Hot-rolled steel sheet, and method for manufacturing same

Abstract

Provided are a hot-rolled steel sheet and a method for manufacturing same. The present invention provides a hot-rolled steel sheet and a method for manufacturing same, the hot-rolled steel sheet satisfying a core section hardness of 40-50 HrC and having excellent bendability (the bending radius at which surface ridges do not form during 90° bending / steel sheet thickness ratio is low).

Description

Hot-rolled steel sheet and its manufacturing method

[0001] The present invention relates to a high-strength hot-rolled steel sheet having high hardness and bending workability, which is mainly used in a frame to protect the battery of an eco-friendly vehicle, and a method for manufacturing the same.

[0002] Conventional hot-rolled high-strength hot-rolled steel sheets primarily utilize a martensitic microstructure, resulting in high strength and hardness. However, their high strength leads to poor bending workability, limiting their processing. Consequently, they are typically used through roll forming or minimal machining. This limitation also applies to wear-resistant steels utilizing martensite as the primary phase, and various technologies have been proposed to overcome this limitation.

[0003] In patent document 1, the bendability was increased by limiting the alloy components C, Si, Mn, etc. of the steel and changing the microstructure composition according to the thickness position.

[0004] In addition, in patent document 2, a high-strength steel plate was designed to increase bendability by configuring tempered martensite as the main phase in the center of the plate thickness and ferrite and pearlite as the main phase in the soft surface region.

[0005] However, Patent Document 1 defined the surface area as up to 100㎛ of the steel plate thickness, and did not consider the microstructure of the extreme surface layer (10㎛) where deformation is concentrated, and only introduced a soft structure, and only achieved a degree of bendability that can be easily obtained in a region with relatively low strength.

[0006] In addition, Patent Document 2 has a disadvantage in that the surface soft part is composed of ferrite and pearlite, which are structures that are excessively soft compared to the core part, which is the center of thickness, and thus has low hardness, and when bending is performed, deformation is concentrated in the surface soft part, which causes rapid fracture.

[0007] In addition, alloying elements such as Si, Mn, Mo, Cr, Cu, and Ni, which are mainly used to manufacture the above high-strength steels, are effective in improving the hardness and formability of the hot-rolled steel sheets, but if a large amount of alloying elements is added to improve these properties, segregation of the alloying elements and unevenness of the microstructure occur, resulting in poor bending workability. In particular, steels with high hardenability have a problem in that their microstructure changes sensitively when cooled, so that the low-temperature transformation structure phase is formed unevenly, making it difficult to obtain even higher bending workability.

[0008] [Prior Art Literature]

[0009] [Patent Document]

[0010] (Patent Document 1) Japanese Patent Publication No. 2013-117068

[0011] (Patent Document 2) Korean Patent Publication No. 2021-0088646

[0012] The purpose of the present invention is to provide a hot-rolled steel sheet having a deep cross-sectional hardness of 40 to 50 HrC and excellent bending workability (small bending radius / steel sheet thickness ratio that does not cause surface unevenness when bent at 90°), and a method for manufacturing the same.

[0013] Therefore, one aspect of the present invention is,

[0014] In weight %, C: 0.17~0.26%, Si: 0.01~0.5%, Mn: 0.3~2.0%, Cr: 0.005~0.5%, Mo: 0.005~0.3%, Nb: 0.001~0.01%, Ti: 0.005~0.08%, V: 0.005~0.2%, Al: 0.01~0.5%, P: 0.003~0.05%, S: 0.001~0.01%, N: 0.001~0.01%, B: 0.0005~0.005%, the remainder being Fe and other unavoidable impurities, and satisfying the following relational expression 1,

[0015] A hot-rolled steel sheet having a microstructure of a steel sheet, wherein a surface layer from the surface of the steel sheet to a depth of 10 ㎛ includes, in terms of area %, ferrite: 15 to 50%, twinned martensite: 2 to 15%, residual bainite, martensite, and auto-tempered martensite, and a core layer at a point t / 4 of the steel sheet includes, in terms of area %, martensite + auto-tempered martensite: 90% or more, and at least one of ferrite and bainite: 10% or less, and wherein the core layer is a hot-rolled steel sheet having an aspect ratio of old austenite of 10 or less.

[0016] [Relationship 1]

[0017] 30 ≤ T ≤ 80

[0018] T =(([C] / 10) 0.5 *1.07)*(0.7*[Si]+1)*(5*[Mn]+1)*(2.16*[Cr]+1)*(3*[Mo]+1)*25.4

[0019] *In the above equation 1, C, Si, Mn, Cr, and Mo represent the weight percent of the corresponding alloy elements. If not added, 0 is substituted.

[0020] Another aspect of the present invention is

[0021] A step of heating a slab containing, in wt%, C: 0.17 to 0.26%, Si: 0.01 to 0.5%, Mn: 0.3 to 2.0%, Cr: 0.005 to 0.5%, Mo: 0.005 to 0.3%, Nb: 0.001 to 0.01%, Ti: 0.005 to 0.08%, V: 0.005 to 0.2%, Al: 0.01 to 0.5%, P: 0.003 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, B: 0.0005 to 0.005%, the remainder Fe and other unavoidable impurities, and satisfying the following relational expression 1, to 1150 to 1350°C;

[0022] A step of obtaining a bar by completing rough rolling of the above-mentioned heated slab at a rough rolling temperature (RDT) of 900 to 1100℃ based on 1 / 2t (t: thickness of steel);

[0023] A step of performing surface cooling while descaling the entire width of the bar three or more times;

[0024] A step of obtaining a hot rolled steel sheet by completing the finishing rolling at a finishing rolling temperature (FDT) of 800℃ to ST+50℃ that satisfies the following relational expression 2 based on 1 / 2t (t: thickness of steel) of the cooled bar above;

[0025] A step of first cooling the above hot-rolled steel sheet to a first cooling stop temperature of Ms to Ms+50°C at a first average cooling rate of 60 to 90°C / sec: and

[0026] The present invention relates to a method for manufacturing a hot-rolled steel sheet, which includes a step of secondarily cooling the first-cooled hot-rolled steel sheet to a coiling temperature (CT) of 50°C to Ms-50°C at a second average cooling rate of 1 to 50°C / sec and then coiling the second-cooled hot-rolled steel sheet.

[0027] [Relationship 1]

[0028] 30 ≤ T ≤ 80

[0029] T =(([C] / 10) 0.5 *1.07)*(0.7*[Si]+1)*(5*[Mn]+1)*(2.16*[Cr]+1)*(3*[Mo]+1)*25.4

[0030] *In the above equation 1, C, Si, Mn, Cr, and Mo represent the weight percent of the corresponding alloy elements. If not added, 0 is substituted.

[0031] [Relationship 2]

[0032] 800℃~ST+50℃

[0033] ST = 800 + 140 [C] - 145 [Si] + 80 [Mn] - 30 [Cr] + 10 [Mo]

[0034] *C, Si, Mn, Cr, and Mo in the above equation 2 are the weight percent of the corresponding alloy elements. If not added, 0 is substituted.

[0035] The present invention, having a configuration as described above, can effectively provide a hot-rolled high-strength steel sheet that satisfies bendability (R / t) < 4 by controlling the difference in hardness between the steel sheet surface and the t / 4 deep cross-section and the aspect ratio of the deep microstructure and the composition of the surface microstructure, thereby having excellent bending workability (small bending radius / steel sheet thickness ratio that does not cause surface unevenness when bent at 90°).

[0036] Figure 1 is a graph showing the relationship between the bendability (R / t) according to the t / 4 deep cross-sectional hardness of the invention example and the comparative example in the embodiment of the present invention.

[0037] Figure 2 is a photograph of twinned martensite formed on the surface of Invention Example 1, taken using a transmission electron microscope.

[0038] Figure 3 is a photograph of ferrite formed on the surface of Invention Example 1, taken using a transmission electron microscope.

[0039] Figure 4 is a photograph of the surface cross-section microstructure distribution of Invention Example 1 and Comparative Example 3 taken with an electron microscope.

[0040] Figure 5 is a photograph showing carbides in the deep microstructure (auto-tempered martensite) of Invention Example 1.

[0041] Figure 6 is a photograph taken with an optical microscope of a surface with irregularities (bending radius of 9 mm) and a good surface without irregularities (bending radius of 7 mm) during a bending test of Invention Example 1.

[0042] Hereinafter, the present invention will be described.

[0043] To address the aforementioned problems of the prior art, the inventors of the present invention investigated the changes in the core cross-sectional hardness and bending workability of various steels with different compositions and microstructures according to the characteristics of the composition, manufacturing process, and microstructure. As a result, they confirmed that the detailed characteristics and hardness differences of each composition of the surface and core layers were correlated with the core cross-sectional hardness and bending workability, and derived the following equations 1 and 2 from them. And based on the derived relationships, the present invention proposes that when the microstructure of the steel is composed of a martensite phase as the main phase with an aspect ratio (major axis / minor axis) in the range of 1 to 10 based on the grain boundaries of the old austenite in the deep part, and the microstructure of the surface layer up to a depth of 10㎛ includes a certain fraction of ferrite and twinned martensite, the difference between the surface hardness and the deep cross-section hardness satisfies the range of 9 to 25%, and the deep hardness satisfies 40 to 50 HrC, it is possible to manufacture a hot-rolled steel sheet having excellent bending workability compared to the deep cross-section hardness.

[0044]

[0045] Hereinafter, the composition components of the hot-rolled steel sheet of the present invention and the reasons for limiting the content thereof are explained, and here, “%” means weight% unless otherwise specified.

[0046] C: 0.17~0.26%

[0047] The above-mentioned C is the most economical and effective element for strengthening steel, and has a great influence on strength and hardness. As its addition amount increases, the hardenability increases, making it easier to form hard phases such as bainite and martensite in the microstructure, thereby increasing the tensile strength. In addition, since C tends to accumulate in austenite rather than ferrite, ferrite is formed in the surface layer, and when the C accumulated in austenite transforms into martensite, a mechanism is formed that changes into twinned martensite. However, when the C content exceeds 0.26%, there is a problem that the martensite strength increases excessively, resulting in excessive strength increase, lowered bending workability, and poor weldability. On the other hand, when the C content is less than 0.17%, it is difficult to obtain a sufficient strengthening effect. Therefore, the C content is preferably in the range of 0.17 to 0.26%. More preferably, the C content is limited to the range of 0.18 to 0.24%.

[0048] Si: 0.01~0.5%

[0049] The above-mentioned Si is an element that is advantageous in improving formability by deoxidizing molten steel, exerting a solid solution strengthening effect, and delaying the formation of coarse carbides. If the content of the Si is less than 0.01%, the above-mentioned effect cannot be sufficiently obtained. On the other hand, if the content of the Si exceeds 0.5%, a red scale due to Si is formed on the surface of the steel sheet during hot rolling, which not only greatly deteriorates the surface quality of the steel sheet, but also causes problems in that ductility and weldability are reduced. Therefore, the content of the Si is preferably in the range of 0.01 to 0.5%. More preferably, the content of the Si is limited to the range of 0.05 to 0.4%.

[0050] Mn: 0.3~2.0%

[0051] The above manganese, like silicon, is an effective element for strengthening steel through solid solution, and increases the hardenability of steel, facilitating the formation of hard phases such as bainite and martensite during cooling after hot rolling. In addition, when ferrite is formed in the surface layer, it is concentrated in austenite, and when it subsequently transforms into martensite, it forms a mechanism that changes into twinned martensite. If the manganese content is less than 0.3%, the above-mentioned effect cannot be sufficiently obtained. On the other hand, if the manganese content exceeds 2.0%, the grain boundaries become weak, causing problems such as low-temperature cracking. In addition, there are problems such as excessive strength increase and reduced formability, and when slab casting is performed during the continuous casting process, a segregation zone is significantly developed in the center of the thickness, and when cooling after hot rolling, the microstructure is formed unevenly in the thickness direction, resulting in inferior bending workability. In particular, it makes it difficult to manufacture a uniform microstructure during cooling across the full length and full width of the hot-rolled steel sheet. Therefore, it is preferable that the content of Mn be in the range of 0.3 to 2.0%. More preferably, the content of Mn is limited to the range of 0.5 to 1.8%.

[0052] Sol.Al: 0.01~0.5%

[0053] The above Al is an element added mainly for deoxidation. When the content of the Al is less than 0.01%, the above-described effect cannot be sufficiently obtained. On the other hand, when the content of the Al exceeds 0.5%, it combines with nitrogen to form excessive AlN, which easily causes corner cracks to occur in the slab during continuous casting and easily causes defects due to the formation of inclusions. Therefore, the content of the Al is preferably in the range of 0.01 to 0.5%. The lower limit of the Al content is more preferably 0.015%, and more preferably 0.02%. The upper limit of the Al content is more preferably 0.1%, and more preferably 0.08%, and most preferably 0.05%.

[0054] Cr: 0.005~0.5%

[0055] The above Cr strengthens the steel and delays the ferrite phase transformation during cooling, thereby helping the formation of martensite and bainite. If the Cr content is less than 0.005%, the above-described effect cannot be sufficiently obtained. On the other hand, if the Cr content exceeds 0.5%, similar to Mn, the segregation zone in the center of the thickness is greatly developed, and the microstructure in the thickness direction becomes non-uniform, thereby reducing the bending workability. In addition, due to excessively high hardenability, it is impossible to form a ferrite structure on the surface. Therefore, the Cr content is preferably in the range of 0.005 to 0.5%. More preferably, the Cr content is limited to the range of 0.08 to 0.4%.

[0056] Mo: 0.005~0.3%

[0057] Mo increases the hardenability of steel, facilitating the formation of martensite and bainite. If the content of Mo is less than 0.005%, the above-described effect cannot be sufficiently obtained. On the other hand, if the content of Mo exceeds 0.3%, the bending workability is rapidly deteriorated due to excessive increase in hardenability, which is economically disadvantageous and detrimental to weldability. In addition, the excessively high hardenability makes it impossible to form a ferrite structure on the surface. Therefore, the content of Mo is preferably in the range of 0.005 to 0.3%. More preferably, the content of Mo is limited to the range of 0.01 to 0.25%.

[0058] P: 0.003~0.05%

[0059] The above P, like Si, has both the effect of strengthening the solid solution and promoting ferrite transformation. However, controlling the P content to less than 0.003% requires a lot of manufacturing cost, which is economically disadvantageous and insufficient for obtaining strength. On the other hand, if the P content exceeds 0.05%, embrittlement due to grain boundary segregation may occur, microcracks are likely to occur during bending, and ductility and impact resistance properties are significantly reduced. Therefore, the P content is preferably in the range of 0.003 to 0.05%. The lower limit of the P content is more preferably 0.005%, more preferably 0.007%, and most preferably 0.01%. The upper limit of the P content is more preferably 0.03%.

[0060] S: 0.001~0.01%

[0061] The above S is an impurity present in steel, and if its content exceeds 0.01%, it combines with Mn, etc. to form non-metallic inclusions, which easily causes fine cracks to occur during bending of the steel and significantly reduces impact resistance. In the present invention, there is no particular limitation on the lower limit of the above S content, but since controlling it to less than 0.001% takes a lot of time during steelmaking, which reduces productivity, the lower limit of the above S content may be limited to 0.001% in consideration of this. Therefore, the above S content is preferably in the range of 0.001 to 0.01%.

[0062] N: 0.001~0.01%

[0063] The above N, together with C, is a representative solid solution strengthening element, and forms coarse precipitates together with Ti, Al, etc. When the content of the above N is less than 0.001%, it is difficult to sufficiently obtain the above-mentioned effect, and in order to control the content of the above N to less than 0.001%, a lot of time is required during the steelmaking process, which reduces productivity. Meanwhile, although the solid solution strengthening effect of the above N is generally superior to that of carbon, there is a problem that the toughness is greatly reduced when the content of the above N exceeds 0.01%. Therefore, the content of the above N is preferably in the range of 0.001 to 0.01%.

[0064] Ti: 0.005~0.08%

[0065] The above Ti, along with Nb and V, is a representative precipitation strengthening element, and forms coarse TiN in steel due to its strong affinity with N. TiN has the effect of inhibiting grain growth during the heating process for hot rolling. In addition, the remaining Ti after reacting with nitrogen becomes a solid solution in the steel and combines with carbon to form TiC precipitates, which is a useful component for improving the strength of the steel. However, if the Ti content is less than 0.005%, the above effect cannot be obtained, and if the Ti content exceeds 0.08%, there is a problem that the bending workability is inferior during forming due to the generation of coarse TiN and the coarsening of the precipitates. Therefore, in the present invention, it is preferable to limit the Ti content to 0.005 to 0.08%. More preferably, the Ti content is limited to the range of 0.01 to 0.06%.

[0066] Nb: 0.001~0.01%

[0067] The above-mentioned Nb, together with Ti and V, is a representative precipitation strengthening element, and is effective in improving the strength and impact toughness of steel due to the grain refinement effect by precipitating during hot rolling and delaying recrystallization. However, if the Nb content is less than 0.001%, the above-mentioned effect cannot be obtained, and if the Nb content exceeds 0.01%, there is a problem of inferior bending workability due to the formation of coarse composite precipitates during hot rolling. In addition, if it exceeds 0.01%, it has the effect of making the aspect ratio of the parent austenite before martensite phase transformation excessively large, so that the low aspect ratio aimed for in the present invention cannot be achieved. Therefore, in the present invention, it is preferable to limit the Nb content to 0.001 to 0.01%. More preferably, the Nb content is limited to the range of 0.005 to 0.01%.

[0068] V: 0.005~0.05%

[0069] The above-mentioned V, along with Nb and Ti, is a representative precipitation strengthening element, and it hardly precipitates during hot rolling and forms precipitates after high-temperature coiling, thereby improving the strength of steel. Therefore, it is effective in additionally improving strength without increasing deformation resistance and rolling load due to delayed recrystallization during hot rolling. In order to obtain this effect in the present invention, it is preferable that V is contained in an amount of 0.005% or more. However, if the content is excessive, there is a problem of inferior bending workability due to the formation of coarse precipitates, and like Nb, it is difficult to maintain a low aspect ratio of the matrix austenite, and it is also economically disadvantageous. Therefore, in the present invention, it is preferable to limit the upper limit to 0.05%, and more preferably to 0.03%.

[0070] B: 0.0005~0.005%

[0071] The above B, when present in a solid solution state in steel, mainly segregates at grain boundaries and stabilizes grain boundaries, thereby improving the brittleness of steel. Furthermore, it plays a role in suppressing the formation of coarse AlN nitrides by stabilizing solid solution N. In addition, it is effective in delaying ferrite phase transformation and forming hard phases such as bainite and martensite. In the present invention, the above effect is achieved only when the content is 0.0005% or more, and when the content exceeds 0.005%, the effect no longer increases, and there is a problem that ductility decreases, resulting in poor formability. Therefore, in the present invention, it is preferable to limit the upper limit to 0.005%.

[0072] In addition to the above composition, the remainder is Fe. However, during the normal manufacturing process, unavoidable impurities from raw materials or the surrounding environment may be mixed in, and thus, this cannot be ruled out. Since these impurities are readily apparent to anyone skilled in the art, not all details are specifically mentioned in this specification. Furthermore, the addition of other effective ingredients beyond the above composition is not excluded.

[0073] Relationship 1

[0074] Meanwhile, in the present invention, in the steel having the component range as described above, it is preferable to control the value of the following relational expression 1 composed of C, Si, Mn, Cr, and Mo to 30 or more and 80 or less. The following relational expression 1 factors out a combination of alloy elements that can maintain the formation of hard phases, bainite and lath martensite, among the steel microstructure of the present invention at an appropriate level. That is, as the "T" value of the following relational expression 2 increases, the formation of hard phases, bainite and lath martensite, increases, and the hardness value of each hard phase also increases. Therefore, if the value defined by the following relational expression 1 is less than 10, the steel plate hardness value decreases, and therefore, the larger this value is, the more advantageous it is for securing strength and hardness, but if it exceeds 80, ferrite and twinned martensite cannot be formed in the surface layer, which deteriorates the bending workability of the steel, and there is a problem that the material deviation increases in the overall length and overall width of the hot-rolled steel plate.

[0075] [Relationship 1]

[0076] 30 ≤ T ≤ 80

[0077] T =(([C] / 10) 0.5 *1.07)*(0.7*[Si]+1)*(5*[Mn]+1)*(2.16*[Cr]+1)*(3*[Mo]+1)*25.4

[0078] *In the above equation 1, C, Si, Mn, Cr, and Mo represent the weight percentage of the corresponding alloy elements. If not added, 0 is substituted.

[0079] Hereinafter, the microstructure of the high-hardness hot-rolled steel sheet of the present invention will be described.

[0080] The steel plate proposed in the present invention is characterized in that the microstructure of the t / 4 deep layer includes, in terms of area%, the sum of martensite (lath, twinned) and auto-tempered martensite: 90% or more, and at least one of bainite and ferrite: 10% or less, and the microstructure of the surface layer up to a depth of 10㎛ includes, in terms of area%, ferrite: 15 to 50%, twinned martensite: 2 to 15%, and residual bainite, martensite, and auto-tempered martensite.

[0081] Twinned martensite is broadly classified as martensite, but it is distinguished from lath martensite, which has a plate-like substructure of several ㎛ in width, and auto-tempered martensite, which contains epsilon carbide as a micro-carbide, by having a twin substructure of a fine thickness (tens of nm) that is crystallographically twin-like.

[0082] And the steel plate of the present invention can satisfy an aspect ratio calculated by the major axis / minor axis of the old austenite grain size of the deep martensite in the range of 1 to 10.

[0083] When the steel plate microstructure satisfies the above-described range, the target deep cross-sectional hardness is 40-50 HrC, and the difference between the surface cross-sectional hardness and the deep cross-sectional hardness can be satisfied as 9-25% of the deep cross-sectional hardness. Therefore, the bendability is excellent compared to the deep cross-sectional hardness, and R / t 4 or less can be satisfied.

[0084] However, if the ferrite content of the surface layer is too low, the surface layer, which is mostly composed of a hard phase, has a disadvantage in that it has insufficient deformation resistance during bending, easily forming surface irregularities and ultimately leading to fracture. On the other hand, if the ferrite content is too high, the deformation that should be distributed throughout the entire thickness is concentrated only in a limited surface layer, which can form surface irregularities despite a large bending radius.

[0085] Meanwhile, the inventors of the present invention confirmed that among the microstructural characteristics of steel, not only the fraction of surface ferrite but also the fraction of twinned martensite formed by the influence of carbon and hardenable elements concentrated in austenite that has not yet been transformed while forming ferrite are important influencing factors.

[0086] The presence of twinned martensite is essential for preventing excessive strain concentration and ensuring adequate workability. Because the twinned martensite is difficult to distinguish in FE-SEM, it can be analyzed using a high-resolution transmission electron microscope, as shown in Fig. 2. Meanwhile, TEM measurements were performed on specimens taken from a cross-section parallel to the rolling direction at a position 10 μm from the surface of the hot-rolled steel sheet.

[0087] In addition, the microstructure of the above-mentioned deep part is mainly composed of martensite phase, and carbides can be observed like tempered martensite due to the auto (self) tempering effect depending on the coiling temperature control. In addition, when observing with SEM, the old austenite grain size that is long and elongated in the rolling direction can be observed, and it was confirmed that the effect of improving the bending workability was exerted when the aspect ratio was in the range of 1 to 10. This is because the bending line during bending is parallel to the rolling direction, so the deformation is in the form of elongation perpendicular to the rolling direction, and at this time, the microstructure formed long in the rolling direction has a very small grain size when viewed perpendicular to the rolling direction, which can be judged to be due to the low deformation resistance. When all of these microstructural characteristics are combined, excellent bending workability can be exhibited.

[0088] Next, the method for manufacturing a high-strength hot-rolled steel sheet of the present invention will be described in detail.

[0089] The method for manufacturing a hot-rolled steel sheet of the present invention comprises the steps of: heating a slab having the above-described composition and satisfying the above-described relational expression 1 to 1150 to 1350°C; completing rough rolling of the heated slab at a rough rolling temperature (RDT) of 900 to 1100°C based on 1 / 2t (t: thickness of steel material) to obtain a bar; performing surface cooling while descaling the entire width of the bar three or more times; completing finish rolling of the cooled bar at a finish rolling temperature (FDT) of 800°C to ST+50°C satisfying the following relational expression 2 based on 1 / 2t (t: thickness of steel material) to obtain a hot-rolled steel sheet; The method includes a step of first cooling the hot-rolled steel sheet to a first cooling stop temperature of Ms to Ms+50°C at a first average cooling rate of 60 to 90°C / sec; and a step of second cooling the first-cooled hot-rolled steel sheet to a coiling temperature (CT) of 50°C to Ms-50°C at a second average cooling rate of 1 to 50°C / sec, and then coiling the hot-rolled steel sheet.

[0090] First, in the present invention, a steel slab having the above-described composition and satisfying the above-described relationship 1 is reheated to 1150 to 1350°C. At this time, if the reheating temperature is lower than 1150°C, the precipitates are not sufficiently re-dissolved, which reduces the formation of precipitates in the process after hot rolling. In addition, coarse TiN remains, and the slab is not sufficiently aged, making it difficult to control the temperature of the steel sheet at a constant level during hot rolling. On the other hand, if it exceeds 1350°C, the strength is reduced due to abnormal grain growth of austenite crystals, so it is preferable to limit the reheating temperature to 1150 to 1350°C.

[0091] Next, in the present invention, the heated slab is subjected to rough rolling at a rough rolling temperature (RDT) of 900 to 1100°C based on 1 / 2t (t: thickness of steel) to obtain a bar.

[0092] In the present invention, surface cooling is performed while descaling the entire width of the bar three or more times, and then the finish rolling is terminated at the finish rolling temperature (FDT) defined by the following equation 2 based on 1 / 2t (t: thickness of steel).

[0093] [Relationship 2]

[0094] 800℃~ST+50℃

[0095] ST = 800 + 140 [C] - 145 [Si] + 80 [Mn] - 30 [Cr] + 10 [Mo]

[0096] *In the above equation 2, C, Si, Mn, Cr, and Mo represent the weight percentage of the corresponding alloy elements. If not added, 0 is substituted.

[0097] At this time, the surface temperature of the steel is controlled lower than the core temperature while performing descaling using a water spray device. This is to maintain a low surface temperature during the cooling stage after rolling, thereby forming some ferrite, and thereby inducing the carbon released from the ferrite to accumulate in the untransformed austenite, which then transforms into twinned martensite during the steel cooling process.

[0098] In order to improve the bending workability in relation to the hardness, the difference in the microstructure composition and hardness between the surface and core is essential, and the aspect ratio calculated by the major axis / minor axis of the old austenite grain size of the core martensite must be controlled within a predetermined range. In order to satisfy all of these conditions, the core temperature FDT at the end of the finish rolling has a temperature limit according to the above calculation formula 2 composed of the steel components. Specifically, if the finish rolling temperature is higher than the relation 2, the microstructure formation of the surface and core cannot be induced to the desired composition, and the steel becomes a typical martensitic high-strength steel, resulting in poor bending workability. On the other hand, if the finish rolling temperature is lower than the relation 2, the aspect ratio of the core structure cannot be achieved as desired, which also adversely affects the bending workability.

[0099] And in the present invention, the above-mentioned hot-rolled steel sheet is first cooled at an average cooling rate of 50 to 100°C / sec to a temperature range of Ms+50°C to Ms. Here, Ms can be defined by the following equation 3.

[0100] [Relationship 3]

[0101] Ms=430-380*[C]-13.5*[Si]-47.5*[Mn]-16*[Cr]-24[Mo]

[0102] *In the above equation 3, C, Si, Mn, Cr, and Mo are the weight% of the corresponding alloy elements.

[0103] The above cooling temperature is a condition for forming sufficient deep martensite and surface ferrite + twin martensite in a balanced manner. If the primary cooling is stopped at a temperature higher than Ms+50℃ and switched to slow cooling, the cooling will be insufficient, so that sufficient deep martensite will not be formed, resulting in insufficient deep cross-sectional hardness and thus inability to obtain the desired high-strength steel sheet. On the other hand, if rapid cooling is performed to a temperature lower than the Ms temperature, the combination of surface ferrite + twin martensite will not be obtained, resulting in inability to obtain bendability.

[0104] Not only in cooling temperature but also in cooling speed, an excessive cooling speed causes the superficial tissue to be formed differently from the intended one, and an insufficient cooling speed causes the deep tissue to be formed differently from the intended one.

[0105] In the present invention, it is preferable to control the primary cooling rate in the range of 60 to 90°C / sec.

[0106] Finally, in the present invention, the first-cooled hot-rolled steel sheet is secondarily cooled to a coiling temperature (CT) of 50°C to Ms-50°C at a second average cooling rate of 1 to 50°C / sec, and then coiled.

[0107] A lower cooling rate compared to the primary cooling described above is recommended for productivity reasons, as further rapid cooling once martensite has formed in the core degrades the plate's dimensional quality. Excessive rapid cooling inhibits surface ferrite formation, so it's best to employ as slow a cooling rate as possible.

[0108] Meanwhile, hot-rolled high-strength steels of giga grade or higher are usually heat-treated to overcome poor workability and are manufactured in the form of tempered martensite containing cementite. However, if heat treatment is omitted for productivity, there is a disadvantage of poor bending workability. Therefore, in the present invention, the coiling temperature is limited to Ms-50℃ to 50℃ to add an auto-tempering effect to the martensite in the core. The steel sheet of the present invention manufactured through this process can obtain auto-tempered martensite without a separate heat treatment process because fine epsilon carbide is observed.

[0109] Furthermore, the present invention may additionally include a step of pickling and oiling the steel sheet rolled after the secondary cooling.

[0110] Each steel slab having the composition shown in Table 1 below was prepared. Then, each of the prepared steel slabs was used to produce a hot-rolled steel sheet using the manufacturing conditions shown in Table 2 below. Meanwhile, in Table 2 below, RDT is the rough rolling end core temperature, FDT is the finish rolling end core temperature, FSB is the number of descaling cycles in finish rolling, CT is the coiling temperature, and ST is the temperature defined by Equation 2.

[0111] Classification Note CSiMnCrMoNbTiVAlPSNBT(hardenability) Invention Example 1 Invention Steel 0.210.051.30.20.010.0050.030.010.030.010.0020.0040.001545.09 Invention Example 20.190.20.90.30.020.0030.020.020.030.010.0020.0040.00241.03 Invention Example 30.220.031.10.010.150.0050.020.020.030.010.0020.0050.00239.63 Invention Example 40.20.051.20.30.020.007 0.040.010.030.010.0050.0060.00248.64Invention Example 50.190.020.70.20.20.0050.030.020.020.020.0030.0040.00239.17Invention Example 60.250.030.80.20.050.007 0.030.010.020.0050.0030.0030.00236.13Invention Example 70.230.021.70.020.010.0050.0250.020.030.0150.0040.0050.002542.66Invention Example 80.180.051.60.010. 250.0030.0250.010.030.0150.0040.0050.00360.72Invention Example 90.240.151.30.020.050.0020.0450.030.020.010.0040.0050.00241.86Invention Example 100.210.011 .60.010.020.0050.030.010.040.010.0020.0050.00238.65Comparative Example 1Comparative Strength 0.190.21.70.20.10.0050.020.010.020.010.0030.0040.00275.53Comparative Example 20.190. 21.70.20.10.0050.030.010.040.0150.0050.0030.00375.53Comparative Example 30.20.22.40.70.50.0050.040.010.030.0150.0050.0060.002357.72Comparative Example 40.180.0 50.90.0050.10.010.030.010.040.010.0040.0060.00127.28Comparative Example50.210.051.20.10.20.010.0250.010.040.0150.0030.0040.002555.52Comparative Example60.20.021.70.40.020.0050.020.020.040.0150.0020.0040.00173.16Comparative Example70.250.030.80.250.20.0180.020.010.040.010.0030.0030.002854.05Comparative Example80.280.021.80.50.10.0250.030.010. 030.0150.0010.0050.0025124.69Comparative Example 90.30.0610.40.40.020.040.030.030.020.0020.0050.0015120.69Comparative Example 10Inventive Steel 0.190.11.60.30.10.0050.020.010.020.010.0030.0040.00277.29.

[0112] *In Table 1, the unit of content of each component is weight%, and the remaining components are Fe and inevitable impurities. Also, in Table 1, T represents the T value defined by Equation 1.

[0113] The microstructure and properties of each hot-rolled steel sheet manufactured as described above are analyzed in detail and shown in Table 3 below. The aspect ratio is the aspect ratio calculated as the major axis / minor axis of the old austenite grain size of the deep martensite. The microstructure was observed using a transmission electron microscope and a general electron microscope at magnifications of 1000x and 3000x, respectively, and measured when possible to distinguish them. The surface hardness and deep cross-section hardness are Rockwell C hardness. The surface hardness was measured at the surface, and the deep cross-section hardness was measured at the thickness t / 4 of the deep cross-section. R / t is the bending workability and is the ratio of the bending radius to the steel sheet thickness. It is expressed based on the minimum bending radius at which no unevenness or cracks occur on the surface even after 90° bending. At this time, the bent specimen was processed to be long in the direction perpendicular to the rolling direction so that the bend line was parallel to the rolling direction. YS and TS are the yield strength and tensile strength measured in a tensile test in a direction perpendicular to the rolling direction, and the unit is MPa. The tensile test standard used was a JIS No. 5 standard test specimen.

[0114] Slab heating temperature RDT (℃) FDT (℃) FSB (times) ST (℃) Ms (℃) Primary cooling rate (℃ / s) Primary cooling stop temperature (℃) Secondary cooling rate (℃ / s) CT (℃) Invention example 1 1 2 3 1 1 0 1 9 3 6 3 9 2 0 2 8 4 7 9 2 9 5 1 6 1 2 2 Invention example 2 1 1 6 0 9 3 8 6 4 2 8 6 1 3 0 7 8 6 3 2 6 2 4 9 4 Invention example 3 1 2 1 7 1 0 2 9 3 1 3 9 1 6 2 9 0 7 3 0 8 3 0 5 1 Invention example 4 1 2 1 6 1 0109163908291683102253Invention Example 5119598090538763166733724106Invention Example 6118097988928892926431224138Invention Example 7127510619863965261762883088Invention Example 81269105996329482798229224148Invention Example 91227100091939162747728623 65 Invention Example 101261103697129562737028821105 Comparative Example 11229101793219292697229416340 Comparative Example 21256103995819292697629322280 Comparative Example 3128310659853975214662321866 Comparative Example 4118898590918913168233522139 Comparative Example 512161007923 0920279782921893Comparative Example 61312108910193949266822771756Comparative Example 712049798981889288722992899Comparative Example 81280107799219662277924726135Comparative Example 91240101192819052527926921111Comparative Example 101240103094919322731201200.594

[0115] Aspect Ratio (major axis / minor axis) Surface hardness (HrC) Core cross-sectional hardness (HrC) Surface-core hardness difference (%) Bendability (R / t) YS (MPa) TS (MPa) Core bainite + ferrite fraction (%) Core martensite fraction (%) Surface twin martensite fraction (%) Surface ferrite fraction (%)Invention Example 1440.947.313.52.3123914742981225Invention Example 2539.445.012.42.7114713974961120Invention Example 3541.547.813.13.01263151001001025Invention Example 44.439.646.314.43.3120014502981125Invention Example 5337.643.814.22.7111613676 941230Invention Example 62.444.449.610.43.713431584397920Invention Example 72.242.349.314.23.3133115740100930Invention Example 8440.344.79.83.31135138901001220Invention Example 94.441.449.015.53.3132215660100725Invention Example 10542.247.511.22.71247 14950100920Comparative Example 1538.942.58.54.3103313201382011Comparative Example 2740.043.37.74.7108413451183113Comparative Example 3541.547.813.25.312601507010000Comparative Example 42.428.431.28.91.0674920752842Comparative Example 5545.146.83.64.712181468010 008Comparative Example 61.544.347.46.65.712421491010000Comparative Example 71644.549.610.34.3134415856941230Comparative Example 81351.452.92.86.315351762010000Comparative Example 91451.152.83.26.015211750010000Comparative Example 10945.846.10.64.711881440010000

[0116] As shown in Table 1-3 above, it can be confirmed that all of the invention examples 1-10, whose steel composition components and manufacturing process conditions meet the scope of the present invention, have excellent mechanical properties.

[0117] In contrast, Comparative Examples 1 and 2 were composed of components satisfying Equation 1, but were cooled to a high temperature exceeding the coiling temperature range during the manufacturing process. As a result, a large amount of martensite in the core of the final microstructure was replaced by bainite, resulting in lower strength compared to the components, and ferrite was not formed within an appropriate range in the surface layer. Consequently, as the coiling temperature increased, strength and hardness decreased, but since the microstructural conditions were not satisfied, no effect of improving bendability was observed.

[0118] Comparative Example 3 used an alloy component exceeding the relational expression 1 related to the components. In addition, hardenable elements such as Mn, Cr, and Mo were synthetic components exceeding the scope of the invention, and the T value of relational expression 1 was excessively high. As a result, ferrite and twinned martensite were not formed at all on the surface layer, and although other manufacturing process conditions were satisfied, the difference between the surface hardness and the deep hardness exceeded 10%, and no improvement in bendability was observed.

[0119] Comparative Example 4, which is the exact opposite of Comparative Example 3, satisfies the range of individual alloy components under the condition that it falls short of the range of relational expression 1, but does not satisfy the component relational expression, and thus it was confirmed that a sufficient high-strength martensitic steel could not be produced. Most of the core formed a bainite structure, and accordingly, the strength and hardness were significantly lower than those of other inventions / comparative examples. Of course, the strength and hardness were lowered, but the workability was improved, and a large amount of ferrite and twinned martensite were formed in the surface layer, so the bendability was very good at 1.0, but it cannot be included in the invention examples that aim to improve bendability with high-strength steel.

[0120] Comparative Example 5 was a case where the number of FSB (finish rolling descaler water spray device) used to satisfy the microstructure composition of the surface and center under the manufacturing process conditions was insufficient, 0, and this was a case where ferrite + twin martensite could not be formed in the surface layer due to insufficient surface cooling. As a result, the difference in cross-sectional hardness between the surface and the core did not fall within the required range, resulting in poor bendability.

[0121] Comparative Example 6 is a case where the FDT (finish rolling end temperature) range calculated from the ST temperature defined in Equation 2 in the manufacturing process conditions was exceeded, and as in Comparative Example 5, ferrite + twin martensite was not formed in the surface layer, and as a result, the difference in hardness between the surface and deep cross-section did not fall within the required range, and as a result, the bendability also showed poor results.

[0122] Comparative Examples 7-9 are cases where the range of components such as Nb and Mo was exceeded, and the aspect ratio of the deep austenite became excessively large, and the effect of tissue refinement in a specific direction was increased, resulting in a deterioration in bendability. In addition, Comparative Examples 8-9 showed a result where the microstructure composition of the surface layer was formed differently from the intended one due to a large amount of alloy components exceeding the range of relation 1, resulting in even worse bendability.

[0123] In Comparative Example 10, an excessively high cooling rate was applied in the first cooling rate associated with phase transformation among the manufacturing process conditions, and thus a combination of surface ferrite and twin martensite was not formed in the microstructure. Consequently, even though the components and other manufacturing process conditions were all within the range, the result was extremely poor bendability.

[0124] Meanwhile, Fig. 1 is a graph showing the relationship between the bendability (R / t) according to the t / 4 deep cross-sectional hardness of the inventive examples and the comparative examples in the embodiments of the present invention. As shown in Fig. 1, as the hardness increases, the R / t increases overall, showing a tendency for the bendability to become inferior. However, the inventive examples have relatively low R / t values ​​for bendability compared to the tensile strength, and thus deviate from the continuous line of the graph, confirming that the bendability is superior to that of the comparative examples.

[0125] Figure 2 is a photograph of twinned martensite formed on the surface of Invention Example 1, taken using a transmission electron microscope.

[0126] Figure 3 is a photograph of ferrite formed on the surface of Invention Example 1, taken using a transmission electron microscope.

[0127] Figure 4 is a photograph of the surface cross-section microstructure distribution of Invention Example 1 and Comparative Example 3 taken with an electron microscope.

[0128] Figure 5 is a photograph showing carbides in the deep microstructure (auto-tempered martensite) of Invention Example 1.

[0129] Figure 6 is a photograph taken with an optical microscope of a surface with irregularities (bending radius of 9 mm) and a good surface without irregularities (bending radius of 7 mm) during a bending test of Invention Example 1.

[0130] As described above, the detailed description of the present invention has described preferred embodiments of the present invention. However, it will be apparent to those skilled in the art that various modifications may be made without departing from the scope of the present invention. Therefore, the scope of the rights of the present invention should not be limited to the described embodiments, but should be determined not only by the claims described below but also by equivalents thereof.

Claims

1. Contains, in wt%, C: 0.17 to 0.26%, Si: 0.01 to 0.5%, Mn: 0.3 to 2.0%, Cr: 0.005 to 0.5%, Mo: 0.005 to 0.3%, Nb: 0.001 to 0.01%, Ti: 0.005 to 0.08%, V: 0.005 to 0.2%, Al: 0.01 to 0.5%, P: 0.003 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, B: 0.0005 to 0.005%, the remainder being Fe and other unavoidable impurities, and satisfying the following Relationship Formula 1: A hot-rolled steel sheet having a microstructure of a steel sheet, wherein a surface layer from a surface of the steel sheet to a depth of 10 ㎛ includes, in terms of area %, ferrite: 15 to 50%, twinned martensite: 2 to 15%, residual bainite, martensite, and auto-tempered martensite, a core layer at a point t / 4 of the steel sheet includes, in terms of area %, martensite + auto-tempered martensite: 90% or more and at least one of ferrite and bainite: 10% or less, and an aspect ratio of old austenite of the core layer is 10 or less. [Relationship 1] 30 ≤T ≤ 80 T =(([C] / 10) 0.5 *1.07)*(0.7*[Si]+1)*(5*[Mn]+1)*(2.16*[Cr]+1)*(3*[Mo]+1)*25.4 *In the above equation 1, C, Si, Mn, Cr, and Mo are the weight% of the corresponding alloy elements. If not added, 0 is substituted.

2. A hot-rolled steel sheet in accordance with paragraph 1, wherein the hardness measured on the surface of the steel sheet is 9 to 25% lower than the deep cross-sectional hardness measured at t / 4 of the steel sheet.

3. A hot-rolled steel sheet satisfying bendability (R / t) <4 in the first paragraph.

4. A step of heating a slab containing, by weight%, C: 0.17 to 0.26%, Si: 0.01 to 0.5%, Mn: 0.3 to 2.0%, Cr: 0.005 to 0.5%, Mo: 0.005 to 0.3%, Nb: 0.001 to 0.01%, Ti: 0.005 to 0.08%, V: 0.005 to 0.2%, Al: 0.01 to 0.5%, P: 0.003 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, B: 0.0005 to 0.005%, the remainder being Fe and other unavoidable impurities, and satisfying the following relationship 1, to 1150 to 1350°C; A step of obtaining a bar by completing rough rolling of the heated slab at a rough rolling temperature (RDT) of 900 to 1100℃ based on 1 / 2t (t: thickness of steel); A step of performing surface cooling while descaling the entire width of the bar three or more times; A step of obtaining a hot rolled steel sheet by completing the finishing rolling at a finishing rolling temperature (FDT) of 800℃ to ST+50℃ that satisfies the following relational expression 2 based on 1 / 2t (t: thickness of steel) of the cooled bar above; A step of first cooling the hot-rolled steel sheet to a first cooling stop temperature of Ms~Ms+50℃ at a first average cooling rate of 60~90℃ / sec: and A method for manufacturing a hot-rolled steel sheet, comprising the step of secondarily cooling the first-cooled hot-rolled steel sheet at a second average cooling rate of 1 to 50°C / sec to a coiling temperature (CT) of 50°C to Ms-50°C and then coiling the hot-rolled steel sheet. [Relationship 1] 30 ≤ T ≤ 80 T =(([C] / 10) 0.5 *1.07)*(0.7*[Si]+1)*(5*[Mn]+1)*(2.16*[Cr]+1)*(3*[Mo]+1)*25.4 *In the above equation 1, C, Si, Mn, Cr, and Mo are the weight% of the corresponding alloy elements. If not added, 0 is substituted. [Relationship 2] 800℃~ST+50℃ ST = 800 + 140 [C] - 145 [Si] + 80 [Mn] - 30 [Cr] + 10 [Mo] * In the above equation 2, C, Si, Mn, Cr, and Mo are the weight% of the corresponding alloy elements. If not added, 0 is substituted.

5. A method for manufacturing a hot-rolled steel sheet, further comprising the step of pickling and oiling the steel sheet coiled after the second cooling in paragraph 4.

6. In the fourth paragraph, the coiled hot-rolled steel sheet has a steel sheet microstructure in which a surface layer from the steel sheet surface to a depth of 10 ㎛ includes, in terms of area %, ferrite: 15 to 50%, twinned martensite: 2 to 15%, residual bainite, martensite, and auto-tempered martensite, a deep layer at a point t / 4 of the steel sheet includes, in terms of area %, martensite + auto-tempered martensite: 90% or more and at least one of ferrite and bainite: 10% or less, and a hardness measured at the surface of the steel sheet is 9 to 25% lower than a deep layer cross-sectional hardness measured at t / 4 of the steel sheet, and an aspect ratio of old austenite in the deep layer is 10 or less.

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