Steel sheet and method for manufacturing same

Abstract

The present invention pertains to a high-strength steel sheet and a method for manufacturing same. More specifically, the present invention pertains to: a steel sheet having high strength, excellent shear formability, and a low yield ratio; and a method for manufacturing same.

Description

Technical Field

[0001] The present disclosure relates to a steel sheet and a method for manufacturing the steel sheet, and more specifically, to a high-strength steel sheet with excellent shear formability and a method for manufacturing the same.Background Art

[0002] Among conventional chassis components, high-strength hot-rolled steel sheets with a tensile strength of 540 to 590 MPa were mainly used for cross members and subframes having relatively large forming volumes. Such hot-rolled steel sheets having a low yield ratio are dual-phase composite structures of ferrite and martensite, and exhibit continuous yield behavior and low yield strength characteristics due to the movable dislocation introduced during martensite transformation, and have excellent elongation and elongation formability.

[0003] In order to improve an elongation rate and elongation formability, patent Documents 1 to 3 utilize a method of controlling a temperature to being below the martensite transformation initiation temperature (Ms) after performing a hot-rolling process based on Si-Mn, Mn-P-Cr component systems and maintaining the temperature in a ferrite transformation region for several seconds. Additionally, Patent Document 4 uses a method of maintaining a temperature in a ferrite transformation region for several seconds using a Si-Mn-Cr or Si-Mn-Cr-Mo system and then performing coiling at a temperature higher than or equal to a martensite transformation initiation temperature. Additionally, Patent Document 5 proposes a technique for applying a manner of suppressing the formation of coarse carbonitrides when Ti, Nb, V, and the like, are added, to obtain improved high-strength steel.

[0004] However, in high-strength hot-rolled steel sheets with a low yield ratio, alloying elements such as Si, Al, Mn, Cr, and Mo, which are mainly used to manufacture ferrite-martensite dual-phase composite structure steel with higher strength, cause severe segregation in a slab after casting, which may lead to the formation of cracks or defects during forming, and may worsen the fatigue resistance and impact resistance. Additionally, when the alloying elements are excessively added, the hot deformation resistance increases, and when Ti, Nb, V, and W are added together, the deformation resistance changes rapidly due to dynamic deformation-induced precipitation during hot rolling, resulting in inferior shape quality of a rolled sheet, uneven microstructure, and ultimately inferior physical properties of a final part.

[0005] Additionally, the techniques propose manufacturing methods by considering only the alloying elements and the ratio of each component, even though high-strength hot-rolled steel sheets with low yield ratios are steels that utilize the mobile dislocation formed in a boundary between soft and hard phases in the microstructure, and therefore there is no solution for cases in which actual shear formability is inferior. Here, shear formability is the first operation in the forming process of parts, and high-strength steels have a problem in that cracks occur during shear formability when the uniformity of a microstructure and components in a thickness direction is inferior. Such cracks may cause serious cracks during forming parts or adversely affect durability during use.[Prior art Document][Patent Document]

[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. 4502897 (Patent Document 5) Korean Patent Publication No. 1543838 Summary of InventionTechnical Problem

[0007] An aspect of the present disclosure is to provide steel sheet and a method for manufacturing the steel sheet.

[0008] An aspect of the present disclosure is to provide a high strength steel sheet having excellent shear formability and a method for manufacturing the steel sheet.

[0009] The aspects of the present disclosure are not limited to the above-described contents. Those skilled in the art will have no difficulty in understanding additional aspects of the present disclosure from the overall contents of this specification.Solution to Problem

[0010] According to an embodiment of the present disclosure, provided is a steel sheet including, by wt%, 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%, and a balance of Fe and inevitable impurities, wherein a value X defined in the following relational expression 1 may be 0.010 to 0.200, a value T defined in the following relational expression 2 may be 1.500 to 4.200, a microstructure may include, in area%, a hard phase including bainite and martensite of 30 to 70%, and a soft phase including ferrite of 30~70%, and pearlite of 3% or less, and an average dislocation density of a hard phase is 2.0 to 3.0x10 14< m -2< , and an average dislocation density of a soft phase is 0.50 to 2.00x10 14< m -2< . X = Nb / 93 + A / 48 + V / 51 / C / 12 + N / 14 A = Ti − 3.42 N − 1.5 S (in the relational expression, [Nb], [V], [C], [N], [Ti] and [S] are wt% of each element.) T = Mn + 2.8 Mo + 1.5 Cr + 500 B (in the relational expression, [Mn], [Mo], [Cr], and [B] are wt% of each element.)

[0011] The steel sheet may further include, by wt%, at least one 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%.

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

[0013] The steel sheet may be configured so that when punching with a punching clearance of 5 to 20%, microcracks with a length of 0.1mm or more in a shear plane are 10 pieces / cm 2< or less, and a maximum length of cracks is 1mm or less.

[0014] According to an embodiment, provided is a method for manufacturing a steel sheet including: reheating a steel slab including, by wt%, 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%, and a balance of Fe and inevitable impurities, wherein a value X defined in the following relational expression 1 is 0.010 to 0.200, and a value T defined in the following relational expression 2 is 1.500 to 4.200, hot-rolling the reheated steel slab; primarily cooling the steel sheet manufactured in the hot rolling 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 sheet for 4.0 to 10.0 seconds; and secondly cooling and coiling 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, wherein during the primary cooling, based on a steel sheet width direction, edge portions corresponding to each 30% region in a direction oriented from both ends to the other end are cooled to a temperature range of 500 to 600°C in a surface temperature (TE), and a central portion of a central 40% region corresponding to a region excluding both edge portions is cooled to a temperature range of 430 to 500°C in a surface temperature (TC). X = Nb / 93 + A / 48 + V / 51 / C / 12 + N / 14 A = Ti − 3.42 N − 1.5 S (in the relational expression, [Nb], [V], [C], [N], [Ti] and [S] are wt% of each element.) T = Mn + 2.8 Mo + 1.5 Cr + 500 B (in the relational expression, [Mn], [Mo], [Cr], and [B] are wt% of each element.)

[0015] The steel slab may further include, by wt%, at least one 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 may be performed in a temperature range of 1150 to 1350°C, and the hot rolling may be performed at a finishing rolling temperature of 850 to 1150°C.

[0017] After the air cooling, an average temperature of the steel sheet may be 550 to 650°C.Advantageous Effects of Invention

[0018] According to an embodiment of the present disclosure, a steel sheet and a method for manufacturing the steel sheet may be provided.

[0019] According to an embodiment of the present disclosure, a high-strength steel sheet having excellent shear formability and a method for manufacturing the steel sheet may be provided.

[0020] According to an embodiment of the present disclosure, a steel sheet that may be used for a cross member and a subframe, which have a relatively large forming amount among automobile chassis components, and a method for manufacturing the steel sheet may be provided.Brief Description of Drawings

[0021] FIGS. 1A and 1B illustrate the Relational Expression between the number of occurrences by crack size of a shear plane of the inventive example and the comparative example according to the punching clearance of 10% and 20%, respectively.Best Mode for Invention

[0022] Hereinafter, preferred embodiments of the present disclosure will be described. The embodiments of the present disclosure may be modified in various forms, and the scope of the present disclosure should not be construed as being limited to the embodiments described below. These embodiments are provided to explain the present disclosure in more detail to those skilled in the art to which the present disclosure pertains.

[0023] The present disclosure has researched a method to overcome the above-described conventional problems, prevent fractures from occurring in the shear and a cross-section of punching forming portion, and prevent fatigue failure from occurring during use of a product. To this end, a method has been researched to suppress the formation of microcracks and enable smooth forming of the product.

[0024] Accordingly, the inventor of the present disclosure has investigated the shear formability and a change in microcracks formed on a shear surface according to the characteristics of the components and microstructures for steels having different alloy compositions and microstructures. As a result, the inventor has confirmed that the shear formability is superior when a soft phase and a hard phase forming a microstructure satisfy a specific dislocation density range rather than a composition ratio of the microstructure, and has completed the present disclosure.

[0025] Hereinafter, the present disclosure will be described in detail.

[0026] Hereinafter, a steel composition of the present disclosure will be described in detail.

[0027] Unless otherwise specifically stated in the present disclosure, the % indicating the content of each element is based on weight.

[0028] According to an embodiment of the present disclosure, a steel sheet may include, by wt%, 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%, and a balance of Fe and inevitable impurities.Carbon (C): 0.030 to 0.150%

[0029] Carbon (C) is the most economical and effective element for strengthening steel, and the hardness value of each component and carbon (C) are the most economical and effective elements for strengthening steel, and have a great influence on a hardness value of each component and the dislocation density. With an increase in an added amount thereof, the hardenability increases, and the fraction of hard phases such as bainite and martensite in the microstructure increases, thereby increasing the dislocation density and tensile strength. Additionally, fine precipitates are formed together with Ti and Nb, which have a high affinity for carbon (C), so that the grain size becomes fine and the precipitation strengthening effect also increases, thereby increasing both the yield strength and the tensile strength. When the content of carbon (C) is less than 0.030%, it may be difficult to obtain a sufficient strengthening effect. According to an embodiment of the present disclosure, in order to stably secure a higher level of strength, carbon (C) may be included in the content of 0.050% or more. On the other hand, if the content thereof exceeds 0.150%, the fraction of each phase including bainite and martensite increases, and a hardness value of the phase also increases, which may cause excessive strength increase, and there may be a problem that an elongation rate and formability decrease, and the weldability may also be inferior. In the present disclosure, in order to secure formability more stably, carbon (C) may be included in the content of 0.120% or less.Silicon (Si): 0.01 to 1.00%

[0030] Silicon (Si) deoxidizes molten steel, has a solid solution strengthening effect, and delays the formation of coarse carbides, which is advantageous in improving formability. In the present disclosure, silicon (Si) may be included in the content of 0.01% or more in order to obtain the above-described effect. According to an embodiment of the present disclosure, silicon (Si) may be included in the content of 0.10% or more. However, when the content thereof exceeds 1.00%, during hot rolling, a red scale due to silicon (Si) may be formed on a surface of the steel sheet, which not only greatly deteriorates the surface quality of the steel sheet, but also reduces ductility and weldability. In an embodiment of the present disclosure, silicon (Si) may be included in the content of 0.90% or less.Manganese (Mn): 1.00 to 2.50%

[0031] Manganese (Mn), like Si, is an element that is effective in solid solution strengthening steel and increases the hardenability of steel, facilitating the formation of hard phases, bainite and martensite, during cooling after hot rolling. However, when the content thereof is less than 1.00%, the above-described effect due to addition may not be obtained. In an embodiment of the present disclosure, manganese (Mn) may be included in the content of 1.40% or more. On the other hand, when the content thereof exceeds 2.50%, the hardenability may increase significantly, so that the fraction of each phase, including bainite and martensite, and the hardness value of the phase increase, which may cause problems such as excessive strength increase and reduced formability. Additionally, when casting a slab in a casting process, a segregation portion may be greatly developed in a center of a thickness thereof, and during cooling after hot rolling, a microstructure in a thickness direction may be formed unevenly, which may cause poor elongation flangeability. Specifically, it may be difficult to uniformly manufacture the microstructure in the full length and full width of a hot-rolled sheet when cooling. In an embodiment of the present disclosure, an upper limit of the content of manganese (Mn) may be limited to 2.30%.Aluminum (Al): 0.01 to 0.80%

[0032] Aluminum (Al) is a component mainly added for deoxidation and has a ferrite transformation promotion effect. When the content thereof is less than 0.01%, the above-described addition effect may be insufficient. According to an embodiment of the present disclosure, aluminum (Al) may be included in the content of 0.02% or more. On the other hand, when the content thereof exceeds 0.80%, aluminum (Al) may be combined with N to form AlN, which may easily cause corner cracks in a slab during casting and may easily cause defects due to inclusion formation. According to an embodiment of the present disclosure, an upper limit of the content of aluminum (Al) may be limited to 0.50%.Chromium (Cr): 0.005 to 0.500%

[0033] Chromium (Cr) strengthens steel with solid solution and, during cooling, delays ferrite phase transformation, thereby assisting in form bainite. However, when the content of chromium (Cr) is less than 0.005%, the above-described effect due to addition may not be obtained. In order to more effectively secure the above-described effect, the present disclosure may include chromium (Cr) in the content of 0.100% or more. On the other hand, when the content thereof exceeds 0.500%, ferrite transformation may be excessively delayed, resulting in inferior elongation due to the formation of martensite. Additionally, similar to Mn, a segregation portion in a center of a thickness thereof is greatly developed, and the microstructure in a thickness direction becomes non-uniform, resulting in inferior elongation flangeability. According to an embodiment of the present disclosure, an upper limit thereof may be limited to 0.300%.Molybdenum (Mo): 0.005 to0.300%

[0034] Molybdenum (Mo) increases the hardenability of steel and facilitates the formation of bainite structure. However, when the content thereof is less than 0.005%, the above-described effect due to addition may not be obtained. In an embodiment of the present disclosure, the content thereof may be 0.050% or more. On the other hand, when the content of molybdenum (Mo) exceeds 0.300%, martensite is formed due to excessive increase in quenchability, resulting in rapid inferior formability. Additionally, this may be economically disadvantageous and may be detrimental to weldability. In an embodiment of the present disclosure, an upper limit thereof may be limited to 0.200%.Phosphorus (P): 0.001 to 0.050%

[0035] Phosphorus (P), like Si, simultaneously has the effects of solid solution strengthening and promoting ferrite transformation. However, since a large ament of manufacturing costs are required to manufacture below 0.001%, this may be economically disadvantageous and insufficient to obtain strength, so that a lower limit thereof may be limited to 0.001%. On the other hand, when phosphorus (P) exceeds 0.050%, brittleness occurs due to grain boundary segregation, and fine cracks are likely to occur during molding, and ductility, elongation flangeability, and impact resistance may be significantly deteriorated.Sulfur (S): 0.001 to 0.010%

[0036] Sulfur (S) is an impurity present in steel, and when the content thereof exceeds 0.010%, sulfur (S) may be combined with Mn, or the like, to form non-metallic inclusions, which may easily cause fine cracks to occur during cutting and processing of steel, and may significantly reduce the elongation flangeability and impact resistance. In the present disclosure, sulfur (S) may be included in the content of 0.005% or less. In the present disclosure, there is no particular limitation on a lower limit of the content of sulfur (S), but in order to manufacture the content of sulfur (S) below 0.001%, a significant amount of time is required during steelmaking, which may reduce productivity, so that a lower limit of the content of sulfur (S) may be limited to 0.001% in consideration thereof.Nitrogen (N): 0.001 to 0.010%

[0037] Nitrogen (N) is a representative solid-solution strengthening element along with C, and forms coarse precipitates along with Ti, Al, and the like. In general, the strengthening effect of nitrogen (N) is better than that of carbon, but there is a problem that the toughness decreases significantly as the amount of nitrogen (N) in steel increases. Additionally, in order to manufacture the content of nitrogen (N) below 0.001%, it takes much time during steelmaking, which reduces productivity, so that the lower limit may be limited to 0.001%.

[0038] The steel of the present disclosure may include the remaining iron (Fe) and inevitable impurities in addition to the composition described above. Since the inevitable impurities may be unintentionally mixed in during a normal manufacturing process, they may not be excluded. Since these impurities are known to anyone skilled in the art of normal steel manufacturing, not all of the contents are specifically mentioned in this specification.

[0039] The steel sheet according to an embodiment of the present disclosure may further include at least one 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%.Niobium (Nb): 0.005 to 0.030%

[0040] Niobium (Nb) is a representative precipitation strengthening element along with Ti and V, and is effective in improving the strength and impact toughness of steel due to the effect of grain refinement by precipitation and recrystallization delay during hot rolling. When the content of niobium (Nb) is less than 0.005%, the above-described effect may not be obtained. On the other hand, if the content thereof exceeds 0.030%, there is a problem of poor elongation flangeability due to the formation of elongated grains and coarse composite precipitates due to excessive recrystallization delay during hot rolling.Titanium (Ti): 0.005 to 0.120%

[0041] Ti is a representative precipitation strengthening element along with Nb and V, and forms coarse TiN in steel due to strong affinity with N. TiN has the effect of suppressing grain growth during the heating process for hot rolling. Additionally, the titanium (Ti) remaining after reacting with nitrogen is dissolved in the steel and is combined with carbon, thereby forming TiC precipitates, which is a useful component for improving the strength of the steel. When the content of titanium (Ti) is less than 0.005%, the above-described effect may not be obtained. On the other hand, when the content thereof exceeds 0.120%, there is a problem of inferior elongation flangeability during forming due to the occurrence of coarse TiN and coarsening of precipitates. In an embodiment of the present disclosure, an upper limit thereof may be 0.105%. In an embodiment of the present disclosure, an upper limit thereof may be 0.100%.Vanadium (V): 0.005 to 0.200%

[0042] Vanadium (V) is a representative precipitation strengthening element along with Nb and Ti, and rarely precipitates during hot rolling and forms precipitates after coiling to improve the strength of the steel. Accordingly, vanadium (V) is effective in improving additional strength without increasing deformation resistance and rolling load due to recrystallization delay during hot rolling. In order to obtain this effect in the present disclosure, the content of vanadium (V) may be included at 0.005% or more. However, when the content is excessive, there is a problem of inferior elongation flangeability due to the formation of coarse precipitates, and this is also economically disadvantageous. Accordingly, in the present disclosure, an upper limit thereof may be limited to 0.200%, and in an embodiment, an upper limit thereof may be limited to 0.150%.Boron (B): 0.0003 to 0.0030%

[0043] When boron (B) exists in a solid solution state in steel, boron (B) is mainly segregated in grain boundaries and has the effect of improving the brittleness of steel by stabilizing grain boundaries, and plays a role in stabilizing solid solution N and suppressing the formation of coarse AlN nitride. Additionally, this is effective in the formation of bainite and martensite, which are hard phases, by delaying ferrite phase transformation. In order to secure the above-described effect, the present disclosure may include boron (B) in an amount of 0.0003% or more. On the other hand, when the content thereof exceeds 0.0030%, the effect due to addition no longer increases, and there is a problem that ductility decreases and formability becomes inferior. According to an embodiment of the present disclosure, boron (B) may be included in the content of 0.0020% or less.

[0044] The steel sheet according to an embodiment of the present disclosure may have a value X defined in the following relational expression 1 of 0.010 to 0.200, and a value T defined in the following relational expression 2 of 1.500 to 4.200. X = Nb / 93 + A / 48 + V / 51 / C / 12 + N / 14 A = Ti − 3.42 N − 1.5 S (where [Nb], [V], [C], [N], [Ti], and [S] are wt% of each element.) T = Mn + 2.8 Mo + 1.5 Cr + 500 B (where [Mn], [Mo], [Cr], and [B] are wt% of each element.)

[0045] When the value X defined in Relational Expression 1 exceeds 0.200, the formation of precipitates may increase and the strength may increase somewhat, but due to the delay in recrystallization during hot rolling, a microstructure elongated in the rolling direction may be likely to be formed, which may decrease an elongation rate in a vertical direction of rolling. Additionally, when the hot-rolled steel sheet is cooled, since solid C and solid N atoms in the non-transformed phase are insufficient, it may be difficult for the hard phase to be stably formed, and the grain boundary becomes weak, which may cause the quality of a shear plane to be inferior. In an embodiment of the present disclosure, the value X may be 0.180 or less. On the other hand, when the value X is less than 0.010, grain growth during reheating may be easily performed, recrystallization during hot rolling may be uneven, and coarse grains may be formed locally, and since solid C and solid N are excessively excessive, a hardness value of the hard phase tends to increase, which may ultimately cause the elongation rate to be inferior. In an embodiment of the present disclosure, the value may be 0.030 or more. Meanwhile, when the corresponding alloy element of Relational Expression 1 is not added, 0 may be substituted.

[0046] Relational expression 2 is a factorization of a combination of alloy elements that may maintain the formation of bainite, martensite, and MA phases, which are hard phases in the microstructure, at an appropriate level. As the value X defined in relational expression 2 increases, the formation of hard phases such as bainite, martensite, and MA may increase, and a hardness value of each hard phase may also increase. Accordingly, in the present disclosure, the value X may be limited to 1.500 or more for the desired strength. According to an embodiment of the present disclosure, the value X may be limited to 2.000 or more. On the other hand, as the value X increases, this is more advantageous for securing strength, but when the value is excessive, the ductility of the steel may decrease and a hardness difference between the soft phase and the hard phase increases more than necessary, which may result in poor shear formability. Additionally, there is a problem that the material deviation increases in an overall length and width of the hot-rolled steel sheet. Accordingly, in the present disclosure, an upper limit of the value may be limited to 4.200. According to an embodiment of the present disclosure, an upper limit of the value X may be limited to 4.000.

[0047] Hereinafter, the steel microstructure of the present disclosure will be described in detail.

[0048] Unless otherwise specifically stated in the present disclosure, % indicating the fraction of microstructure is based on area.

[0049] The inventor of the present disclosure found that it may be difficult to clearly distinguish how much the inherent shear formability of steel is excellent based only on the area ratio of the microstructure composition and whether stable results are shown even when punching clearance changes. Specifically, it was confirmed that the dislocation density and physical characteristics change greatly depending on the components constituting the steel.

[0050] Accordingly, the inventor of the present disclosure has confirmed, as a result of research, that the dislocation density (Geometrical Necessary Dislocation) of the microstructure is an important factor affecting the occurrence of microcracks, which is a quality in the shear plane of the steel, and has proposed the present disclosure.

[0051] The microstructure of the steel sheet according to an embodiment of the present disclosure may include, in terms of area %, a hard phase including bainite and martensite of 30 to 70%, a soft phase including ferrite of 30 to 70%, and pearlite of 3% or less.

[0052] In the present disclosure, the microstructure is controlled to secure a yield ratio and elongation formability. Accordingly, in the present disclosure, bainite and martensite may be distinguished as hard phases, and ferrite may be distinguished as soft phases, and area fractions thereof may be limited.

[0053] In the present disclosure, bainite may include upper bainite and lower bainite, and may be distinguished from a ferritic low-temperature transformation phase in that fine carbides are formed in a lath-shaped structure.

[0054] In the present disclosure, ferrite may include equiaxed ferrite and a ferritic low-temperature transformation phase. The ferritic low-temperature transformation phase may include acicular ferrite, bainitic ferrite, and granular bainitic ferrite, and may denote ferrite having an uneven grain boundary, a high intragranular dislocation density, and a high intragranular low-angle grain boundary structure density, as compared to equiaxed ferrite.

[0055] In the present disclosure, the microstructure may be observed in a cross-section, perpendicular to a rolling direction of the steel sheet, and may be analyzed at a point 1 / 4 to 1 / 2t (t is a thickness of the steel sheet) in the thickness direction. The distinction of microstructure and measurement of area fraction may be analyzed at 3000 to 5000 times magnification using Electron Back Scattered Diffraction (EBSD, (JEOL JSM-1001F)).

[0056] When an area fraction of the hard phase exceeds 70%, the elongation rate may be greatly reduced, and in the shear plane, a ratio of fractured portions may increase, and cracks of 1 mm or more in length may significantly increase. Additionally, the dependence of the shear plane quality on the change in punching clearance may increase, and thus, when actually forming a part, an occurrence of defects may increase. On the other hand, when the area fraction is less than 30%, it may be difficult to secure target strength. Meanwhile, according to an embodiment of the present disclosure, since martensite is a relatively hard phase as compared to bainite, an increase in martensite may lead to a decrease in ductility, and in consideration thereof, an upper limit of the area fraction of martensite may be limited to 60%. Meanwhile, according to an embodiment of the present disclosure, martensite among the hard phases may be 0%.

[0057] In the present disclosure, although a small amount of MA phase may be observed, this is considered as martensite because this does not have a particularly different effect on the properties and cross-sectional quality of a punched portion suggested in the present disclosure. Additionally, since tempered martensite including fine carbides has a lower dislocation density than martensite, a hardness value of a phase decreases, and the tempered martensite may assist in improving the ductility of the entire steel, but when the size of the formed carbides increases, this may have a negative effect such as brittleness, so that and the tempered martensite may be distinguished from general martensite. However, in the present disclosure, since the dislocation density of the phase is measured and described together, it is not necessary to distinguish the martensite from the tempered martensite, so that the tempered martensite is considered as martensite.

[0058] The soft phase may help with the ductility of the steel and the formation of fine precipitates, and thus, a lower limit of an area fraction thereof may be limited to 30%. According to an embodiment of the present disclosure, an upper limit thereof may be limited to 70%.

[0059] Meanwhile, pearlite may be further included as a structure other than the hard phase and the soft phase. However, when the area fraction of pearlite exceeds 3%, this corresponds to a weak tissue during formation of a shear of the steel, and therefore, cracks of 1 mm or more in length may increase, so that an upper limit thereof may be limited to 3%.

[0060] According to an embodiment of the present disclosure, the steel sheet may have an average dislocation density (Geometrical Necessary Dislocation) of 2.0 to 3.0x10 14< m -2< in the hard phase, and an average dislocation density of 0.50 to 2.00x10 14< m -2< in the soft phase.

[0061] The average dislocation density (Geometrical Necessary Dislocation) may be calculated using kernel average misorientation (KAM) data after measuring a cross-section parallel to an rolling direction at a point 1 / 4 in a steel sheet thickness direction with EBSD, and may be calculated by the following equation. For convenience, such calculations may be performed using software such as OIM analysis ™< (EDAX) that analyzes the EBSD measurement results. DISLOCATION DENSITY GND m − 2 = 2 θ ub (where θ is average misorientation (KAM values), u is a unit length (step size in the EBSD measurement), and b is a burgers vector.)

[0062] When the average dislocation density of the hard phase is less than 2.0x10 14< m -2< , the strength may be significantly reduced, and when a value thereof exceeds 3.0x10 14< m -2< , the ductility may be reduced and a shear plane quality may be inferior.

[0063] When the average dislocation density of a soft phase is less than 0.50x10 14< m -2< , the strength may not reach a level required by the present disclosure, and burr occurrence may be severe during shear forming. On the other hand, when a value thereof exceeds 2.00x10 14< m -2< , there may be a problem of increased yield strength and decreased elongation.

[0064] Hereinafter, a method for manufacturing a steel sheet of the present disclosure will be described in detail.

[0065] According to an embodiment of the present disclosure, a steel sheet may be manufactured by reheating, hot rolling, primary cooling, air cooling, secondary cooling, and coiling a steel slab satisfying the alloy composition of the present disclosure.Reheating

[0066] A steel slab satisfying the alloy composition of the present disclosure may be reheated at a temperature range of 1150 to 1350°C.

[0067] When the reheating temperature is less than 1150°C, precipitates may not be sufficiently re-dissolved, so that the formation of precipitates may decrease in the process after hot rolling, and coarse TiN may remain, and since the steel slab is not sufficiently heated, it may be difficult to control a temperature of the steel sheet at a constant level during hot rolling. On the other hand, when the temperature exceeds 1350°C, the strength may be reduced due to abnormal grain growth of austenite crystal grains.Hot rolling

[0068] The reheated steel slab may be hot-rolled at a finishing rolling temperature of 850 to 1150°C.

[0069] During hot rolling, when the finishing rolling temperature exceeds 1150°C, the temperature of a hot-rolled steel sheet increases, which may cause the grain size to become coarse and the surface quality of the hot-rolled steel sheet to become poor. On the other hand, when the temperature is less than 850°C, the elongated grains may develop due to excessive recrystallization delay, which may cause severe anisotropy and poor formability.Primary cooling

[0070] A steel sheet manufactured in the hot rolling operation may be primarily cooled at an average cooling rate of 50 to 100°C / s to a temperature range of 430 to 600°C. During the primary cooling, an edge portion corresponding to 30% of a region from both ends toward the other end in a width direction of the steel sheet may be cooled to a temperature range of 500 to 600°C in the surface temperature (TE), and a central portion of a central 40% region corresponding to a region excluding both edges may be cooled to a temperature range of 430 to 500°C in the surface temperature (TC).

[0071] In the present disclosure, a region corresponding to 30% of each region from both ends toward the other end or center in a width direction of the steel sheet, that is, a region corresponding to a total of 60% of the entire steel sheet, is identified as an edge portion, and the central 40% region excluding the edge portion is identified as a center portion.

[0072] In the present disclosure, a soft phase among the microstructures of the steel may be formed during the primary cooling and air cooling, and a hard phase may be formed during the secondary cooling and coiling. However, since an untransformed phase just before the secondary cooling is formed into a hard phase after the secondary cooling, the soft phase should be formed uniformly at each width position of the steel sheet immediately after the primary cooling and air cooling. Normally, when the steel sheet is cooled, heat transfer proceeds quickly in the edge portion of the steel sheet, so that more hard phases are formed as compared to the center. Accordingly, in order to provide a uniform cooling speed at each width position of the steel sheet, the present disclosure aims to control cooling termination temperatures of the center and the edge portion differently. As suggested in the present disclosure, when the soft phase is formed uniformly at each width position during the primary cooling and air cooling, a ratio of the hard phase formed during the secondary cooling becomes constant, and the dependence on the cooling conditions after coiling may be reduced. When the soft phase and the hard phase are formed uniformly, the excellent material uniformity and shear formability aimed at in the present disclosure may be secured.

[0073] During the primary cooling, when the cooling rate is less than 50°C / s, the ferrite fraction may be formed excessively, and the average dislocation density may be below a target level, which may be detrimental to securing strength. On the other hand, when the cooling rate exceeds 100°C / s, during the primary cooling, the surface temperature (TC) of the central portion may be excessively low, which may significantly reduce the ferrite fraction and increase the hard phase more than necessary, and thus may result in insufficient elongation.

[0074] On the other hand, in the present disclosure, there may be a concern that a difference in a cooling rate may become severe depending on a width position of the steel sheet, and thus, in order to suppress the formation of the hard phase differently depending on a width position, the edge portion and the central portion may be cooled separately to different temperature ranges.

[0075] Specifically, in the present disclosure, during the subsequent air cooling process, a temperature of the steel sheet is reheated so that the steel sheet has a uniform temperature of 550 to 650°C. Accordingly, the temperature of the central portion needs to be supercooled below the target temperature of 550 to 650°C, and preferably, the edge portion, which has a high cooling rate, may be cooled to a higher temperature range than the central portion. Accordingly, for this purpose, in the present disclosure, a surface temperature (TE) of the edge portion may be cooled to 500 to 600°C, and a surface temperature (TC) of the central portion may be cooled to 430 to 500°C.

[0076] When the surface temperature of the edge portion is less than 500°C, the formation of the soft phase may be insufficient, and there may be a problem of falling short of a target temperature range of the hot-rolled steel sheet during air cooling. On the other hand, when the temperature exceeds 600°C, the surface temperature may exceed the target temperature range during air cooling, and there may be a problem of excessive formation of the soft phase among the final microstructures. According to an embodiment of the present disclosure, the surface temperature of the edge portion may be 510°C or higher during the primary cooling.

[0077] When the surface temperature of the central portion is below 430°C, there is a problem that the phase transformation of the hard phase bainite occurs. On the other hand, when the temperature exceeds 500°C, the effect of the heat recovery may be reduced, and there may be a problem of falling short of the target temperature range.Air cooling

[0078] A steel sheet primarily cooled may be air-cooled for 4.0 to 10.0 seconds.

[0079] The intentional cooling of the primarily-cooled steel plate may be terminated so that the temperature of the steel plate may be restored to a target temperature by internal latent heat and transformation heat generation. When the primarily-cooled steel plate is air-cooled, the steel plate may be restored to an average temperature of the steel plate in the temperature range of 550 to 650°C.

[0080] During air-cooling, when the time is less than 4.0 seconds, the restoration effect may not be effective. On the other hand, when the time exceeds 10.0 seconds, the ferrite fraction in the microstructure may significantly increase, and hard phases of bainite and martensite may decrease. Additionally, a pearlite structure and coarse carbides may be formed in a central portion of a thickness of the high-temperature region of the steel plate, which may cause the cross-sectional quality to be inferior after shear forming.Secondary cooling and coiling

[0081] The air-cooled steel plate may be secondarily-cooled and coiled at an average cooling rate of 10 to 100°C / s to a temperature range of 50 to 200°C.

[0082] When the coiling temperature exceeds 200°C, the average dislocation density of the hard phase may exceed a range suggested by the present disclosure, making it difficult to secure strength. According to an embodiment of the present disclosure, the coiling temperature may be limited to 150°C or less. On the other hand, when the temperature is less than 50°C, martensite may be formed in an excessive amount, and the average dislocation density of the hard phase may exceed the range suggested by the present disclosure, resulting in poor elongation of the steel. Additionally, there may be a problem in which the cooling water remains and the steel plate corrodes. According to an embodiment of the present disclosure, the coiling temperature may be limited to 70°C or more.

[0083] When the cooling rate exceeds 100°C / s, the average dislocation density of the hard phase may be excessively high, resulting in a problem in which an elongation rate thereof decreases. There is no particular limitation on the lower limit of the cooling rate, but in order to control the cooling rate to be less than 10°C / s, there may be a problem in that a length of the cooling zone equipment should be elongated, and there may be a problem in that it may be difficult to manufacture at the target coiling temperature of 200°C or less.

[0084] The steel plate manufactured in this manner has a tensile strength of 780 MPa or more, and a yield ratio of 0.70 to 0.85, and when punching formed with a punching clearance of 5 to 20%, the number of microcracks with a length of 0.1 mm or more in the shear plane is 10 pieces / cm 2< or less, and a maximum crack length is 1 mm or less, which may secure excellent strength and shear formability and a low yield ratio.Mode for Invention

[0085] Hereinafter, the present disclosure will be described more specifically through examples. However, it should be noted that the following examples are only intended to illustrate the present disclosure and explain the same in more detail, and are not intended to limit the scope of the rights of the present disclosure.[Inventive Example]

[0086] A steel slab having a composition in Table 1 below was manufactured into a steel plate under the conditions described in Table 2 below. A reheating temperature not disclosed in Table 2 below was 1250°C, and a thickness of the steel plate immediately after hot rolling was manufactured to be 3.2 mm. [Table 1]Steel TypeAlloy composition (wt%)Relati onal Expres sion 1 (X)Relati onal Expres sion 2 (T)CSiMnCrMoNbTiVAlPSNBA0.07 50.701.700.20 00.02 00.03 00.10 00.00 50.050.01 90.00 10.00 50.00 040.3202.256B0.07 50.702.300.70 00.20 00.02 50.05 00.00 50.030.00 90.00 20.00 40.00 100.1624.410C0.07 50.701.200.00 50.00 10.02 00.06 00.00 50.040.01 00.00 20.00 40.00 040.1861.410D0.07 50.701.700.20 00.02 00.00 50.02 50.00 50.050.01 90.00 10.00 50.00 040.0432.256E0.07 50.702.200.30 00.15 00.02 00.06 00.00 10.030.01 00.00 30.00 40.00 250.1694.320F0.07 50.701.100.10 00.01 00.02 00.06 00.00 10.030.01 00.00 30.00 40.00 030.1691.428G0.09 50.251.400.40 00.01 00.00 50.03 00.00 50.100.01 00.00 30.00 40.00 040.0492.228H0.08 00.801.900.50 00.02 00.03 00.00 50.00 50.20.01 00.00 30.00 40.00 040.0212.906I0.06 00.201. 800.10 00.20 00.02 00.05 00.00 50.030.01 00.00 30.00 40.00 040.1852.710J0.07 50.701.700.20 00.02 00.00 50.02 50.00 50.050.01 90.00 10.00 50.00 040.0432.256K0.11 00.081.450.40 00.01 00.00 50.02 00.00 50.200.01 00.00 20.00 40.00 040.0232.278L0.08 00.902.300.20 00.20 00.01 50.02 00.00 50.040.01 00.00 20.00 40.00 200.0474.160M0.15 00.902.200.20 00.20 00.01 50.02 00.00 50.040.01 00.00 20.00 40.00 040.0263.260 X=Nb / 93+A / 48+V / 51 / C / 12+N / 14A=Ti−3.42N−1.5S (where [Nb], [V], [C], [N], [Ti], and [S] are wt% of each element.)T=Mn+2.8Mo+1.5Cr+500B (were [Mn], [Mo], [Cr], and [B] are wt% of each element.) [Table 2] Spec imen Numb erSteel TypeHot RollingPrimary CoolingAir-CoolingSecondary coolingFinishing Rolling Temperatu re (°C)Rate (°C / s)Temperature (°C)Time (secon d)Temperat ure after air-cooling (°C)Rate (°C / s)Coiling Temperatu re (°C)Edge portio nCentra l Portio n1A874635384865.560155952B865745224475.7590491103C860725304886.260553784D873256425746.264045855D866984704026.055856896D8906557049012.064355927D8746658348512.0668481158D855625664762.552550889D885685454885.85941122610D881655534966.2607852711D866715584826.56034822012E875625334624.8610488513F890555754725.26055510514G873655204896.96045510515H880715384956.5625628816I879775544706.35895311017J887685494925.8595587518K892625774606.2611609219L885795714624.8593539020M870655624585.86084982

[0087] Tables 3 and 4 below illustrate a microstructure and mechanical properties of the manufactured steel sheets. The microstructure was observed in a cross-section, perpendicular to a rolling direction of the steel sheet, and was analyzed at 1 / 4 to 1 / 2 points based on a thickness direction. The classification and area fraction of ferrite, ferrite-type low-temperature transformation phase, bainite, martensite, and pearlite formed in the steel were measured using electron back-scattered diffraction (EBSD, (JEOL JSM-1001F)) at a magnification of 3000 to 5000. Additionally, the average dislocation density (Geometrical Necessary Dislocation, GND) was measured using OIM analysis ™< (EDAX) after EBSD measurement based on a cross-section, parallel to the rolling direction, at 1 / 4 points of the steel sheet in a thickness direction.

[0088] The mechanical properties were measured and presented in terms of yield strength, tensile strength, fracture elongation, and yield ratio. Here, the 0.2% off-set yield strength (YS), tensile strength (TS), and fracture elongation (T-El) are the results of testing the JIS No. 5 standard specimen taken in a direction, perpendicular to the rolling direction. Additionally, the number of cracks in the punched cross-section was measured and displayed as a physical property. This was evaluated by punching a hole with a diameter of 10 mm, and in this case, the number of microcracks observed in the cross-section, parallel to the rolling direction, and the cross-section, perpendicular to the rolling direction, were averaged after punching with different punching clearances of 5, 10, and 20%, and each result value thereof is the number of occurrences in each crack length. [Table 3]Specimen NumberSteel TypeMicrostructurePropertiesDivisionFraction (area %)Dislocation Density (10 14< ,m -2< )Yield Stren gth (MPa)Tensi le Stren gth (MPa)Elonga tion at Break (%)Yield RatioSoft PhaseHard PhasePSoft PhaseHard PhaseFBFBM1A244127802.12.9728837180.87Comparative Example 12B027621101.93.31068128950.83Comparative Example 23C33590800.91.8517765210.68Comparative Example 34D29529550.81.9584751210.78Comparative Example 45D027621101.72.4725892100.81Comparative Example 56D234223841.52.3667768180.87Comparative Example 67D214921541.22.2660762190.87Comparative Example 78D523324001.32.879695090.84Comparative Example 89D1938133001.73.1680843140.81Comparative Example 910D2135162801.63.1655880130.74Comparative Example 1011D2333242001.51.9723812170.89Comparative Example 1112E026591501.93.41013122550.83Comparative Example 1213F28609300.81.8504750220.67Comparative Example 1314G1938192401.22.1609825190.74Inventive Example 115H732283301.92.78271056120.78Inventive Example 216I1536282101.62.3715909140.79Inventive Example 317J2534251601.552.4633854180.74Inventive Example 418K2133341111. 622.2628831170.76Inventive Example 519L332333201.82.9950120970.79Inventive Example 620M530273801.752.8975121360.80Inventive Example 7

[0089] * F: equiaxed ferrite, BF: ferritic low-temperature transformation phase, B: bainite, M: martensite, P: pearlite [Table 4]Specimen NumberSteel TypePropertiesDivisionNumber of cracks by length in a punched cross-section (piece / cm 2< )Clearance 5%Clearance 10%Clearance 20%<0 .1 (m m)0.1~1.0 (mm)>1.0 (mm)<0.1 (mm)0.1~1.0 (mm)>1.0 (mm)<0.1 (mm)0.1~1.0 (mm)>1.0 (mm)1A1501802150Comparative Example 12B38141128132Comparative Example 23C100020310Comparative Example 34D232131182Comparative Example 45D2404805122Comparative Example 56D240021272Comparative Example 67D240362562Comparative Example 78D570462692Comparative Example 89D1701802111Comparative Example 910D130121552Comparative Example 1011D220530731Comparative Example 1112E381310110153Comparative Example 1213F100110220Comparative Example 1314G200100010Inventive Example 115H130130350Inventive Example 216I100220010Inventive Example 317J100110210Inventive Example 418K000000210Inventive Example 519L220220550Inventive Example 620M330230460Inventive Example 7

[0090] As illustrated in Tables 3 and 4, in the case of the inventive examples satisfying the alloy composition and manufacturing conditions of the present disclosure, the microstructure characteristics proposed by the present disclosure were satisfied, and the physical properties targeted by the present disclosure were also secured.

[0091] FIGS. 1A and 2B illustrate a relationship between the number of occurrences of crack sizes in the shear plane of the inventive examples and the comparative examples according to the punching clearance of 10% and 20%, respectively. Specifically, for the punching clearances of 10% and 20%, it may be confirmed that the number of cracks occurring in the comparative examples is greater than that in the inventive examples, and it may also be confirmed that no cracks exceeding 1.0 mm occurred in the inventive examples, and the number of occurrences of cracks of 1.0 mm or less was significantly lower than that in the comparative examples.

[0092] On the other hand, Comparative Example 1 is an example that satisfies a content range of the alloy element proposed by the present disclosure, but does not satisfy the condition of Relational Expression 1. As a result, the average dislocation density of the soft phase exceeded the range proposed by the present disclosure, and this is determined to have occurred due to the increased formation of fine precipitates in the soft phase. No coarse cracks exceeding 1 mm in length were found in a punched portion, but when the clearance was 20%, the occurrence of cracks of 0.1 to 1.0 mm in length increased significantly. Additionally, there was also a problem that the formability was inferior because the yield strength increased excessively due to work hardening during forming due to the high yield ratio.

[0093] Comparative Examples 2, 3, 12 and 13 are examples that did not satisfy Relational Expression 2, and Comparative Examples 2 and 12 included an excessive amount of alloy components having high hardenability effects, and thus stably secured the strength, but an elongation rate thereof was insufficient. As a result, the quality of the shear plane was also inferior. Comparative Examples 3 and 13 lacked alloy components having excellent hardenability effects, so that the hard phase was not formed to the level proposed by the present disclosure, and as a result, the target strength was not secured. Additionally, as the clearance increased, the occurrence of cracks became more severe, and cracks exceeding 1 mm in length were also confirmed.

[0094] Comparative Examples 4 and 5 are cases in which a cooling termination temperature was outside the range suggested by the present disclosure during the primary cooling after hot rolling. In Comparative Example 4, during the primary cooling, the termination temperature exceeded an upper limit thereof, and thus, the formation of the hard phase was insufficient, and unnecessary pearlite was also formed, which caused the cross-sectional quality after punching to be inferior. In Comparative Example 5, during the primary cooling, the termination temperature exceeded a lower limit thereof, which caused the fraction of the soft phase to be insufficient, and the hard phase was excessively formed, which caused the cross-sectional quality after punching to be inferior.

[0095] Comparative Examples 6 and 7 are cases in which the air-cooling time after the first cooling was outside the range of the present disclosure. The exposure time in the high temperature zone was long, so that the soft phase fraction increased significantly due to latent heat and heat generation due to phase transformation inside the steel plate, and pearlite was also formed, resulting in poor cross-sectional quality after punching, and the yield ratio also exceeded the range proposed by the present disclosure. Specifically, in the case of Comparative Example 7, it may be confirmed that the average temperature of the steel plate after air cooling exceeded the temperature range proposed by the present disclosure.

[0096] Comparative Example 8 is an example in which the air-cooling time after the first cooling fell short of the range of the present disclosure. The secondary cooling was performed before the steel plate was reheated, so that the soft phase fraction fell short of the target level, and the hard phase fraction exceeded the target range. As a result, the cross-sectional quality after punching was poor.

[0097] Comparative Example 9 is a case in which, during the secondary cooling, the cooling speed was excessively fast, and overcooling was performed, which failed to satisfy the coiling temperature range targeted by the present disclosure. As a result, the dislocation density of the hard phase exceeded the target range, and the cross-sectional quality after punching was poor. The main cause is determined to have been an increased difference in properties between the soft phase and the hard phase.

[0098] Comparative Examples 10 and 11 are cases in which the cooling termination temperature during the secondary cooling was outside the range suggested by the present disclosure. Comparative Example 10 is a case in which the coiling temperature after the secondary cooling fell short of the suggested temperature range, and the dislocation density of the hard phase was excessively high, and as a result, the cross-sectional quality of after punching was inferior. Comparative Example 11 is a case in which a coiling temperature after the secondary cooling exceeded the suggested temperature range, and the dislocation density of the hard phase is below the suggested level. As a result, the yield ratio was excessively high, and the cross-sectional quality was also inferior.

[0099] Although the present disclosure has been described in detail through examples, other forms of examples are also possible. Therefore, the technical idea and scope of the claims described below are not limited to the embodiments.

Claims

1. A steel sheet, comprising: by wt%, 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%, and a balance of Fe and inevitable impurities, wherein a value X defined in the following relational expression 1 is 0.010 to 0.200, a value T defined in the following relational expression 2 is 1.500 to 4.200, a microstructure includes, in area%, a hard phase including bainite and martensite of 30 to 70%, and a soft phase including ferrite of 30~70%, and pearlite of 3% or less, and an average dislocation density of a hard phase is 2.0 to 3.0x1014m-2, and an average dislocation density of a soft phase is 0.50 to 2.00x1014m-2, X = Nb / 93 + A / 48 + V / 51 / C / 12 + N / 14 A = Ti − 3.42 N − 1.5 S where [Nb], [V], [C], [N], [Ti] and [S] are wt% of each element, T = Mn + 2.8 Mo + 1.5 Cr + 500 B where [Mn], [Mo], [Cr], and [B] are wt% of each element.

2. The steel sheet of claim 1, wherein the steel sheet further includes, by wt%, at least one 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 sheet of claim 1, wherein the steel sheet is a steel sheet having a tensile strength of 780 MPa or more and a yield ratio of 0.70 to 0.85.

4. The steel sheet of claim 1, wherein the steel sheet is configured so that, when punching with a punching clearance of 5 to 20%, microcracks with a length of 0.1mm or more in a shear plane are 10 pieces / cm2 or less, and a maximum length of cracks is 1mm or less.

5. A method for manufacturing a steel sheet, comprising: reheating a steel slab including, by wt%, 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%, and a balance of Fe and inevitable impurities, wherein a value X defined in the following relational expression 1 is 0.010 to 0.200, and a value T defined in the following relational expression 2 is 1.500 to 4.200, hot-rolling the reheated steel slab; primarily cooling the steel sheet manufactured in the hot rolling 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 sheet for 4.0 to 10.0 seconds; and secondly cooling and coiling 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, wherein during the primary cooling, based on a width direction of the steel sheet, edge portions corresponding to each 30% region in a direction oriented from both ends to the other end are cooled to a temperature range of 500 to 600°C in a surface temperature (TE), and a central portion of a central 40% region corresponding to a region excluding both edge portions is cooled to a temperature range of 430 to 500°C in a surface temperature (TC), X = Nb / 93 + A / 48 + V / 51 / C / 12 + N / 14 A = Ti − 3.42 N − 1.5 S where [Nb], [V], [C], [N], [Ti] and [S] are wt% of each element T = Mn + 2.8 Mo + 1.5 Cr + 500 B where [Mn], [Mo], [Cr], and [B] are wt% of each element.

6. The method for manufacturing a steel sheet of claim 5, wherein the steel slab further includes, by wt%, at least one 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%.

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

8. The method for manufacturing a steel sheet of claim 5, wherein after the air cooling, an average temperature of the steel sheet is 550 to 650°C.

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