WO2024128710A1 - Steel sheet and method for manufacturing same - Google Patents

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

Steel plate and its manufacturing method

The present invention relates to a steel sheet and a manufacturing method thereof, and more specifically, to a high-strength steel sheet with excellent shear formability and a manufacturing method thereof.

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 with a relatively large amount of forming. This type of resistance compound hot-rolled steel sheet is an abnormal composite structure steel of ferrite-martensite, and exhibits continuous yield behavior and low yield strength characteristics due to the moving dislocation introduced during martensite transformation, and has excellent elongation and expansion formability. has

In this way, in order to improve elongation and stretch formability, Patent Documents 1 to 3 report hot rolling based on the Si-Mn and Mn-P-Cr composition system, then maintaining the ferrite transformation zone for several seconds and then measuring the martensite transformation start temperature (Ms). ) The following control method is used. In addition, Patent Document 4 uses a method of holding Si-Mn-Cr or Si-Mn-Cr-Mo in the ferrite transformation zone for several seconds and then winding at a temperature higher than the martensite transformation start temperature. In addition, Patent Document 5 proposed a technique for suppressing the formation of coarse carbonitrides when Ti, Nb, V, etc. are added to obtain improved high-strength steel.

However, in high-strength hot-rolled steel sheets with a high resistance compound ratio, alloying components such as Si, Al, Mn, Cr, and Mo, which are mainly used to manufacture ferrite-martensite abnormal composite structure steel with higher strength, are used in the slab after casting. Severe segregation occurs during molding, causing cracks and defects to form, worsening fatigue resistance and impact resistance. In addition, if the above alloy components are added excessively, the hot deformation resistance increases, and if Ti, Nb, V, and W are added together, the deformation resistance of the rolled sheet changes rapidly due to dynamic deformation organic precipitation during hot rolling. The shape quality becomes inferior, the microstructure becomes uneven, and eventually the physical properties of the final part also become inferior.

In addition, the above technologies propose a manufacturing method considering only the alloy components and the ratio of each component, even though the high-strength hot-rolled steel sheet with a high resistance compound ratio is a steel material that utilizes the moving dislocations formed at the boundary between the soft phase and the hard phase in the microstructure. There is no solution for cases where shear formability is poor. Here, shear formability is a process that is performed as the first step in the forming process of parts, and high-strength steel has a problem of cracks occurring during shear forming if the microstructure and uniformity of components in the thickness direction are poor. Such cracks can cause serious cracks during part molding or have a negative effect on durability during use.

[Prior art literature]

[Patent Document]

(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 Registration No. 4502897

(Patent Document 5) Korean Patent Publication No. 1543838

According to one embodiment of the present invention, it is intended to provide a steel plate and a method of manufacturing the same.

According to one embodiment of the present invention, the object is to provide a high-strength steel sheet with excellent shear formability and a manufacturing method thereof.

The object of the present invention is not limited to the above-described content. A person skilled in the art will have no difficulty in understanding the additional problems of the present invention from the overall content of the present specification.

According to one embodiment of the present invention, in weight percent, C: 0.030-0.150%, Si: 0.01-1.00%, Mn: 1.00-2.50%, Al: 0.01-0.80%, Cr: 0.005-0.500%, Mo: 0.005~0.300%, P: 0.001~0.050%, S: 0.001~0.010%, N: 0.001~0.010%, including the balance Fe and inevitable impurities,

The X value defined in equation 1 below is 0.010 to 0.200,

The T value defined in equation 2 below is 1.500 to 4.200,

The microstructure is expressed in area% and includes 30 to 70% of the hard phase containing bainite and martensite, 30 to 70% of the soft phase containing ferrite, and 3% or less of pearlite.

A steel sheet may be provided in which the hard phase has an average dislocation density of 2.0 to 3.0x10 ¹⁴ m ^-2 and the soft phase has an average dislocation density of 0.50 to 2.00x10 ¹⁴ m ^-2 .

[Relational Expression 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 percent of each element.)

[Relational Expression 2]

T = [Mn]+2.8[Mo]+1.5[Cr]+500[B]

(In the formula, [Mn], [Mo], [Cr], and [B] are the weight percent of each element.)

The steel sheet may further include, in weight percent, one or more selected from among Nb: 0.005-0.03%, Ti: 0.005-0.120%, V: 0.005-0.200%, and B: 0.0003-0.0030%.

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

When the steel plate is punched with a punching clearance of 5 to 20%, the shear surface may have 10 microcracks of 0.1 mm or more in length/cm ² or less, and the maximum crack length may be 1 mm or less.

According to one embodiment of the present invention, in weight percent, C: 0.030-0.150%, Si: 0.01-1.00%, Mn: 1.00-2.50%, Al: 0.01-0.80%, Cr: 0.005-0.500%, Mo: 0.005 to 0.300%, P: 0.001 to 0.050%, S: 0.001 to 0.010%, N: 0.001 to 0.010%, including the balance Fe and inevitable impurities, and the X value defined in equation 1 below is 0.010 to 0.200, and Reheating the steel slab with a T value of 1.500 to 4.200 defined in Equation 2;

hot rolling the reheated steel slab;

Primary cooling the steel sheet manufactured in the hot rolling step at an average cooling rate of 50 to 100° C./s to a temperature range of 430 to 600° C.;

Air cooling the primary cooled steel sheet for 4.0 to 10.0 seconds; and

Secondary cooling and winding 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,

During the first cooling, the surface temperature (TE) of the edge portion corresponding to 30% of the area from both ends to the other end, based on the width direction of the steel sheet, is in the temperature range of 500 to 600 ℃, excluding both edge portions. It is possible to provide a steel sheet manufacturing method in which the central portion of the central 40% area corresponding to is cooled to a temperature range of 430 to 500°C.

[Relational Expression 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 percent of each element.)

[Relational Expression 2]

T = [Mn]+2.8[Mo]+1.5[Cr]+500[B]

(In the formula, [Mn], [Mo], [Cr], and [B] are the weight percent of each element.)

The steel slab may further include one or more selected from among Nb: 0.005-0.030%, Ti: 0.005-0.120%, V: 0.005-0.200%, and B: 0.0003-0.0030% in weight percent.

The reheating step is performed at a temperature range of 1150 to 1350°C,

The hot rolling step can be performed at a finish rolling temperature of 850 to 1150°C.

After the air cooling step, the average temperature of the steel sheet may be 550 to 650°C.

According to an embodiment of the present invention, a steel plate and a manufacturing method thereof can be provided.

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

According to an embodiment of the present invention, a steel plate that can be used in cross members and subframes with relatively large amounts of forming among automobile chassis parts and a method for manufacturing the same can be provided.

Figures 1 (a) and (b) show the relationship between the number of occurrences by size of shear surface cracks in the invention example and the comparative example according to the punching clearance of 10% and 20%, respectively.

Below, preferred embodiments of the present invention will be described. Embodiments of the present invention may be modified in various forms, and the scope of the present invention should not be construed as limited to the embodiments described below. These embodiments are provided to explain the present invention in more detail to those skilled in the art.

The present invention studied a method to overcome the above-described conventional problems, prevent fractures from occurring in the cross section of the shear and punching molded parts, and prevent fatigue failure from occurring when the product is used. To this end, we studied ways to suppress the formation of microcracks and ensure smooth molding of products.

Accordingly, the inventor of the present invention investigated changes in shear formability and microcracks formed on the shear surface according to the characteristics of the composition and microstructure of steels with various alloy compositions and different microstructures. From the results, it was confirmed that shear formability is superior when the soft phase and hard phase constituting the microstructure satisfy a specific dislocation density range, rather than the composition ratio of the microstructure, and the present invention was completed.

Hereinafter, the present invention will be described in detail.

Below, the steel composition of the present invention will be described in detail.

In the present invention, unless otherwise specified, the % indicating the content of each element is based on weight.

The steel sheet according to an embodiment of the present invention has, in weight percent, C: 0.030-0.150%, Si: 0.01-1.00%, Mn: 1.00-2.50%, Al: 0.01-0.80%, Cr: 0.005-0.500%, 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 may contain the remainder of Fe and unavoidable impurities.

Carbon (C): 0.030~0.150%

Carbon (C) is the most economical and effective element in strengthening steel, and carbon (C) is the most economical and effective element in strengthening steel, and has a great influence on the hardness value and dislocation density of each composition. It's crazy. As the addition amount increases, the hardenability increases and the fraction of hard phases such as bainite and martensite in the microstructure increases, resulting in an increase in dislocation density and tensile strength. In addition, by forming fine precipitates with Ti and Nb, which have high affinity for carbon (C), the grain size becomes finer and the precipitation strengthening effect increases, resulting in an increase in both yield strength and tensile strength. If the carbon (C) content is less than 0.030%, it may be difficult to obtain a sufficient strengthening effect. According to one embodiment of the present invention, in order to stably secure a higher level of strength, it may be included in an amount of 0.050% or more. On the other hand, if the content exceeds 0.150%, the fraction of each phase, including bainite and martensite, increases, and the hardness value of the phases also increases, resulting in excessive strength increase, lower elongation and formability, and poor weldability. You can. In order to ensure more stable moldability in the present invention, carbon (C) may be included at 0.120% or less.

Silicon (Si): 0.01~1.00%

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 invention, 0.01% or more of silicon (Si) may be included to obtain the above-described effects. According to one embodiment of the present invention, it may be included at 0.10% or more. However, if the content exceeds 1.00%, red scale due to silicon (Si) is formed on the surface of the steel sheet during hot rolling, which not only deteriorates the surface quality of the steel sheet, but also reduces ductility and weldability. In one embodiment of the present invention, it may be included at 0.90% or less.

Manganese (Mn): 1.00~2.50%

Manganese (Mn), like Si, is an effective element 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, if the content is less than 1.00%, the above effect cannot be achieved by addition. According to one embodiment of the present invention, it may contain 1.40% or more of manganese (Mn). On the other hand, if the content exceeds 2.50%, the hardenability increases significantly, and the fraction and hardness value of each phase, including bainite and martensite, increase, which may lead to problems such as excessive increase in strength and deterioration in formability. In addition, when casting a slab in a continuous casting process, a large segregation area is developed in the center of the thickness, and when cooling after hot rolling, the microstructure in the thickness direction is formed unevenly, which may result in poor elongation flangeability. In particular, it may be difficult to manufacture the microstructure uniformly during cooling, even in the full length and full width of the hot-rolled sheet. In one embodiment of the present invention, the upper limit of the manganese (Mn) content may be limited to 2.30%.

Aluminum (Al): 0.01~0.80%

Aluminum (Al) is mainly added for deoxidation and has the effect of promoting ferrite transformation. If the content is less than 0.01%, the above-described addition effect may be insufficient. According to one embodiment of the present invention, it may contain 0.02% or more of aluminum (Al). On the other hand, if the content exceeds 0.80%, AlN is formed by combining with N, and corner cracks are likely to occur in the slab during continuous casting and defects due to inclusion formation are likely to occur. According to one embodiment of the present invention, the upper limit of aluminum (Al) content may be limited to 0.50%.

Chromium (Cr): 0.005~0.500%

Chromium (Cr) strengthens steel by solid solution and, when cooled, delays ferrite phase transformation and helps form bainite. However, if the chromium (Cr) content is less than 0.005%, the above effect cannot be achieved by addition. In order to more effectively secure the above effect, in the present invention, it may be included at 0.100% or more. On the other hand, if the content exceeds 0.500%, ferrite transformation is excessively delayed, resulting in poor elongation due to the formation of martensite. In addition, similar to Mn, the segregation zone at the center of the thickness develops significantly, and the microstructure in the thickness direction becomes non-uniform, which may result in poor elongation flangeability. According to one embodiment of the present invention, the upper limit can be limited to 0.300%.

Molybdenum (Mo): 0.005~0.300%

Molybdenum (Mo) increases the hardenability of steel and facilitates the formation of bainite structure. However, if the content is less than 0.005%, the above effect cannot be achieved by addition. In one embodiment of the present invention, the content may be 0.050% or more. On the other hand, if the molybdenum (Mo) content exceeds 0.300%, martensite may be formed due to an excessive increase in hardenability, resulting in a sharp deterioration in formability. In addition, it is economically disadvantageous and can be detrimental to weldability. In one embodiment of the present invention, the upper limit may be limited to 0.200%.

Phosphorus (P): 0.001~0.050%

Phosphorus (P), like Si, has both solid solution strengthening and ferrite transformation promotion effects. However, manufacturing at less than 0.001% requires a lot of manufacturing cost, which is economically disadvantageous and is insufficient to obtain strength, so the lower limit can be limited to 0.001%. On the other hand, if phosphorus (P) exceeds 0.050%, brittleness occurs due to grain boundary segregation, microcracks are likely to occur during molding, and ductility, elongation flangeability, and impact resistance characteristics can significantly deteriorate.

Sulfur (S): 0.001~0.010%

Sulfur (S) is an impurity present in steel. When its content exceeds 0.010%, it combines with Mn to form non-metallic inclusions. As a result, microcracks are likely to occur during cutting and processing of steel, and the extension flangeability and impact resistance are greatly improved. There is a problem with dropping it. In the present invention, sulfur (S) may be included in an amount of 0.005% or less. In the present invention, there is no particular limitation on the lower limit of the sulfur (S) content. However, in order to manufacture the content to less than 0.001%, a lot of time is required during steelmaking operation, which reduces productivity, so taking this into consideration, the lower limit of the sulfur (S) content is set. It can be limited to 0.001%.

Nitrogen (N): 0.001~0.010%

Nitrogen (N) is a representative solid solution strengthening element along with C and forms coarse precipitates together with Ti and Al. In general, the solid solution strengthening effect of nitrogen (N) is superior to that of carbon, but as the amount of nitrogen (N) increases in steel, there is a problem in that toughness decreases significantly. In addition, since manufacturing with a nitrogen (N) content of less than 0.001% requires a lot of time during steelmaking operations and reduces productivity, the lower limit can be limited to 0.001%.

The steel material of the present invention may contain remaining iron (Fe) and inevitable impurities in addition to the composition described above. Since unavoidable impurities may be unintentionally introduced during the normal manufacturing process, they cannot be excluded. Since these impurities are known to anyone skilled in the field of steel manufacturing, all of them are not specifically mentioned in this specification.

The steel sheet according to an embodiment of the present invention may further include 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%. .

Niobium (Nb): 0.005~0.030%

Niobium (Nb), along with Ti and V, is a representative precipitation strengthening element and is effective in improving the strength and impact toughness of steel through the effect of grain refinement by precipitating during hot rolling and delaying recrystallization. If the niobium (Nb) content is less than 0.005%, the above effect cannot be obtained. On the other hand, if the content exceeds 0.030%, there is a problem in that the stretched flangeability is inferior due to the formation of stretched grains and the formation of coarse composite precipitates due to excessive recrystallization delay during hot rolling.

Titanium (Ti): 0.005~0.120%

Titanium (Ti) is a representative precipitation strengthening element along with Nb and V, and forms coarse TiN in steel due to its strong affinity with N. TiN has the effect of suppressing grain growth during the heating process for hot rolling. In addition, titanium (Ti) remaining after reacting with nitrogen is dissolved in solid solution in the steel and combines with carbon to form TiC precipitates, which are useful ingredients for improving the strength of steel. If the titanium (Ti) content is less than 0.005%, the above effect cannot be achieved. On the other hand, if the content exceeds 0.120%, there is a problem of poor elongation flangeability during molding due to the generation of coarse TiN and coarsening of precipitates. In one embodiment of the present invention, the upper limit may be 0.105%. In one embodiment of the present invention, the upper limit may be 0.100%.

Vanadium (V): 0.005~0.200%

Vanadium (V) is a representative precipitation strengthening element along with Nb and Ti. It hardly precipitates during hot rolling and forms precipitates after coiling to improve 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. In order to achieve this effect in the present invention, the vanadium (V) content may be 0.005% or more. However, if the content is excessive, there is a problem that the elongation flangeability is inferior due to the formation of coarse precipitates and it is economically disadvantageous. Therefore, in the present invention, the upper limit can be limited to 0.200%, and in one embodiment, the upper limit can be limited to 0.150%.

Boron (B): 0.0003~0.0030%

When boron (B) exists in a solid solution state in steel, it is mainly segregated at grain boundaries and has the effect of improving the brittleness of steel by stabilizing grain boundaries. It also plays a role in suppressing the formation of coarse AlN nitride by stabilizing dissolved N. In addition, it is effective in the formation of hard phases bainite and martensite by delaying the ferrite phase transformation. In the present invention, boron (B) may be included in an amount of 0.0003% or more to ensure the above-described effects. On the other hand, if the content exceeds 0.0030%, the effect of addition no longer increases and ductility decreases, leading to poor formability. According to one embodiment of the present invention, boron (B) may be included in an amount of 0.0020% or less.

The steel plate according to an embodiment of the present invention may have an

[Relational Expression 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 percent of each element.)

[Relational Expression 2]

T = [Mn]+2.8[Mo]+1.5[Cr]+500[B]

(In the formula, [Mn], [Mo], [Cr], and [B] are the weight percent of each element.)

When the value of This may decrease. In addition, when the hot rolled steel sheet is cooled, there is a risk that the dissolved C and dissolved N atoms in the untransformed phase become insufficient, making it difficult to form a hard phase stably, and the grain boundaries become weak, resulting in 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 Ultimately, there is a risk that the elongation rate will become inferior. In one embodiment of the present invention, the value may be 0.030 or more. On the other hand, if the corresponding alloy element in relational equation 1 is not added, 0 can be substituted.

Equation 2 is a factorization of the combination of alloy elements that can maintain the formation of bainite, martensite, and MA phases, which are hard phases in the microstructure, at an appropriate level. As the T value defined in Equation 2 increases, the formation of hard phases such as bainite, martensite, and MA phases increases, and the hardness value of each hard phase may also increase. Therefore, in the present invention, for the desired strength, the T value can be limited to 1.500 or more. According to one embodiment of the present invention, it may be limited to 2.000 or more. On the other hand, the larger the T value, the more advantageous it is to secure strength, but if the value is excessive, the ductility of the steel decreases and the difference in hardness between the soft phase and the hard phase increases more than necessary, which may result in poor shear formability. Additionally, there is a problem in that material variation increases in the overall length and overall width of the hot rolled steel sheet. Therefore, in the present invention, the upper limit of the 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.

Below, the steel microstructure of the present invention will be described in detail.

In the present invention, unless specifically stated otherwise, the % indicating the fraction of microstructure is based on area.

The inventor of the present invention found that it is difficult to clearly distinguish whether the inherent shear formability of the steel is excellent and whether the results are stable despite changes in the punching clearance, just based on the area ratio of the microstructure. In particular, it was confirmed that dislocation density and physical characteristics vary greatly depending on the components that make up the steel.

Accordingly, as a result of research, the present inventor has confirmed that the geometrical dislocation density of the microstructure is an important factor affecting the occurrence of microcracks, which is the quality of the shear surface of steel, and proposes the present invention.

The microstructure of the steel sheet according to an embodiment of the present invention is, in terms of area percentage, 30 to 70% of the hard phase containing bainite and martensite, 30 to 70% of the soft phase containing ferrite, and 3% of pearlite. It may include the following.

In the present invention, the microstructure is controlled to ensure resistance compound ratio and extrusion formability. Accordingly, in the present invention, bainite and martensite are divided into hard phases, and ferrite is divided into soft phases, and their area fractions can be limited.

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

In the present invention, ferrite may include equiaxed ferrite and a ferritic low-temperature transformation phase. The ferritic low-temperature transformation phase may include acyclic ferrite, bainitic ferrite, granular bainitic ferrite, etc., and has non-uniform grain boundaries, a high dislocation density within grains, and a high density of low-angle grain boundaries within grains compared to equiaxed ferrite. It may mean ferrite.

In the present invention, the microstructure can be observed in a cross section perpendicular to the rolling direction of the steel sheet, and can be analyzed at a point of 1/4 to 1/2t (t is the thickness of the steel sheet) based on the thickness direction. Classification of microstructure and measurement of area fraction can be analyzed at a magnification of 3000 to 5000 using Electron Back Scattered Diffraction (EBSD, (JEOL JSM-1001F)).

When the area fraction of the hard phase exceeds 70%, the elongation rate is greatly reduced, the proportion of fractured areas in the shear section increases, and cracks with a length of 1 mm or more may significantly increase. In addition, the dependence of shear surface quality on changes in punching clearance increases, which may increase the occurrence of defects during actual part molding. On the other hand, if the area fraction is less than 30%, it may be difficult to secure 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 lead to a decrease in ductility, so taking this into account, the upper limit of the area fraction of martensite is set to 60%. It can be limited. Meanwhile, according to one embodiment of the present invention, martensite in the hard phase may be 0%.

In the present invention, a small amount of the MA phase may be observed, but it is considered martensite because its effect on the physical properties and cross-sectional quality of the punching part proposed in the present invention is not particularly different. In addition, tempered martensite, which contains fine carbides, has a lower dislocation density than martensite, which reduces the hardness value of the phase and helps improve the ductility of the overall steel. However, as the size of the formed carbides increases, it has negative effects such as brittleness. It can be distinguished from martensite. However, in the present invention, since the dislocation density of the phases is measured and described together, there is no need to distinguish between tempered martensite and it is regarded as martensite.

Since the soft phase can help the ductility of steel and the formation of fine precipitates, the lower limit of the area fraction can be limited to 30%. According to one embodiment of the present invention, the upper limit can be limited to 70%.

Meanwhile, as a structure other than the hard phase and soft phase, pearlite may be further included. However, if the area fraction of pearlite exceeds 3%, during shear forming of steel, it may correspond to a weak structure and increase cracks of 1 mm or more in length, so the upper limit can be limited to 3%.

In the steel sheet according to an embodiment of the present invention, the average dislocation density (Geometrical Necessary Dislocation) of the hard phase may be 2.0~3.0x10 ¹⁴ m ^-2 and the average dislocation density of the soft phase may be 0.50~2.00x10 ¹⁴ m ^-2 there is.

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 1/4 point in the thickness direction of the steel sheet with EBSD, using the formula below: can be calculated. For convenience, such calculations can be made using OIM analysis™ (EDAX), a software that analyzes the EBSD measurement results.

[ceremony]

(In the equation, θ is average misorientation (KAM values), u is unit length (step size in the EBSD measurement), and b is burgers vector.)

If the average dislocation density of the hard phase is less than 2.0x10 ¹⁴ m ^-2 , the strength may be greatly reduced, and if the value exceeds 3.0x10 ¹⁴ m ^-2 , ductility may decrease and the shear surface quality may be poor.

If the average dislocation density of the soft phase is less than 0.50x10 ¹⁴ m ^-2 , there is a risk that the strength may not meet the level required by the present invention, and burr generation may become severe during shear forming. On the other hand, if the value exceeds 2.00x10 ¹⁴ m ^-2 , there may be a problem of increased yield strength and decreased elongation.

Below, the steel sheet manufacturing method of the present invention will be described in detail.

The steel plate according to one embodiment of the present invention can be manufactured by reheating, hot rolling, primary cooling, air cooling, secondary cooling, and winding a steel slab that satisfies the alloy composition of the present invention.

reheat

Steel slabs satisfying the alloy composition of the present invention can be reheated to a temperature range of 1150 to 1350°C.

If the reheating temperature is less than 1150℃, the precipitates are not sufficiently re-dissolved, so the formation of precipitates decreases in the process after hot rolling, coarse TiN remains, and the tempering of the steel slab is not sufficient, so the temperature of the steel sheet during hot rolling increases. It may be difficult to control it consistently. On the other hand, if the temperature exceeds 1350°C, the strength may decrease due to abnormal grain growth of austenite grains.

hot rolling

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

During hot rolling, if the finish rolling temperature exceeds 1150°C, the temperature of the hot rolled steel sheet increases, the grain size becomes coarse, and the surface quality of the hot rolled steel sheet may deteriorate. On the other hand, if the temperature is less than 850°C, the anisotropy may worsen and formability may worsen due to the development of stretched crystal grains due to excessive recrystallization delay.

Primary cooling

The steel sheet manufactured in the hot rolling step can be first cooled to a temperature range of 430 to 600°C at an average cooling rate of 50 to 100°C/s. During the first cooling, the edge portion corresponding to 30% of the area from both ends to the other end based on the width direction of the steel sheet has a surface temperature (TE) in the temperature range of 500 to 600°C, corresponding to the area excluding both edge portions. The central part of the central 40% area can be cooled to a surface temperature (TC) of 430~500℃.

In the present invention, an area corresponding to 30% of each end from both ends toward the other end or center based on the width direction of the steel plate, that is, an area corresponding to a total of 60% of the entire steel plate, is divided into an edge portion, and the edge portion is The excluded central 40% area is divided into the central region.

In the present invention, the soft phase of the microstructure of steel may be formed during primary cooling and air cooling, and the hard phase may be formed during secondary cooling and winding. However, since the untransformed phase immediately before secondary cooling is formed into a hard phase after secondary cooling, a uniform soft phase must be formed at each width position of the steel sheet immediately after primary cooling and air cooling. Normally, when cooling a steel sheet, heat transfer proceeds quickly at the edge of the steel sheet, forming more hard phases compared to the central part. Therefore, in order to provide a uniform cooling rate for each position of the steel sheet width, the present invention seeks to control the cooling end temperature of the center portion and the edge portion differently. As proposed in the present invention, if the soft phase is uniformly formed at each width position during primary cooling and air cooling, the ratio of the hard phase formed during secondary cooling becomes constant, and the dependence on cooling conditions after winding is reduced. You can. When the soft phase and the hard phase are uniformly formed, excellent material uniformity and shear formability desired in the present invention can be secured.

During primary cooling, if the cooling rate is less than 50°C/s, too much ferrite fraction may be formed, and the average dislocation density may fall below the desired level, which may be disadvantageous in securing strength. On the other hand, if the cooling rate exceeds 100°C/s, during the first cooling, the surface temperature (TC) of the central part is excessively lowered, the ferrite fraction is greatly reduced, and the hard phase increases more than necessary, which may result in insufficient elongation.

Meanwhile, in the present invention, there is a risk that the difference in cooling rate depending on the width position of the steel sheet may become severe, so in order to prevent the hard phase from being formed differently depending on the width position, the edge portion and the center portion can be separated and cooled to different temperature ranges.

In particular, the present invention aims to ensure that the temperature of the steel sheet is reheated during air cooling in the subsequent process so that the steel sheet has a uniform temperature of 550 to 650°C. Accordingly, the temperature of the central portion needs to be subcooled below the desired temperature of 550 to 650°C, and it is desirable that the edge portion, which has a high cooling rate, be cooled to a higher temperature range than the central portion. Therefore, for this purpose, in the present invention, the surface temperature (TE) of the edge portion can be cooled to 500-600°C, and the surface temperature (TC) of the central portion can be cooled to 430-500°C.

If the surface temperature of the edge portion is less than 500°C, the formation of a soft phase may be insufficient, and there may be a problem of falling short of the desired temperature range of the hot rolled steel sheet during air cooling. On the other hand, if the temperature exceeds 600°C, there is a risk of exceeding the desired temperature range during air cooling, and there may be a problem of excessive formation of a soft phase in the final microstructure. According to an embodiment of the present invention, during primary cooling, the surface temperature of the edge portion may be 510°C or higher.

If the surface temperature of the central part is less than 430°C, there is a problem in that phase transformation of the hard phase bainite occurs. On the other hand, if the temperature exceeds 500°C, the effect of recuperation decreases, which may result in the temperature falling short of the desired temperature range.

air cooling

The primary cooled steel sheet may be air cooled for 4.0 to 10.0 seconds.

By stopping the intentional cooling of the primary cooled steel sheet, the temperature of the steel sheet can be restored to the desired temperature by internal latent heat and transformation heat. When the primarily cooled steel sheet is air-cooled, the steel sheet can recuperate to an average temperature of 550 to 650°C.

When air cooling, if the time is less than 4.0 seconds, the recuperation effect may not be effective. On the other hand, if the time exceeds 10.0 seconds, the ferrite fraction in the microstructure may greatly increase, and the hard phases bainite and martensite may decrease. In addition, pearlite structure and coarse carbides may be formed in the center of the thickness of the high temperature area of the steel sheet, resulting in poor cross-sectional quality after shear forming.

Secondary cooling and winding

The air-cooled steel sheet can be secondary cooled and wound at an average cooling rate of 10 to 100°C/s up to a temperature range of 50 to 200°C.

If the coiling temperature exceeds 200°C, the average dislocation density of the hard phase may be outside the range suggested by the present invention, making it difficult to secure strength. According to one embodiment of the present invention, the coiling temperature can be limited to 150°C or lower. On the other hand, if the temperature is less than 50°C, martensite is formed in an excessive amount, so that the average dislocation density of the hard phase exceeds the range proposed by the present invention, resulting in inferior elongation of the steel, and there may be a problem of corrosion of the steel sheet due to residual coolant. You can. According to one embodiment of the present invention, the coiling temperature can be limited to 70°C or higher.

If the cooling rate exceeds 100°C/s, the average dislocation density of the hard phase may become excessively high, resulting in a decrease in elongation. There is no particular limitation on the lower limit of the cooling rate, but in order to control the cooling rate to less than 10℃/s, the length of the cooling zone equipment must be long, and it is difficult to manufacture at the target coiling temperature of 200℃ or less. .

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%, there are ^no more than 10 microcracks with a length of 0.1 mm or more on the shear surface. , the maximum crack length is less than 1 mm, which ensures excellent strength and shear formability and low yield ratio.

Hereinafter, the present invention will be described in more detail through examples. However, it should be noted that the examples below are only for illustrating and explaining the present invention in more detail, and are not intended to limit the scope of the present invention.

(Example)

Steel slabs having the compositions shown in Table 1 below were manufactured into 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 sheet immediately after hot rolling was manufactured to be the same at 3.2 mm.

river
bell Alloy composition (% by weight) relational expression
One
(X) relational expression
2
(T) C Si Mn Cr Mo Nb Ti V Al P S N B A 0.075 0.70 1.70 0.200 0.020 0.030 0.100 0.005 0.05 0.019 0.001 0.005 0.0004 0.320 2.256 B 0.075 0.70 2.30 0.700 0.200 0.025 0.050 0.005 0.03 0.009 0.002 0.004 0.0010 0.162 4.410 C 0.075 0.70 1.20 0.005 0.001 0.020 0.060 0.005 0.04 0.010 0.002 0.004 0.0004 0.186 1.410 D 0.075 0.70 1.70 0.200 0.020 0.005 0.025 0.005 0.05 0.019 0.001 0.005 0.0004 0.043 2.256 E 0.075 0.70 2.20 0.300 0.150 0.020 0.060 0.001 0.03 0.010 0.003 0.004 0.0025 0.169 4.320 F 0.075 0.70 1.10 0.100 0.010 0.020 0.060 0.001 0.03 0.010 0.003 0.004 0.0003 0.169 1.428 G 0.095 0.25 1.40 0.400 0.010 0.005 0.030 0.005 0.10 0.010 0.003 0.004 0.0004 0.049 2.228 H 0.080 0.80 1.90 0.500 0.020 0.030 0.005 0.005 0.2 0.010 0.003 0.004 0.0004 0.021 2.906 I 0.060 0.20 1.80 0.100 0.200 0.020 0.050 0.005 0.03 0.010 0.003 0.004 0.0004 0.185 2.710 J 0.075 0.70 1.70 0.200 0.020 0.005 0.025 0.005 0.05 0.019 0.001 0.005 0.0004 0.043 2.256 K 0.110 0.08 1.45 0.400 0.010 0.005 0.020 0.005 0.20 0.010 0.002 0.004 0.0004 0.023 2.278 L 0.080 0.90 2.30 0.200 0.200 0.015 0.020 0.005 0.04 0.010 0.002 0.004 0.0020 0.047 4.160 M 0.150 0.90 2.20 0.200 0.200 0.015 0.020 0.005 0.04 0.010 0.002 0.004 0.0004 0.026 3.260

[Relational Expression 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 percent of each element.)

[Relational Expression 2]

T = [Mn]+2.8[Mo]+1.5[Cr]+500[B]

(In the formula, [Mn], [Mo], [Cr], and [B] are the weight percent of each element.)

Psalter
number steel grade hot rolling Primary cooling air cooling secondary cooling finishing rolling temperature
(℃) speed
(℃/s) Temperature (℃) hour
(candle) Steel plate temperature after air cooling
(℃) speed
(℃/s) Winding temperature
(℃) Edge part central part One A 874 63 538 486 5.5 601 55 95 2 B 865 74 522 447 5.7 590 49 110 3 C 860 72 530 488 6.2 605 53 78 4 D 873 25 642 574 6.2 640 45 85 5 D 866 98 470 402 6.0 558 56 89 6 D 890 65 570 490 12.0 643 55 92 7 D 874 66 583 485 12.0 668 48 115 8 D 855 62 566 476 2.5 525 50 88 9 D 885 68 545 488 5.8 594 112 26 10 D 881 65 553 496 6.2 607 85 27 11 D 866 71 558 482 6.5 603 48 220 12 E 875 62 533 462 4.8 610 48 85 13 F 890 55 575 472 5.2 605 55 105 14 G 873 65 520 489 6.9 604 55 105 15 H 880 71 538 495 6.5 625 62 88 16 I 879 77 554 470 6.3 589 53 110 17 J 887 68 549 492 5.8 595 58 75 18 K 892 62 577 460 6.2 611 60 92 19 L 885 79 571 462 4.8 593 53 90 20 M 870 65 562 458 5.8 608 49 82

Tables 3 and 4 below show the microstructure and mechanical properties of the manufactured steel sheets. The microstructure was observed in a cross section perpendicular to the rolling direction of the steel sheet, and was analyzed at 1/4 to 1/2 points based on the thickness direction. Electron Back Scattered Diffraction (EBSD, (JEOL JSM-1001F)) was used to classify and measure the area fraction of ferrite, ferritic low-temperature transformation phase, bainite, martensite, and pearlite formed in steel. Analyzed at ~5000 magnification. In addition, the average dislocation density (Geometrical Necessary Dislocation, GND) was measured using OIM analysis ^TM (EDAX) after EBSD measurement based on a cross section parallel to the rolling direction at 1/4 of the thickness direction of the steel sheet.

Mechanical properties were measured and expressed as yield strength, tensile strength, elongation at break, and yield ratio. Here, the 0.2% off-set yield strength (YS), tensile strength (TS), and elongation at break (T-El) are the results of testing by collecting JIS No. 5 standard test specimens in a direction perpendicular to the rolling direction. In addition, the physical properties were measured and shown by measuring the number of cracks in the punched cross-section by length. This was evaluated by punching a hole with a diameter of 10 mm. In this case, the number of micro cracks 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%. Each result value is the number of occurrences by crack length.

Psalter
number river
bell microstructure Properties division Fraction (area%) dislocation density
(10 ¹⁴ ,m ^-2 ) surrender
robbery
(MPa) Seal
robbery
(MPa) Destruction
elongation
(%) surrender fee soft phase Hard phase P soft phase Hard phase F BF B M One A 24 41 27 8 0 2.1 2.9 728 837 18 0.87 Comparative Example 1 2 B 0 27 62 11 0 1.9 3.3 1068 1289 5 0.83 Comparative example 2 3 C 33 59 0 8 0 0.9 1.8 517 765 21 0.68 Comparative Example 3 4 D 29 52 9 5 5 0.8 1.9 584 751 21 0.78 Comparative example 4 5 D 0 27 62 11 0 1.7 2.4 725 892 10 0.81 Comparative Example 5 6 D 23 42 23 8 4 1.5 2.3 667 768 18 0.87 Comparative Example 6 7 D 21 49 21 5 4 1.2 2.2 660 762 19 0.87 Comparative Example 7 8 D 5 23 32 40 0 1.3 2.8 796 950 9 0.84 Comparative example 8 9 D 19 38 13 30 0 1.7 3.1 680 843 14 0.81 Comparative Example 9 10 D 21 35 16 28 0 1.6 3.1 655 880 13 0.74 Comparative Example 10 11 D 23 33 24 20 0 1.5 1.9 723 812 17 0.89 Comparative Example 11 12 E 0 26 59 15 0 1.9 3.4 1013 1225 5 0.83 Comparative Example 12 13 F 28 60 9 3 0 0.8 1.8 504 750 22 0.67 Comparative Example 13 14 G 19 38 19 24 0 1.2 2.1 609 825 19 0.74 Invention Example 1 15 H 7 32 28 33 0 1.9 2.7 827 1056 12 0.78 Invention example 2 16 I 15 36 28 21 0 1.6 2.3 715 909 14 0.79 Invention Example 3 17 J 25 34 25 16 0 1.55 2.4 633 854 18 0.74 Invention Example 4 18 K 21 33 34 11 One 1.62 2.2 628 831 17 0.76 Invention Example 5 19 L 3 32 33 32 0 1.8 2.9 950 1209 7 0.79 Invention Example 6 20 M 5 30 27 38 0 1.75 2.8 975 1213 6 0.80 Invention Example 7

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

Psalter
number river
bell Properties division Number of cracks in the punched cross-section by length (pieces/cm ² ) Clearance 5% Clearance 10% Clearance 20% <0.1
(mm) 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) One A One 5 0 One 8 0 2 15 0 Comparative Example 1 2 B 3 8 One 4 11 2 8 13 2 Comparative example 2 3 C One 0 0 0 2 0 3 One 0 Comparative Example 3 4 D 2 3 2 One 3 One One 8 2 Comparative example 4 5 D 2 4 0 4 8 0 5 12 2 Comparative Example 5 6 D 2 4 0 0 2 One 2 7 2 Comparative Example 6 7 D 2 4 0 3 6 2 5 6 2 Comparative Example 7 8 D 5 7 0 4 6 2 6 9 2 Comparative example 8 9 D One 7 0 One 8 0 2 11 One Comparative Example 9 10 D One 3 0 One 2 One 5 5 2 Comparative Example 10 11 D 2 2 0 5 3 0 7 3 One Comparative Example 11 12 E 3 8 One 3 10 One 10 15 3 Comparative Example 12 13 F One 0 0 One One 0 2 2 0 Comparative Example 13 14 G 2 0 0 One 0 0 0 One 0 Invention Example 1 15 H One 3 0 One 3 0 3 5 0 Invention example 2 16 I One 0 0 2 2 0 0 One 0 Invention Example 3 17 J One 0 0 One One 0 2 One 0 Invention Example 4 18 K 0 0 0 0 0 0 2 One 0 Invention Example 5 19 L 2 2 0 2 2 0 5 5 0 Invention Example 6 20 M 3 3 0 2 3 0 4 6 0 Invention Example 7

As shown in Tables 3 and 4, in the case of the invention example that satisfies the alloy composition and manufacturing conditions of the present invention, the microstructure characteristics proposed in the present invention were satisfied, and the physical properties desired in the present invention were also secured.

Figures 1 (a) and (b) show the relationship between the number of occurrences by size of shear surface cracks in the invention example and the comparative example according to the punching clearance of 10% and 20%, respectively. Specifically, at punching clearances of 10% and 20%, it can be seen that the number of cracks in the comparative example was greater than that in the inventive example, and no cracks exceeding 1.0 mm in size occurred in the inventive example, and 1.0 It can be seen that the number of cracks smaller than mm is significantly less than in the comparative example.

On the other hand, Comparative Example 1 is an example that satisfies the content range of alloy elements proposed in the present invention, but does not satisfy the conditions of relational equation 1. As a result, the average dislocation density of the soft phase exceeded the range proposed in the present invention, and this was believed to be caused by increased formation of fine precipitates in the soft phase. No coarse cracks exceeding 1 mm in length were found in the punching area, but when the clearance was 20%, the occurrence of cracks with a length of 0.1 to 1.0 mm increased significantly. In addition, there is a problem of poor formability because the yield strength increases excessively due to work hardening during molding due to the high yield ratio.

Comparative Examples 2, 3, 12, and 13 are examples that did not satisfy Relation 2. Comparative Examples 2 and 12 contained an excessive amount of alloy components with high hardenability effect, and thus the strength was stably secured, but the elongation was insufficient. Because of this, the quality of the shear surface was also inferior. In Comparative Examples 3 and 13, the hard phase was not formed at the level proposed in the present invention due to the lack of alloy components with excellent hardenability effect, and as a result, the target strength was not secured. In addition, as the clearance increased, the occurrence of cracks became more severe, and cracks exceeding 1 mm in length were also confirmed.

Comparative Examples 4 and 5 are cases where the cooling end temperature during primary cooling after hot rolling is outside the range proposed in the present invention. In Comparative Example 4, during the first cooling, the end temperature exceeded the upper limit standard, so the formation of a hard phase was insufficient, and unnecessary pearlite was also formed, resulting in poor cross-sectional quality after punching. In Comparative Example 5, during the first cooling, the end temperature exceeded the lower limit standard and the soft phase fraction was insufficient, and the hard phase was excessively formed, resulting in poor cross-sectional quality after punching.

Comparative Examples 6 and 7 are cases where the air cooling time after primary cooling is outside the scope of the present invention. Due to the long exposure time in the high temperature range, the soft phase fraction increased significantly due to latent heat inside the steel sheet and heat generation due to phase transformation. Pearlite was also formed, so the cross-sectional quality after punching was inferior, and the yield ratio also exceeded the range proposed in the present invention. did. 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 in the present invention.

Comparative Example 8 is an example in which the air cooling time after primary cooling is below the range of the present invention. Secondary cooling occurred before the steel sheet was reheated, so the soft phase fraction fell below the desired level, and the hard phase fraction exceeded the desired range. As a result, the cross-sectional quality after punching was inferior.

Comparative Example 9 was a case in which the cooling rate was excessively fast during the secondary cooling, and it was overcooled and did not satisfy the coiling temperature range desired in the present invention. Because of this, the dislocation density of the hard phase exceeded the suggested range, and the cross-sectional quality after punching was inferior. It is believed that the main cause is the increase in the difference in physical properties between the soft phase and the hard phase.

Comparative Examples 10 and 11 are cases where the cooling end temperature during secondary cooling is outside the range suggested by the present invention. Comparative Example 10 is a case where the coiling temperature after secondary cooling was below the suggested temperature range, and the dislocation density of the hard phase was excessively high, resulting in poor cross-sectional quality after punching. Comparative Example 11 is a case where the coiling temperature after secondary cooling exceeded the suggested temperature range, and the dislocation density of the hard phase did not reach the suggested level. As a result, the yield ratio was excessively high and the cross-sectional quality was poor.

Although the present invention has been described in detail through examples above, other forms of embodiments are also possible. Therefore, the technical spirit and scope of the claims set forth below are not limited to the embodiments.

Claims (8)

In weight percent, C: 0.030~0.150%, Si: 0.01~1.00%, Mn: 1.00~2.50%, Al: 0.01~0.80%, Cr: 0.005~0.500%, Mo: 0.005~0.300%, P: 0.001~ 0.050%, S: 0.001~0.010%, N: 0.001~0.010%, including the balance Fe and inevitable impurities, The X value defined in equation 1 below is 0.010 to 0.200, The T value defined in equation 2 below is 1.500 to 4.200, The microstructure is expressed in area% and includes 30 to 70% of the hard phase containing bainite and martensite, 30 to 70% of the soft phase containing ferrite, and 3% or less of pearlite. A steel sheet in which the hard phase has an average dislocation density of 2.0~3.0x10 ¹⁴ m ^-2 and the soft phase has an average dislocation density of 0.50~2.00x10 ¹⁴ m ^-2 . [Relational Expression 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 percent of each element.) [Relational Expression 2] T = [Mn]+2.8[Mo]+1.5[Cr]+500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percent of each element.) In claim 1, The steel sheet further contains one or more selected from the group consisting of Nb: 0.005-0.030%, Ti: 0.005-0.120%, V: 0.005-0.200%, and B: 0.0003-0.0030% in weight percent. In claim 1, The steel plate has a tensile strength of 780 MPa or more and a yield ratio of 0.70 to 0.85. In claim 1, The steel sheet is a steel sheet in which, when punching formed with a punching clearance of 5 to 20%, the shear surface has 10 microcracks with a length of 0.1 mm or more/cm ² or less, and the maximum crack length is 1 mm or less. In weight percent, C: 0.030~0.150%, Si: 0.01~1.00%, Mn: 1.00~2.50%, Al: 0.01~0.80%, Cr: 0.005~0.500%, Mo: 0.005~0.300%, P: 0.001~ 0.050%, S: 0.001-0.010%, N: 0.001-0.010%, including the balance Fe and inevitable impurities, the X value defined in equation 1 below is 0.010-0.200, and the T value defined in equation 2 below is 1.500 Reheating the steel slab to ~4.200; hot rolling the reheated steel slab; Primary cooling the steel sheet manufactured in the hot rolling step at an average cooling rate of 50 to 100° C./s to a temperature range of 430 to 600° C.; Air cooling the primary cooled steel sheet for 4.0 to 10.0 seconds; and Secondary cooling and winding 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, During the first cooling, the surface temperature (TE) of the edge portion corresponding to 30% of the area from both ends to the other end, based on the width direction of the steel sheet, is in the temperature range of 500 to 600 ℃, excluding both edge portions. A steel sheet manufacturing method in which the central part of the central 40% area is cooled to a temperature range of 430 to 500 ℃. [Relational Expression 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 percent of each element.) [Relational Expression 2] T = [Mn]+2.8[Mo]+1.5[Cr]+500[B] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percent of each element.) In claim 5, The steel slab is a steel plate manufacturing method further comprising at least one selected from the group consisting of Nb: 0.005-0.030%, Ti: 0.005-0.120%, V: 0.005-0.200%, and B: 0.0003-0.0030%, in weight percent. In claim 5, The reheating step is performed at a temperature range of 1150 to 1350°C, A steel sheet manufacturing method in which the hot rolling step is performed at a finish rolling temperature of 850 to 1150°C. In claim 5, A method of manufacturing a steel sheet where the average temperature of the steel sheet after the air cooling step is 550 to 650°C.

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