Steel sheet and method for manufacturing same

Abstract

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

Description

Technical Field

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

[0002] In existing chassis components, cross members and subframes with a large amount of forming mainly use low yield ratio high strength hot rolled steel plates with a tensile strength of 540 - 590 MPa. The low yield ratio hot rolled steel plate is a ferritic-martensitic dual-phase composite structure steel. During martensitic transformation, due to the introduced slip dislocations, it exhibits continuous yielding behavior and low yield strength characteristics, and has excellent elongation and bulge formability.

[0003] As described above, in order to improve elongation and bulge formability, Patent Documents 1 to 3 utilize a method based on the Si-Mn, Mn-P-Cr composition system, maintaining for several seconds in the ferrite transformation region after hot rolling and then controlling to below the martensite transformation start temperature (Ms). In addition, Patent Document 4 utilizes a method of coiling at a temperature above the martensite transformation start temperature after maintaining for several seconds in the ferrite transformation region by using Si-Mn-Cr or Si-Mn-Cr-Mo systems. In addition, Patent Document 5 proposes a scheme for suppressing the formation of coarse carbonitrides when adding Ti, Nb, V, etc. to obtain an improved high strength steel technology.

[0004] However, in high strength hot rolled steel plates with a low yield ratio, alloying elements such as Si, Al, Mn, Cr, Mo, etc., which are mainly used to manufacture ferritic-martensitic dual-phase composite structure steels with higher strength, will cause severe segregation in the slab after casting, forming cracks or defects during the forming process, resulting in poor fatigue resistance and impact resistance characteristics. In addition, when the addition amount of the above alloying elements is too high, the hot deformation resistance increases. When adding Ti, Nb, V, and W at the same time, during the hot rolling process, due to dynamic deformation-induced precipitation, there is a sharp change in the deformation resistance, the shape quality of the rolled plate becomes poor, and the fine microstructure becomes uneven. As a result, the physical properties of the final components will also become poor.

[0005] In addition, the above technologies only propose manufacturing methods considering alloying elements and the proportion of each constituent phase. High strength hot rolled steel plates with a low yield ratio are steels that utilize mobile dislocations formed at the grain boundaries of soft and hard phases in the fine microstructure, and there is no solution for the poor actual shear formability. Among them, for shear formability, when the uniformity of the fine microstructure and composition in the thickness direction of high strength steel during the process which is the first step in the forming process of components is poor, there is a problem of crack generation during shear forming. The cracks may cause serious cracks during the component forming process or have an adverse impact on durability during use.

[0006] [Prior Art Documents]

[0007] [Patent Documents]

[0008] (Patent Document 1) Japanese Published Patent Gazette No. 1995 - 278731

[0009] (Patent Document 2) Japanese Published Patent Gazette No. 1997 - 241790

[0010] (Patent Document 3) Japanese Published Patent Gazette No. 1994 - 049591

[0011] (Patent Document 4) US Patent Granted Gazette No. 4502897

[0012] (Patent Document 5) Korean Patent Granted Gazette No. 1543838 Summary of the Invention

[0013] (1) Technical Problem to be Solved

[0014] According to an embodiment of the present invention, an object is to provide a steel sheet and a method for manufacturing the same.

[0015] According to an embodiment of the present invention, an object is to provide a low yield ratio high strength steel sheet with excellent shear formability and a method for manufacturing the same.

[0016] The technical problems of the present invention are not limited to the above. Those skilled in the technical field to which the present invention pertains can easily understand additional technical problems of the present invention from the overall content of this specification.

[0017] (2) Technical Solution

[0018] According to an embodiment of the present invention, a steel sheet can be provided which, by weight%, comprises: 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%, the balance being Fe and inevitable impurities, the value of X defined in the following relational expression 1 being 0.010 to 0.200, the value of T defined in the following relational expression 2 being 1.500 to 4.200, and by area%, the fine structure comprising 30 - 70% of a hard phase containing bainite and martensite, 30 - 70% of a soft phase containing ferrite, and 3% or less of pearlite, the average dislocation density of the hard phase being 2.0×10 14 m -2 to 3.0×1014 m -2 , the average dislocation density of the soft phase is 0.50×10 14 m -2 to 2.00×10 14 m -2 .

[0019] [Relationship 1]

[0020] X = ([Nb] / 93 + A / 48 + [V] / 51) / ([C] / 12 + [N] / 14)

[0021] A = [Ti] - 3.42[N] - 1.5[S]

[0022] (In the formula, [Nb], [V], [C], [N], [Ti] and [S] are the weight percentages of each element.)

[0023] [Relationship 2]

[0024] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B]

[0025] (In the formula, [Mn], [Mo], [Cr] and [B] are the weight percentages of each element.)

[0026] By weight, the steel plate may further contain one or more selected from 0.005 - 0.03% of Nb, 0.005 - 0.120% of Ti, 0.005 - 0.200% of V, and 0.0003 - 0.0030% of B.

[0027] The tensile strength of the steel plate may be 780 MPa or more, and the yield ratio may be 0.70 to 0.85.

[0028] When the steel plate is stamped with a stamping gap of 5 - 20%, the number of microcracks with a length of 0.1 mm or more on the shear surface may be 10 per cm 2 or less, and the length of the maximum crack may be 1 mm or less.

[0029] According to an embodiment of the present invention, a method for manufacturing a steel plate can be provided. The manufacturing method includes the following steps: reheating a steel slab, which, by weight percentage, contains: 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%, the balance being Fe and unavoidable impurities, the X value defined in the following relational expression 1 being 0.010 to 0.200, and the T value defined in the following relational expression 2 being 1.500 to 4.200; hot rolling the reheated steel slab; performing primary cooling on the steel plate manufactured in the hot rolling step at an average cooling rate of 50 - 100 °C / second to a temperature range of 430 - 600 °C; air-cooling the steel plate after the primary cooling for 4.0 - 10.0 seconds; and performing secondary cooling on the air-cooled steel plate at an average cooling rate of 10 - 100 °C / second to a temperature range of 50 - 200 °C and coiling it. Wherein, during the primary cooling, based on the width direction of the steel plate, the surface temperature (TE) of the edge portions of the regions corresponding to 30% from both ends along the direction of the other end is cooled to a temperature range of 500 - 600 °C, and the surface temperature (TC) of the central portions of the regions corresponding to 40% of the center of the region other than the two edge portions is cooled to a temperature range of 430 - 500 °C.

[0030] [Relational expression 1]

[0031] X = ([Nb] / 93 + A / 48 + [V] / 51) / ([C] / 12 + [N] / 14)

[0032] A = [Ti] - 3.42[N] - 1.5[S]

[0033] (In the formula, [Nb], [V], [C], [N], [Ti], and [S] are the weight percentages of the respective elements.)

[0034] [Relational expression 2]

[0035] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B]

[0036] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of the respective elements.)

[0037] By weight percentage, the steel billet may further contain one or more selected from 0.005 - 0.030% of Nb, 0.005 - 0.120% of Ti, 0.005 - 0.200% of V, and 0.0003 - 0.0030% of B.

[0038] The reheating step may be carried out in the temperature range of 1150 - 1350 °C, and the hot rolling step may be carried out at a finish rolling temperature of 850 - 1150 °C.

[0039] The average temperature of the steel plate after the air cooling step may be 550 - 650 °C.

[0040] (III) Beneficial Effects

[0041] According to an embodiment of the present invention, a steel plate and a manufacturing method thereof can be provided.

[0042] According to an embodiment of the present invention, a low yield ratio high strength steel plate with excellent shearing formability and a manufacturing method thereof can be provided.

[0043] According to an embodiment of the present invention, a steel plate that can be used for cross members and subframes with a large forming amount in automotive chassis components and a manufacturing method thereof can be provided. Description of the Drawings

[0044] Figure 1 (a) and (b) thereof respectively show a relationship diagram of the occurrence times of each shear surface crack size of the inventive example and the comparative example according to a stamping gap of 10% and a stamping gap of 20%. Best Mode for Carrying Out the Invention

[0045] Hereinafter, preferred specific embodiments of the present invention will be described. The specific embodiments of the present invention can be modified into various forms, and it should not be construed that the scope of the present invention is limited to the specific embodiments described below. These specific embodiments are provided to more specifically explain the present invention to those skilled in the technical field to which the present invention pertains.

[0046] The present invention has studied a solution to overcome the above-mentioned existing problems, where no fracture occurs in the cross-section of the shearing and stamping forming parts, and fatigue failure during the use of the product is prevented. For this purpose, a solution to suppress the formation of microcracks and enable the product to be successfully formed has been studied.

[0047] Therefore, the inventors of the present invention studied the shear formability according to the characteristics of the composition and microstructure and the changes in the microcracks formed on the shear surface for various steels with different alloy compositions and microstructures. As a result, it was confirmed that when the soft and hard phases constituting the microstructure satisfy specific dislocation density ranges, respectively, rather than the composition ratio of the microstructure, excellent shear formability is obtained, and thus the present invention was completed.

[0048] Hereinafter, the present invention will be described in detail.

[0049] Hereinafter, the composition of the steel of the present invention will be described in detail.

[0050] In the present invention, unless otherwise specifically stated, % representing the content of each element is based on weight.

[0051] By weight %, the steel sheet according to an embodiment of the present invention may contain: 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%, the balance being Fe and inevitable impurities.

[0052] Carbon (C): 0.030 - 0.150%

[0053] Carbon (C) is the most economical and effective element for strengthening steel and is the most economical and effective element for strengthening the hardness values of the respective constituent phases and carbon (C) steel, and has a great influence on the hardness values and dislocation densities of the respective constituent phases. When the addition amount of carbon (C) increases, due to the increase in hardenability, the fraction of hard phases such as bainite and martensite in the microstructure increases, and the dislocation density and tensile strength increase. In addition, carbon (C) forms fine precipitates together with Ti and Nb having high affinity, the grain size becomes fine, the precipitation strengthening effect also increases, and at the same time the yield strength and tensile strength increase. 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 invention, in order to stably ensure a higher level of strength, carbon (C) of 0.050% or more may be contained. On the other hand, when the content of carbon (C) exceeds 0.150%, the fraction of each phase including bainite and martensite increases, the hardness value of the phase also increases, the strength rises excessively, there are problems of reduced elongation and formability, and the weldability may also be poor. In the present invention, in order to ensure more stable formability, carbon (C) of 0.120% or less may be contained.

[0054] Silicon (Si): 0.01 - 1.00%

[0055] Silicon (Si) deoxidizes the molten steel and has a solution strengthening effect, delaying the formation of coarse carbides, and thus is beneficial to improving formability. In the present invention, in order to obtain the above effects, silicon (Si) of 0.01% or more can be included. According to an embodiment of the present invention, silicon (Si) of 0.10% or more can be included. However, when the content of silicon (Si) exceeds 1.00%, during hot rolling, red scale caused by silicon (Si) is formed on the surface of the steel plate, not only the surface quality of the steel plate becomes very poor, but also there may be problems of reduced ductility and weldability. As an embodiment of the present invention, silicon (Si) of 0.90% or less can be included.

[0056] Manganese (Mn): 1.00 - 2.50%

[0057] Like Si, manganese (Mn) is an effective element for solution strengthening of steel, increasing the hardenability of steel, and it is easy to form bainite and martensite as hard phases during the cooling process after hot rolling. However, when the content of manganese (Mn) is less than 1.00%, the above effects brought by the addition cannot be obtained. According to an embodiment of the present invention, manganese (Mn) of 1.40% or more can be included. On the other hand, when the content of manganese (Mn) exceeds 2.50%, the hardenability increases significantly, the fractions of each phase including bainite and martensite and the hardness values of the phases increase, and there may be problems of excessive strength increase and reduced formability. In addition, when casting a slab in the continuous casting process, the segregation part in the thickness center part develops significantly, and the fine microstructure along the thickness direction is formed unevenly during cooling after hot rolling, so the flange stretching property may become poor. In particular, it may be difficult to manufacture a uniform fine microstructure even over the entire length and entire width of the hot rolled sheet during cooling. As an embodiment of the present invention, the upper limit of the content of manganese (Mn) can be limited to 2.30%.

[0058] Aluminum (Al): 0.01 - 0.80%

[0059] Aluminum (Al) is an added component mainly for deoxidation and has the effect of promoting ferrite phase transformation. When the content of aluminum (Al) is less than 0.01%, the above addition effects may be insufficient. According to an embodiment of the present invention, aluminum (Al) of 0.02% or more can be included. On the other hand, when the content of aluminum (Al) exceeds 0.80%, aluminum (Al) combines with N to form AlN, and corner cracks are likely to occur on the slab during continuous casting, and defects are likely to occur due to the formation of inclusions. According to an embodiment of the present invention, the upper limit of the content of aluminum (Al) can be limited to 0.50%.

[0060] Chromium (Cr): 0.005 - 0.500%

[0061] Chromium (Cr) plays a role in solid-solution strengthening of steel and delaying the ferrite phase transformation during cooling, thus contributing to the formation of bainite. However, when the content of chromium (Cr) is less than 0.005%, the above effects brought by the addition cannot be obtained. In order to more effectively ensure the above effects, in the present invention, chromium (Cr) can be included in an amount of 0.100% or more. In addition, when the content of chromium (Cr) exceeds 0.500%, the ferrite phase transformation is overly delayed, resulting in the formation of martensite and a deterioration in elongation. Furthermore, similar to Mn, the segregation part in the thickness center is significantly developed, and the microstructure in the thickness direction is non-uniform, so the flange stretchability may deteriorate. According to an embodiment of the present invention, the upper limit of chromium (Cr) can be limited to 0.300%.

[0062] Molybdenum (Mo): 0.005 - 0.300%

[0063] Molybdenum (Mo) increases the hardenability of steel, thus promoting the formation of bainite structure. However, when the content of molybdenum (Mo) is less than 0.005%, the above effects brought by the addition cannot be obtained. As an embodiment of the present invention, the content of molybdenum (Mo) can be 0.050% or more. On the other hand, when the content of molybdenum (Mo) exceeds 0.300%, the hardenability increases excessively, resulting in the formation of martensite, so the formability may rapidly deteriorate. In addition, it is economically disadvantageous and may also damage the weldability. As an embodiment of the present invention, the upper limit of the content of molybdenum (Mo) can be limited to 0.200%.

[0064] Phosphorus (P): 0.001 - 0.050%

[0065] Similar to Si, phosphorus (P) has both the effects of solid-solution strengthening and promoting ferrite phase transformation. However, in order to make the content of phosphorus (P) less than 0.001%, high manufacturing costs are required, so it is economically disadvantageous and the strength obtained is also insufficient. Therefore, the lower limit of the content of phosphorus (P) can be limited to 0.001%. On the other hand, when phosphorus (P) exceeds 0.050%, brittleness is generated due to grain boundary segregation, and fine cracks are likely to occur during forming, and the ductility, flange stretchability, and impact resistance characteristics may deteriorate significantly.

[0066] Sulfur (S): 0.001 - 0.010%

[0067] Sulfur (S) is an impurity present in steel. When the content of sulfur (S) exceeds 0.010%, it combines with Mn etc. to form non-metallic inclusions. Therefore, during the cutting process of steel, fine cracks are likely to occur, and there are problems of significantly reducing the flange elongation and impact resistance. In the present invention, sulfur (S) with a content of 0.005% or less can be included. In the present invention, there is no particular limitation on the lower limit of the content of sulfur (S). When the content of sulfur (S) is made less than 0.001%, a large amount of time is required during the steelmaking operation, resulting in a reduction in productivity. Therefore, considering this point, the lower limit of the content of sulfur (S) can be limited to 0.001%.

[0068] Nitrogen (N): 0.001 - 0.010%

[0069] Nitrogen (N), together with C, is a representative solid-solution strengthening element, and forms coarse precipitates together with Ti, Al, etc. Generally, the solid-solution strengthening effect of nitrogen (N) is superior to that of carbon, but there is a problem that the toughness is significantly reduced as the amount of nitrogen (N) in the steel increases. In addition, in order to make the content of nitrogen (N) less than 0.001%, a large amount of time is required during the steelmaking operation, resulting in a reduction in productivity. Therefore, the lower limit of nitrogen (N) can be limited to 0.001%.

[0070] In the steel of the present invention, in addition to the above components, the balance of iron (Fe) and inevitable impurities can be included. Inevitable impurities may be inadvertently mixed in during the conventional manufacturing process, so these impurities cannot be excluded. These impurities are well-known to those skilled in the art of conventional steel manufacturing, so not all of their contents are specifically described in this specification.

[0071] The steel plate according to an embodiment of the present invention may further include one or more selected from 0.005 - 0.030% of Nb, 0.005 - 0.120% of Ti, 0.005 - 0.200% of V, and 0.0003 - 0.0030% of B.

[0072] Niobium (Nb): 0.005 - 0.030%

[0073] Niobium (Nb), together with Ti and V, is a representative precipitation strengthening element and precipitates during the hot rolling process. Due to the grain refinement effect brought about by the delay of recrystallization, it is effective in improving the strength and impact toughness of steel. When the content of niobium (Nb) is less than 0.005%, the above effects cannot be obtained. On the other hand, when the content of niobium (Nb) exceeds 0.030%, due to excessive recrystallization delay during the hot rolling process, elongated grains and coarse composite precipitates are formed, resulting in a problem of poor flange elongation.

[0074] Titanium (Ti): 0.005 - 0.120%

[0075] Titanium (Ti), like Nb and V, is a representative precipitation strengthening element. Due to its strong affinity with N, coarse TiN is formed in steel. TiN has the effect of suppressing grain growth during the heating process for hot rolling. In addition, the remaining titanium (Ti) after reacting with nitrogen dissolves in the steel and combines with carbon to form TiC precipitates. Therefore, titanium (Ti) is a useful component for increasing the strength of steel. When the content of titanium (Ti) is less than 0.005%, the above effects cannot be obtained. On the other hand, when the content of titanium (Ti) exceeds 0.120%, coarse TiN is produced, the precipitates become coarse, and there is a problem of poor flange stretchability during forming. As an embodiment of the present invention, the upper limit of the content of titanium (Ti) can be 0.105%. As an embodiment of the present invention, the upper limit of titanium (Ti) can be 0.100%.

[0076] Vanadium (V): 0.005 - 0.200%

[0077] Vanadium (V), together with Nb and Ti, is a representative precipitation strengthening element. It hardly precipitates during hot rolling and forms precipitates after coiling, thereby playing a role in increasing the strength of steel. Therefore, due to the delay of recrystallization during hot rolling, the additional strength is effectively increased without increasing the deformation resistance and rolling load. In the present invention, in order to obtain this effect, the content of vanadium (V) can be 0.005% or more. However, when the content of vanadium (V) is too high, there is a problem of poor flange stretchability due to the formation of coarse precipitates, which is also economically disadvantageous. Therefore, in the present invention, the upper limit of the content of vanadium (V) is limited to 0.200%. As an embodiment, the upper limit of the content of vanadium (V) can be limited to 0.150%.

[0078] Boron (B): 0.0003 - 0.0030%

[0079] When boron (B) exists in a solid solution state in steel, it mainly segregates at grain boundaries, has the effect of stabilizing grain boundaries to improve the brittleness of steel, and plays a role in stabilizing solid solution N to inhibit the formation of coarse AlN nitrides. In addition, by delaying the ferrite phase transformation, bainite and martensite as hard phases are effectively formed. In the present invention, in order to ensure the above effects, boron (B) of 0.0003% or more can be included. On the other hand, when the content of boron (B) exceeds 0.0030%, the effects brought by the addition no longer increase, the ductility decreases, and there is a problem of poor formability. According to an embodiment of the present invention, boron (B) of 0.0020% or less can be included.

[0080] According to an embodiment of the present invention, the X value defined in the following relational expression 1 of the steel plate can be 0.010 to 0.200, and the T value defined in the following relational expression 2 can be 1.500 to 4.200.

[0081] [Relationship 1]

[0082] X = ([Nb] / 93 + A / 48 + [V] / 51) / ([C] / 12 + [N] / 14)

[0083] A = [Ti] - 3.42[N] - 1.5[S]

[0084] (In the formula, [Nb], [V], [C], [N], [Ti] and [S] are the weight percentages of each element.)

[0085] [Relationship 2]

[0086] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B]

[0087] (In the formula, [Mn], [Mo], [Cr] and [B] are the weight percentages of each element.)

[0088] When the X value defined in Relationship 1 exceeds 0.200, the formation of precipitates increases and the strength slightly increases. However, during hot rolling, due to the delay in recrystallization, a tensile fine structure is likely to form in the rolling direction, and the elongation in the direction perpendicular to rolling may decrease. In addition, when the hot-rolled steel plate is cooled, there is a shortage of dissolved C and dissolved N atoms in the untransformed phase, making it difficult to stably form hard phases, resulting in poor grain boundaries and possibly poor shear surface quality. As an embodiment of the present invention, the X value can be 0.180 or less. On the other hand, when the X value is less than 0.010, the grains are likely to grow during the reheating process, and the recrystallization becomes uneven during hot rolling, forming locally coarse grains. The dissolved C and dissolved N are excessive above the required level, and the hardness value of the hard phase shows a tendency to increase, and the final elongation may deteriorate. As an embodiment of the present invention, the X value can be 0.030 or more. In addition, when the corresponding alloying elements in Relationship 1 are not added, 0 can be substituted in.

[0089] Relationship 2 is the factorization of the combination of alloying elements that can keep the formation of bainite, martensite, and MA phase as hard phases in the fine microstructure at an appropriate level. As the value of T defined in Relationship 2 increases, the formation of bainite, martensite, and MA phase as hard phases increases, and the hardness values of the respective hard phases may also increase. Therefore, in the present invention, for the desired strength, the value of T can be limited to 1.500 or more. According to an embodiment of the present invention, the value of T can be limited to 2.000 or more. Additionally, the larger the value of T, the more beneficial it is to ensure strength. However, when the value of T is too high, the ductility of the steel decreases, and the hardness difference between the soft phase and the hard phase increases above the required level, so the shear formability may deteriorate. Furthermore, there is a problem of increased material deviation over the entire length and width of the hot-rolled steel sheet. Therefore, in the present invention, the upper limit of the value of T can be limited to 4.200. According to an embodiment of the present invention, the upper limit of the value of T can be limited to 4.000.

[0090] Next, the fine microstructure of the steel of the present invention will be described in detail.

[0091] In the present invention, unless otherwise specifically stated, % representing the fraction of the fine microstructure is based on area.

[0092] The inventors of the present invention found that it is difficult to clearly distinguish to what extent the inherent shear formability of steel is excellent and whether stable results can be obtained even when the stamping gap changes only based on the area ratio of the constituent phases of the fine microstructure. In particular, it was confirmed that the dislocation density and physical characteristics vary significantly depending on the composition of the steel.

[0093] Therefore, the inventors of the present invention confirmed through research results that the dislocation density (Geometrical Necessary Dislocation) of the fine microstructure is an important factor affecting the generation of microcracks that are the quality of the shear surface of the steel, and thus proposed the present invention.

[0094] By area %, the fine microstructure of the steel sheet according to an embodiment of the present invention may include 30 - 70% of a hard phase containing bainite and martensite, 30 - 70% of a soft phase containing ferrite, and 3% or less of pearlite.

[0095] In the present invention, a low yield ratio and bulge formability are ensured by controlling the fine microstructure. Therefore, in the present invention, bainite and martensite are classified as hard phases, ferrite is classified as a soft phase, and their area fractions can be limited.

[0096] In the present invention, bainite may include upper bainite and lower bainite, and fine carbides are formed in the lath structure, thereby being distinguishable from the low-temperature transformation phase of the ferrite system.

[0097] In the present invention, the ferrite may include equiaxed ferrite and low-temperature transformation phases of the ferrite system. The low-temperature transformation phases of the ferrite system may include acicular ferrite, bainitic ferrite, granular bainitic ferrite, etc., and may refer to ferrite with uneven grain boundaries, a high dislocation density in the grains, and a high density of small-angle grain boundaries in the grains compared to equiaxed ferrite.

[0098] The microstructure in the present invention can be observed on a cross-section perpendicular to the rolling direction of the steel plate, and analysis can be performed at a position of 1 / 4 - 1 / 2t (t is the thickness of the steel plate) based on the thickness direction. The differentiation of the microstructure and the measurement of the area fraction can be analyzed at a magnification of 3000 - 5000 using Electron BackScattered Diffraction (EBSD, (JEOL JSM-1001F)).

[0099] When the area fraction of the hard phase exceeds 70%, the elongation rate decreases significantly. In the shear plane, the proportion of the fracture part increases, and cracks with a length of 1 mm or more may increase significantly. In addition, the dependence of the shear plane quality on the change in the stamping gap increases, and defects may increase during actual component forming. On the other hand, when the area fraction of the hard phase is less than 30%, it may be difficult to ensure the desired strength. Additionally, according to an embodiment of the present invention, martensite is a relatively harder phase compared to bainite. Therefore, an increase in martensite may lead to a decrease in ductility. Considering this, the upper limit of the area fraction of martensite can be limited to 60%. Additionally, according to an embodiment of the present invention, the martensite in the hard phase can be 0%.

[0100] In the present invention, a small amount of MA phase can be observed, but it does not have a special impact on the physical properties and the quality of the stamping part cross-section proposed in the present invention, so it is considered martensite. In addition, tempered martensite containing fine carbides has a lower dislocation density compared to martensite, so the hardness value of the phase decreases. Although it helps to improve the ductility of the entire steel, when the size of the formed carbides increases, adverse effects such as brittleness may occur, and it can usually be distinguished from martensite. However, in the present invention, since the dislocation density of the phase is measured and described together, there is no need to distinguish tempered martensite, so it is regarded as martensite.

[0101] Since the soft phase can contribute to the ductility of the steel and the formation of fine precipitates, the lower limit of the area fraction of the soft phase can be limited to 30%. According to an embodiment of the present invention, the upper limit of the area fraction of the soft phase can be limited to 70%.

[0102] In addition, as a structure other than the hard phase and the soft phase, pearlite may be included. However, when the area fraction of pearlite exceeds 3%, during the shearing forming of the steel, it is equivalent to a weak structure, and cracks with a length of 1 mm or more may increase. Therefore, the upper limit of the area fraction of pearlite can be limited to 3%.

[0103] The average dislocation density of the hard phase of the steel plate according to an embodiment of the present invention may be 2.0×10 14 m -2 to 3.0×10 14 m -2 , and the average dislocation density of the soft phase may be 0.50×10 14 m -2 to 2.00×10 14 m -2 .

[0104] The average dislocation density can be calculated by using kernel average misorientation (KAM) data after EBSD measurement of a cross-section parallel to the rolling direction at the 1 / 4 position along the thickness direction of the steel plate, and can be calculated by the following formula. For convenience, the above calculation can utilize software such as OIM analysis™ (EDAX) for analyzing the EBSD measurement results.

[0105] [Formula]

[0106]

[0107] (In the formula, θ is the average misorientation (KAM value), u is the unit length (step size in the EBSD measurement), and b is the burgers vector.)

[0108] When the average dislocation density of the hard phase is less than 2.0×10 14 m -2 , the strength may be significantly reduced. When this value exceeds 3.0×10 14 m -2 , the ductility decreases, and the quality of the shear surface may be poor.

[0109] When the average dislocation density of the soft phase is less than 0.50×10 14 m -2 , the strength may not reach the level required by the present invention, and the generation of burrs may become serious during shearing forming. On the other hand, when this value exceeds 2.00×1014 m -2 When m, the yield strength increases, and there may be a problem of reduced elongation.

[0110] Next, a method for manufacturing the steel plate of the present invention will be described in detail.

[0111] The steel plate according to an embodiment of the present invention can be manufactured by reheating, hot rolling, primary cooling, air cooling, secondary cooling, and coiling a steel slab that satisfies the alloy composition of the present invention.

[0112] Reheating

[0113] The steel slab that satisfies the alloy composition of the present invention can be reheated in a temperature range of 1150 - 1350 °C.

[0114] When the reheating temperature is lower than 1150 °C, the precipitates cannot be fully redissolved, the formation of precipitates in the subsequent hot rolling process is reduced, coarse TiN remains, and the heat storage of the steel slab is insufficient. Therefore, it may be difficult to control the temperature of the steel plate at a certain level during hot rolling. On the other hand, when the reheating temperature exceeds 1350 °C, abnormal growth of austenite grains may lead to a decrease in strength.

[0115] Hot rolling

[0116] The reheated steel slab can be hot rolled at a finish rolling temperature of 850 - 1150 °C.

[0117] When the finish rolling temperature during hot rolling exceeds 1150 °C, the temperature of the hot rolled steel plate increases, the grain size becomes coarser, and the surface quality of the hot rolled steel plate may deteriorate. On the other hand, when the finish rolling temperature is lower than 850 °C, due to excessive recrystallization delay, developed stretched grains result in severe anisotropy, and the formability may also deteriorate.

[0118] Primary cooling

[0119] The steel plate manufactured in the hot rolling step can be subjected to primary cooling at an average cooling rate of 50 - 100 °C / second to a temperature range of 430 - 600 °C. During the primary cooling, based on the width direction of the steel plate, the surface temperature (TE) of the edge portion corresponding to 30% of the area from both end portions along the direction of the other end portion can be cooled to a temperature range of 500 - 600 °C, and the surface temperature (TC) of the central portion corresponding to 40% of the area of the center except for the two side edge portions can be cooled to a temperature range of 430 - 500 °C.

[0120] In the present invention, based on the width direction of the steel plate, regions corresponding to 30% are respectively defined from both end portions along the direction toward the other end portion or the center, that is, the regions corresponding to 60% of the entire steel plate are classified as edge portions, and the central 40% region excluding the edge portions is classified as the central portion.

[0121] In the present invention, soft phases are formed in the fine structure of the steel during primary cooling and air cooling, and hard phases can be formed during secondary cooling and coiling. However, since the untransformed phase before secondary cooling is transformed into a hard phase after secondary cooling, a uniform soft phase must be formed at each width position of the steel plate after primary cooling and air cooling. Generally, when the steel plate is cooled, heat transfer occurs rapidly at the edge portions of the steel plate, and more hard phases are formed compared to the central portion. Therefore, in the present invention, in order to impart a uniform cooling rate to each position across the width of the steel plate, it is desirable to control the cooling termination temperatures of the central portion and the edge portions to be different. As proposed in the present invention, when a uniform soft phase is formed at each width position during primary cooling and air cooling, during secondary cooling, the proportion of the formed hard phase becomes constant, and the dependence on the cooling conditions after coiling can be reduced. When soft phases and hard phases are formed uniformly, excellent material uniformity and shearing formability as desired in the present invention can be ensured.

[0122] When the cooling rate during primary cooling is less than 50 °C / second, the fraction of ferrite may be too high, and the average dislocation density is equal to or lower than the desired level, which may not be conducive to ensuring strength. On the other hand, when this cooling rate exceeds 100 °C / second, during primary cooling, the surface temperature (TC) of the central portion decreases excessively, the ferrite fraction decreases significantly, and the hard phase increases by more than the necessary amount, which may lead to insufficient elongation.

[0123] In addition, in the present invention, the difference in the cooling rate at each width position of the steel plate may be exacerbated. Therefore, in order to suppress the formation of hard phases from varying depending on the width position, the edge portion and the central portion can be distinguished and cooled to different temperature ranges.

[0124] Specifically, in the present invention, the purpose is to return the temperature of the steel plate and make the steel plate uniformly have a temperature of 550 - 650 °C during air cooling as a subsequent process. Therefore, the temperature of the central portion needs to be overcooled below the desired temperature of 550 - 650 °C, and preferably the edge portion with a high cooling rate is 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.

[0125] When the surface temperature of the edge part is lower than 500 °C, the formation of the soft phase may be insufficient. During air cooling, there may be a problem that the hot-rolled steel sheet does not reach the desired temperature range. On the other hand, when the surface temperature exceeds 600 °C, it may exceed the desired temperature range during air cooling, and there may be a problem of excessive formation of the soft phase in the final fine structure. According to an embodiment of the present invention, during primary cooling, the surface temperature of the edge part can be 510 °C or higher.

[0126] When the surface temperature of the center part is lower than 430 °C, there is a problem of the phase transformation of bainite, which is a hard phase. On the other hand, when the surface temperature exceeds 500 °C, the effect of temperature return decreases, and thus there may be a problem of not reaching the desired temperature range.

[0127] Air cooling

[0128] The steel sheet after the primary cooling can be air-cooled for 4.0 - 10.0 seconds.

[0129] For the steel sheet after the primary cooling, by stopping intentional cooling, the temperature of the steel sheet can return to the desired temperature through internal latent heat and transformation heat release. When the steel sheet after the primary cooling is air-cooled, the temperature of the steel sheet returns, and the average temperature of the steel sheet can return to the temperature range of 550 - 650 °C.

[0130] When the time during air cooling is less than 4.0 seconds, there may be no such temperature return effect. On the other hand, when the time exceeds 10.0 seconds, the fraction of ferrite in the fine structure increases significantly, and bainite and martensite, which are hard phases, may decrease. In addition, pearlite tissue and coarse carbides can be formed in the thickness center part of the area with a high temperature of the steel sheet, so the cross-section quality after shearing and forming may be poor.

[0131] Secondary cooling and coiling

[0132] The steel sheet after the air cooling can be secondarily cooled at an average cooling rate of 10 - 100 °C / second, cooled to the temperature range of 50 - 200 °C and coiled.

[0133] When the coiling temperature exceeds 200 °C, the average dislocation density of the hard phase is out of the range proposed by the present invention, and thus it may be difficult to ensure the strength. According to an embodiment of the present invention, the coiling temperature can be limited to 150 °C or lower. On the other hand, when the coiling temperature is lower than 50 °C, the formation amount of martensite is too large, the average dislocation density of the hard phase exceeds the range proposed by the present invention, the elongation of the steel becomes poor, and residual cooling water exists, so there may be a problem of steel sheet corrosion. According to an embodiment of the present invention, the coiling temperature can be limited to 70 °C or higher.

[0134] When the cooling rate exceeds 100 °C / second, there may be a problem that the average dislocation density of the hard phase becomes too high, resulting in a decrease in elongation. There is no particular limitation on the lower limit of the cooling rate. However, in order to control the cooling rate to less than 10 °C / second, there is a problem that the equipment length of the cooling zone must be increased, and there is a problem that it is difficult to make the target coiling temperature below 200 °C.

[0135] The tensile strength of the steel plate manufactured as described above is 780 MPa or more, the yield ratio is 0.70 to 0.85, and the number of microcracks with a length of 0.1 mm or more on the shear surface during stamping forming with a stamping gap of 5 - 20% is 10 pieces / cm 2 Hereinafter, the length of the maximum crack is 1 mm or less, thereby ensuring excellent strength, shear formability, and a low yield ratio. Detailed implementation mode

[0136] Hereinafter, the present invention will be described in more detail by way of examples. However, it should be noted that the following examples are only used to illustrate the present invention for more detailed description and are not used to limit the scope of the rights of the present invention.

[0137] (Example)

[0138] Under the conditions described in Table 2 below, a steel slab having the composition shown in Table 1 below was made into a steel plate. The reheating temperature not disclosed in Table 2 was 1250 °C, and the thickness of the steel plate after hot rolling was made to be 3.2 mm uniformly.

[0139] [Table 1]

[0140]

[0141] [Relationship 1]

[0142] X = ([Nb] / 93 + A / 48 + [V] / 51) / ([C] / 12 + [N] / 14)

[0143] A = [Ti] - 3.42[N] - 1.5[S]

[0144] (In the formula, [Nb], [V], [C], [N], [Ti], and [S] are the weight percentages of each element.)

[0145] [Relationship 2]

[0146] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B]

[0147] (In the formula, [Mn], [Mo], [Cr], and [B] are the weight percentages of each element.)

[0148] [Table 2]

[0149]

[0150] The microstructure and mechanical properties of the manufactured steel plates are shown in Tables 3 and 4 below. The microstructure is observed in a cross-section perpendicular to the rolling direction of the steel plate, and analyzed at the 1 / 4 - 1 / 2 position based on the thickness direction. The discrimination and measurement of the area fractions of ferrite, ferrite-based low-temperature transformation products, bainite, martensite, and pearlite formed in the steel are analyzed at magnifications of 3000 - 5000 using electron backscatter diffraction (EBSD, (JEOL JSM-1001F)). In addition, the average dislocation density (GND) is based on a cross-section parallel to the rolling direction at the 1 / 4 position along the thickness direction of the steel plate, and is measured using OIManalysis TM (EDAX) after EBSD measurement.

[0151] As mechanical properties, the yield strength, tensile strength, elongation at break, and yield ratio are measured and shown. Among them, the 0.2% offset yield strength (YS), tensile strength (TS), and elongation at fracture (T-El) are the results of testing by collecting JIS No. 5 standard test pieces in the direction perpendicular to the rolling direction. In addition, as physical properties, the number of crack lengths in each stamping cross-section is measured and shown. This is evaluated by stamping holes with a diameter of 10 mm. At this time, after stamping with different stamping clearances of 5%, 10%, and 20%, it is the average result of the number of microcracks observed in the cross-section parallel to the rolling direction and the cross-section perpendicular to the rolling direction, and each result value is the number of occurrences of each crack length.

[0152] [Table 3]

[0153]

[0154] F: Equiaxed ferrite, BF: Ferrite-based low-temperature transformation phase, B: Bainite, M: Martensite, P: Pearlite

[0155] [Table 4]

[0156]

[0157] As shown in Tables 3 and 4, in the case of the inventive examples that satisfy the alloy composition and manufacturing conditions of the present invention, the microstructure characteristics proposed by the present invention are satisfied, and the physical properties desired by the present invention can also be ensured.

[0158] Figure 1Figures (a) and (b) show the relationship diagrams of the occurrence times of the shear plane crack sizes for the inventive example and the comparative example according to a stamping clearance of 10% and a stamping clearance of 20%, respectively. Specifically, for stamping clearances of 10% and 20%, it can be confirmed that the number of crack occurrences in the comparative example is more than that in the inventive example. In the inventive example, there are no cracks with a size exceeding 1.0 mm, and the number of cracks with a size of 1.0 mm or less is significantly reduced compared to the comparative example.

[0159] On the other hand, Comparative Example 1 is an example that satisfies the content range of the alloying elements proposed in the present invention but does not satisfy the conditions of Formula 1. As a result, the average dislocation density of the soft phase exceeds the range proposed in the present invention, which is judged to be caused by an increase in the formation of fine precipitates in the soft phase. No large cracks with a length exceeding 1 mm were found in the stamping part, but when the clearance was 20%, the occurrence of cracks with a size of 0.1 - 1.0 mm increased significantly. In addition, due to the high yield ratio, the yield strength increases excessively due to work hardening during forming, so there is also a problem of poor formability.

[0160] Comparative Example 2, Comparative Example 3, Comparative Example 12, and Comparative Example 13 are examples that do not satisfy Formula 2. In Comparative Example 2 and Comparative Example 12, excessive alloying components with high hardenability effects are included, ensuring strength stably, but the elongation is insufficient. Therefore, the quality of the shear plane is also poor. In Comparative Example 3 and Comparative Example 13, the alloying components with excellent hardenability effects are insufficient, and the hard phase at the level proposed in the present invention is not formed. As a result, the target strength cannot be ensured. In addition, as the clearance increases, the occurrence of cracks becomes more serious, and cracks with a length exceeding 1 mm are confirmed.

[0161] Comparative Example 4 and Comparative Example 5 are cases where the cooling termination temperature during the first cooling after hot rolling deviates from the range proposed in the present invention. In Comparative Example 4, during the first cooling, the termination temperature exceeds the upper limit standard, the formation of the hard phase is insufficient, and unnecessary pearlite is also formed. Therefore, the cross-sectional quality after stamping is poor. In Comparative Example 5, during the first cooling, the termination temperature deviates from the lower limit standard, the fraction of the soft phase is insufficient, and an excessive amount of the hard phase is formed. Therefore, the cross-sectional quality after stamping is poor.

[0162] Comparative Example 6 and Comparative Example 7 are cases where the air cooling time after the first cooling deviates from the range of the present invention. The exposure time in the high-temperature region is long. Due to the latent heat inside the steel plate and the heat generated by the phase transformation, the fraction of the soft phase increases significantly, and pearlite is also formed. The cross-sectional quality after stamping is poor, and the yield ratio also exceeds the range proposed in the present invention. In particular, in the case of Comparative Example 7, it can be confirmed that the average temperature of the steel plate after air cooling exceeds the temperature range proposed in the present invention.

[0163] Comparative Example 8 is an example where the air-cooling time after the first cooling does not reach the range of the present invention. Before the temperature of the returned steel sheet, secondary cooling is performed, and the fraction of the soft phase does not reach the desired level, while the fraction of the hard phase exceeds the desired range. As a result, the cross-sectional quality after stamping is poor.

[0164] Comparative Example 9 is a case where the cooling rate during secondary cooling is too fast. Due to excessive cooling, the coiling temperature range desired by the present invention cannot be satisfied. Therefore, the dislocation density of the hard phase exceeds the proposed range, and the cross-sectional quality after stamping is poor. It is judged that the main reason is the increase in the physical property difference between the soft phase and the hard phase.

[0165] Comparative Example 10 and Comparative Example 11 are cases where the cooling termination temperature during secondary cooling deviates from the range proposed by the present invention. Comparative Example 10 is a case where the coiling temperature after secondary cooling does not reach the proposed temperature range, and the dislocation density of the hard phase is too high. Therefore, the cross-sectional quality after stamping is poor. Comparative Example 11 is a case where the coiling temperature after secondary cooling exceeds the proposed temperature range, and the dislocation density of the hard phase does not reach the proposed level. As a result, the yield ratio is too high, and the cross-sectional quality is also poor.

[0166] The present invention has been described in detail through the above embodiments, but other forms of embodiments can also be implemented. Therefore, the technical idea and scope of the claims are not limited to the embodiments.

Claims

1. A steel plate, by weight %, the steel plate comprises: 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%, the balance Fe and inevitable impurities, the X value defined in the following relationship 1 is 0.010 to 0.200, the T value defined in the following relationship 2 is 1.500 to 4.200, by area %, the microstructure comprises 30 - 70% of a hard phase containing bainite and martensite, 30 - 70% of a soft phase containing ferrite, and 3% or less of pearlite, The average dislocation density of the hard phase is 2.0×10 14 m -2 to 3.0×10 14 m -2 , and the average dislocation density of the soft phase is 0.50×10 14 m -2 to 2.00×10 14 m -2 , [Relationship 1] X = ([Nb] / 93 + A / 48 + [V] / 51) / ([C] / 12 + [N] / 14) A = [Ti] - 3.42[N] - 1.5[S] In the formula, [Nb], [V], [C], [N], [Ti] and [S] are the weight % of each element, [Relationship 2] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] In the formula, [Mn], [Mo], [Cr] and [B] are the weight % of each element.

2. The steel plate according to claim 1, wherein, by weight %, the steel plate further comprises one or more selected from 0.005 - 0.030% of Nb, 0.005 - 0.120% of Ti, 0.005 - 0.200% of V, 0.0003 - 0.0030% of B.

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

85.

4. The steel plate according to claim 1, wherein, When the steel plate is stamped and formed with a stamping clearance of 5-20%, the number of fine cracks with a length of 0.1 mm or more on the shear surface is 10 / cm. 2 Hereinafter, the length of the maximum crack is 1 mm or less.

5. A method for manufacturing a steel plate, which comprises the following steps: Reheat the steel billet. By weight %, the steel billet comprises: 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%, the balance Fe and inevitable impurities, the X value defined in the following relationship 1 is 0.010 to 0.200, and the T value defined in the following relationship 2 is 1.500 to 4.200; Hot-roll the reheated steel billet; Cool the steel plate manufactured in the hot-rolling step at an average cooling rate of 50 - 100 °C / second for the first cooling to a temperature range of 430 - 600 °C; Air-cool the steel plate after the first cooling for 4.0 - 10.0 seconds; And Cool the steel plate after the air-cooling at an average cooling rate of 10 - 100 °C / second for the second cooling to a temperature range of 50 - 200 °C and coil it up, During the primary cooling, with the width direction of the steel plate as the reference, the surface temperature (TE) of the edge parts of the regions corresponding to 30% from both ends along the direction of the other end is cooled to the temperature range of 500 - 600 °C, and the surface temperature (TC) of the central part of the region corresponding to 40% of the center of the region other than the two edge parts is cooled to the temperature range of 430 - 500 °C. [Relationship 1] X = ([Nb] / 93 + A / 48 + [V] / 51) / ([C] / 12 + [N] / 14) A = [Ti] - 3.42[N] - 1.5[S] In the formula, [Nb], [V], [C], [N], [Ti] and [S] are the weight percentages of each element. [Relationship 2] T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] In the formula, [Mn], [Mo], [Cr] and [B] are the weight percentages of each element.

6. The manufacturing method of the steel plate according to claim 5, wherein, By weight percentage, the steel billet further contains one or more selected from 0.005 - 0.030% of Nb, 0.005 - 0.120% of Ti, 0.005 - 0.200% of V, and 0.0003 - 0.0030% of B.

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

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

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