively large ferrite fraction may be formed, and the average dislocation density may fall below the target level, which may make it difficult to ensure strength.On the other hand, if the cooling rate exceeds 100°C / s, the surface temperature (TC) at the center during primary cooling may become excessively low, significantly reducing the ferrite fraction, and the hard phase may increase more than necessary, resulting in insufficient elongation.

[0076] Meanwhile, in the present invention, since there is a risk of a large difference in cooling rate depending on the position in the width of the steel sheet, in order to prevent the hard phase from being formed differently depending on the position in the width, the edge portion and the center portion may be divided and cooled to different temperature ranges.

[0077] In particular, the present invention aims to have the temperature of the steel sheet recover during the subsequent air-cooling process, thereby maintaining a uniform temperature of 550 to 650°C. Therefore, the temperature of the center needs to be cooled below the target temperature of 550 to 650°C, and it is preferable that the edge portions, which have a higher cooling rate, are cooled to a higher temperature range than the center portion. Therefore, in the present invention, the surface temperature TE of the edge portions can be cooled to 500 to 600°C, and the surface temperature (TC) of the center portion can be cooled to 430 to 500°C.

[0078] If the surface temperature of the edge portion is less than 500°C, the formation of soft phases may be insufficient, which may result in the hot-rolled steel sheet not reaching the desired temperature range during air-cooling. On the other hand, if the temperature exceeds 600°C, the desired temperature range may be exceeded during air-cooling, which may result in the formation of excessive soft phases in the final microstructure. According to an embodiment of the present invention, the surface temperature of the edge portion may be 510°C or higher during primary cooling.

[0079] If the surface temperature of the center is below 430°C, a phase transformation of the hard phase, bainite, occurs. On the other hand, if the temperature exceeds 500°C, the effect of recuperation decreases, and the target temperature range may not be reached.

[0080] air-cooled The steel sheet that has been primarily cooled can be air-cooled for 4.0 to 10.0 seconds.

[0081] The intentional cooling of the primarily cooled steel sheet can be stopped, and the temperature of the steel sheet can be restored to a target temperature by internal latent heat and transformation heat. When the primarily cooled steel sheet is air-cooled, the steel sheet can be restored to an average temperature of 550 to 650°C.

[0082] If the air-cooling time is less than 4.0 seconds, the reheating effect may not be achieved. On the other hand, if the time exceeds 10.0 seconds, the ferrite fraction in the microstructure may increase significantly, and the hard phases of bainite and martensite may decrease. Furthermore, pearlite and coarse carbides may form in the center of the thickness of the steel sheet where the temperature is high, which may result in a deterioration in the cross-section quality after shear forming.

[0083] Secondary cooling and winding The air-cooled steel sheet can be subjected to secondary cooling to a temperature range of 50 to 200°C at an average cooling rate of 10 to 100°C / s and then coiled.

[0084] If the coiling temperature exceeds 200°C, the average dislocation density of the hard phase will be outside the range proposed by the present invention, and it may be difficult to ensure strength. According to one embodiment of the present invention, the coiling temperature can be limited to 150°C or less. On the other hand, if the temperature is less than 50°C, an unnecessarily large amount of martensite is formed, the average dislocation density of the hard phase will exceed the range proposed by the present invention, and the elongation rate of the steel will deteriorate, and problems may arise such that cooling water will remain and corrode the steel sheet. According to one embodiment of the present invention, the coiling temperature can be limited to 70°C or more.

[0085] If the cooling rate exceeds 100°C / s, the average dislocation density of the hard phase becomes excessively high, which may result in a problem of a decrease in elongation. Although there is no particular limit to the cooling rate, there is a problem in that the length of the cooling zone equipment must be increased in order to control the cooling rate to less than 10°C / s, which makes it difficult to produce at the target coiling temperature of 200°C or less.

[0086] The steel sheet manufactured in this way has a tensile strength of 780 MPa or more, a yield ratio of 0.70 to 0.85, and when punched with a punching clearance of 5 to 20%, the number of microcracks with a length of 0.1 mm or more on the shear surface is 10 / cm. 2 The maximum crack length is 1 mm or less, and the strength and shear formability are excellent, and the properties of a low yield ratio can be ensured. [Example]

[0087] The present invention will be described in more detail with reference to the following examples. However, it should be noted that the following examples are intended to illustrate and explain the present invention in more detail, and are not intended to limit the scope of the present invention.

[0088] (Example) Steel slabs having the composition shown in Table 1 below were used to produce steel plates under the conditions shown in Table 2 below. The reheating temperature not disclosed in Table 2 below was 1250°C, and the thickness of the steel plates immediately after hot rolling was also 3.2 mm.

[0089] [Table 1]

[0090] [Equation 1] X=([Nb] / 93+A / 48+[V] / 51) / ([C] / 12+[N] / 14) A=[Ti]-3.42[N]-1.5[S] (In the formula, [Nb], [V], [C], [N], [Ti], and [S] are the weight percentages of each element.) [Equation 2] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of each element.)

[0091] [Table 2]

[0092] Tables 3 and 4 below show the microstructure and mechanical properties of the steel sheets manufactured. The microstructure was analyzed by observing a cross section perpendicular to the rolling direction of the steel sheet at the 1 / 4 to 1 / 2 point in the thickness direction. The classification and area fraction of ferrite, ferritic low-temperature transformation-generated phases, bainite, martensite, and pearlite formed in the steel were measured using electron backscattered diffraction (EBSD) (JEOL JSM-1001F) at 3000 to 5000 magnifications. Additionally, the average dislocation density (Geometrical Necessary Dislocation, GND) was measured using OIM analysis after EBSD measurement on a cross section parallel to the rolling direction at the 1 / 4 point in the thickness direction of the steel sheet. TM (EDAX) was used for measurement.

[0093] The mechanical properties were measured and shown as yield strength, tensile strength, elongation at break, and yield ratio. Here, the 0.2% offset yield strength (YS), tensile strength (TS), and elongation at break (T-El) were measured using JIS No. 5 standard test pieces taken perpendicular to the rolling direction. The physical properties were also shown by measuring the number of cracks by length on the punched cross section. These were evaluated by punching holes with a diameter of 10 mm. Punching was performed with different punching clearances of 5, 10, and 20%, and the average number of microcracks observed on cross sections parallel to and perpendicular to the rolling direction was calculated. Each result represents the number of cracks by length.

[0094] [Table 3]

[0095] *F: equiaxed ferrite, BF: ferrite low-temperature transformation phase, B: bainite, M: martensite, P: pearlite

[0096] [Table 4]

[0097] As shown in Tables 3 and 4, in the case of the invention examples that satisfy the alloy composition and manufacturing conditions of the present invention, the microstructure characteristics proposed in the present invention are satisfied, and the physical properties targeted by the present invention are also ensured.

[0098] 1(a) and (b) show the relationship between the number of cracks by size on the shear surface of the inventive example and the comparative example at 10% and 20% punching clearance, respectively. Specifically, it can be seen that at 10% and 20% punching clearance, the comparative example had a higher number of cracks than the inventive example, and that no cracks exceeding 1.0 mm in size occurred in the inventive example, and the number of cracks smaller than 1.0 mm was significantly lower than in the comparative example.

[0099] On the other hand, Comparative Example 1 is an example in which the alloy element content range proposed in the present invention is satisfied, but the condition of Relational Formula 1 is not satisfied. As a result, the average dislocation density of the soft phase exceeds the range proposed in the present invention, which is believed to be due to the increased formation of fine precipitates in the soft phase. Although no coarse cracks exceeding 1 mm in length were observed in the punched section, when the clearance was 20%, the occurrence of cracks measuring 0.1 to 1.0 mm in length significantly increased. In addition, the yield ratio was high, and work hardening during forming excessively increased the yield strength, which also caused the problem of deteriorated formability.

[0100] Comparative Examples 2, 3, 12, and 13 are examples that do not satisfy Relational Formula 2. Comparative Examples 2 and 12 contained excessive amounts of alloy components with high hardening effects, and although strength was stably ensured, the elongation rate was insufficient. This also deteriorated the quality of the shear surface. Comparative Examples 3 and 13 contained insufficient alloy components with excellent hardening effects, and the hard phase was not formed to the level proposed in the present invention. As a result, the target strength could not be ensured. Furthermore, as the clearance increased, cracking became more severe, and cracks longer than 1 mm were observed.

[0101] In Comparative Examples 4 and 5, the cooling end temperature during primary cooling after hot rolling was outside the range proposed in the present invention. In Comparative Example 4, the end temperature during primary cooling exceeded the upper limit standard, resulting in insufficient formation of hard phases and the formation of unnecessary pearlite, resulting in poor cross-sectional quality after punching. In Comparative Example 5, the end temperature during primary cooling was outside the lower limit standard, resulting in an insufficient fraction of soft phases and the formation of excessive hard phases, resulting in poor cross-sectional quality after punching.

[0102] In Comparative Examples 6 and 7, the air-cooling time after primary cooling was outside the range of the present invention. The long exposure time in the high-temperature range resulted in a large increase in the soft phase fraction due to latent heat inside the steel sheet and heat generation due to phase transformation, and pearlite was also formed, resulting in poor cross-section quality after punching and a yield ratio that exceeded the range proposed by the present invention. In particular, in the case of Comparative Example 7, it can be confirmed that the average temperature of the steel sheet after air-cooling exceeded the temperature range proposed by the present invention.

[0103] Comparative Example 8 is an example in which the air-cooling time after primary cooling did not fall within the range of the present invention. Secondary cooling was performed before the steel sheet was reheated, so the soft phase fraction did not reach the desired level and the hard phase fraction exceeded the desired range. As a result, the cross-section quality after punching was deteriorated.

[0104] In Comparative Example 9, the cooling rate during secondary cooling was excessively fast, resulting in overcooling and not meeting the coiling temperature range targeted by the present invention. As a result, the dislocation density of the hard phase exceeded the proposed range, and the cross-section quality after punching deteriorated. The main cause is believed to be an increased difference in physical properties between the soft and hard phases.

[0105] Comparative Examples 10 and 11 are cases where the cooling end temperature during secondary cooling was outside the range proposed in the present invention. In Comparative Example 10, the coiling temperature after secondary cooling did not reach the proposed temperature range, and the dislocation density of the hard phase was excessively high, which resulted in poor cross-section quality after punching. In Comparative Example 11, the coiling temperature after secondary cooling exceeded the proposed temperature range, and the dislocation density of the hard phase did not reach the proposed level. As a result, the yield ratio was excessively high and the cross-section quality was also poor.

[0106] Although the present invention has been described in detail with reference to the above embodiments, other embodiments are possible, and the spirit and scope of the claims set forth below should not be limited to the embodiments.

Claims

1. The composition contains, in weight percent, C: 0.030 to 0.150%, Si: 0.01 to 1.00%, Mn: 1.00 to 2.50%, Al: 0.01 to 0.80%, Cr: 0.005 to 0.500%, Mo: 0.005 to 0.300%, P: 0.001 to 0.050%, S: 0.001 to 0.010%, N: 0.001 to 0.010%, the balance being Fe and inevitable impurities, The X value defined by the following relational expression 1 is 0.010 to 0.200, The T value defined by the following relational expression 2 is 1.500 to 4.200, The microstructure contains, by area percentage, 30 to 70% of a hard phase including bainite and martensite, 30 to 70% of a soft phase including ferrite, and 3% or less of pearlite; The average dislocation density of the hard phase is 2.0 to 3.0 × 10 １４ m －２ and the average dislocation density of the soft phase is 0.50 to 2.00 × 10 １４ m －２ That is, steel plate. [Relationship 1] X=([Nb] / 93+A / 48+[V] / 51) / ([C] / 12+[N] / 14) A=[Ti]-3.42[N]-1.5[S] (In the formula, [Nb], [V], [C], [N], [Ti] and [S] are the weight percentages of each element.) [Relationship 2] T=[Mn]+2.8[Mo]+1.5[Cr]+500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of each element.)

2. The steel plate according to claim 1, further comprising, by weight%, one or more selected from Nb: 0.005 to 0.030%, Ti: 0.005 to 0.120%, V: 0.005 to 0.200%, and B: 0.0003 to 0.0030%.

3. The steel plate according to claim 1, wherein the steel plate has a tensile strength of 780 MPa or more and a yield ratio of 0.70 to 0.

85.

4. When the steel plate is punched with a punching clearance of 5 to 20%, the number of microcracks having a length of 0.1 mm or more on the shear surface is 10 / cm. ２ 2. The steel plate according to claim 1, wherein the maximum crack length is 1 mm or less.

5. reheating a steel slab containing, by weight, C: 0.030 to 0.150%, Si: 0.01 to 1.00%, Mn: 1.00 to 2.50%, Al: 0.01 to 0.80%, Cr: 0.005 to 0.500%, Mo: 0.005 to 0.300%, P: 0.001 to 0.050%, S: 0.001 to 0.010%, N: 0.001 to 0.010%, the balance being Fe and inevitable impurities, and having an X value defined by the following relational expression 1 of 0.010 to 0.200 and a T value defined by the following relational expression 2 of 1.500 to 4.200; hot rolling the reheated steel slab; a step of primarily cooling the steel sheet produced in the hot rolling step to a temperature range of 430 to 600°C at an average cooling rate of 50 to 100°C / s; air-cooling the primarily cooled steel plate for 4.0 to 10.0 seconds; and secondary cooling of the air-cooled steel sheet to a temperature range of 50 to 200°C at an average cooling rate of 10 to 100°C / s and coiling; During the primary cooling, edge portions corresponding to 30% regions from both ends to the other end in the width direction of the steel plate are cooled to a surface temperature (TE) in the range of 500 to 600°C, and a central portion of a central 40% region corresponding to the region excluding both edge portions is cooled to a surface temperature (TC) in the range of 430 to 500°C. [Relationship 1] X=([Nb] / 93+A / 48+[V] / 51) / ([C] / 12+[N] / 14) A=[Ti]-3.42[N]-1.5[S] (In the formula, [Nb], [V], [C], [N], [Ti] and [S] are the weight percentages of each element.) [Relationship 2] T=[Mn]+2.8[Mo]+1.5[Cr]+500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of each element.)

6. The steel slab further contains, by weight%, one or more selected from Nb: 0.005 to 0.030%, Ti: 0.005 to 0.120%, V: 0.005 to 0.200%, and B: 0.0003 to 0.0030%. The method for producing a steel plate according to claim 5.

7. The reheating step is carried out at a temperature in the range of 1150 to 1350°C, The method for manufacturing a steel sheet according to claim 5, wherein the hot rolling step is performed at a finish rolling temperature of 850 to 1150°C.

8. The method for producing a steel plate according to claim 5, wherein the average temperature of the steel plate after the air-cooling step is 550 to 650°C.

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