Hot-rolled steel sheet and its manufacturing method

Abstract

TECHNICAL FIELD The present invention relates to a hot-rolled steel sheet and a manufacturing method thereof. BACKGROUND ART An embodiment of the present invention is to provide a hot-rolled steel sheet having excellent bending workability, strength, and hardness, and a manufacturing method thereof.

Description

[Technical Field]

[0001] The present invention relates to a hot-rolled steel sheet and a method for manufacturing the same. [Background technology]

[0002] Conventionally, high-hardness wear-resistant hot-rolled steel sheets have been used for parts requiring high wear resistance, mainly due to the martensite-based microstructure, which provides high strength and hardness. However, due to poor bending workability caused by high strength, processing is limited, resulting in the drawback of minimal processing. This is also a limiting factor for high-strength hot-rolled steels that utilize martensite as the main phase, and various technologies have been proposed to overcome this limitation.

[0003] In Patent Document 1, the alloy components of the steel are mainly controlled, such as C, Si, and Mn, and rolling at the austenite non-recrystallization temperature during hot rolling is minimized, thereby reducing the aspect ratio of prior austenite grains and suppressing the texture that strengthens anisotropy, thereby attempting to improve bendability.

[0004] In Patent Document 2, an attempt is made to improve bendability by forming a high-strength steel plate with tempered martensite in the center and ferrite and pearlite in the surface layer as the main phases in the thickness direction of the steel plate.

[0005] However, Patent Document 1 requires high-temperature rolling, which is a technique that requires controlling the texture of the steel, which is difficult to control. This technique has the drawback of being difficult to manufacture, and since the bainite phase is contained in the center as the main phase in addition to martensite, which ensures high hardness, it is difficult to ensure uniform hardness, making the steel unsuitable for use as wear-resistant steel.

[0006] The steel in Patent Document 2 has the drawback that the surface layer is composed of ferrite and pearlite, which are structures that are excessively soft compared to the center portion, and therefore has low hardness and is unsuitable for use as wear-resistant steel because deformation is concentrated in the surface layer during bending.

[0007] In addition, alloy elements such as Si, Mn, Mo, Cr, Cu, and Ni, which are mainly used to produce the above-mentioned high-hardness steels, are effective in improving hardness and formability, but adding large amounts of alloy elements to improve physical properties can lead to segregation of the alloy elements and uneven microstructures, resulting in poor bending workability. In particular, steels with high hardenability have the problem that they are sensitive to changes in their microstructures during cooling and form uneven low-temperature transformation structures, making it difficult to achieve better bending workability. [Prior art documents] [Patent documents]

[0008] [Patent Document 1] Japanese Patent Application Laid-Open No. 2013-117068 [Patent Document 2] Korean Patent Publication No. 10-2021-0088646 Summary of the Invention [Problem to be solved by the invention]

[0009] One embodiment of the present invention is to provide a hot rolled steel sheet and a method for manufacturing the same.

[0010] A preferred embodiment of the present invention is to provide a hot-rolled steel sheet having excellent bending workability, strength, and hardness, and a method for manufacturing the same. [Means for solving the problem]

[0011] One embodiment of the present invention is a sintered body 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 0.55%, Nb: 0.005 to 0.05%, 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%, and the balance being Fe and the like. and other unavoidable impurities, and satisfying the following Relational Expressions 1 and 2, the microstructure has, in area %, a center portion containing 90% or more of martensite and autotempered martensite in total and 10% or less of at least one of pearlite and bainite, and a surface layer portion containing 90% or more of bainite and 10% or less of at least one of ferrite, martensite and pearlite, the average thickness of the surface layer portion being 30 to 200 μm, and the average dislocation density of the surface layer portion being 1.7 × 10 14 ~3.0×10 14 m -2 and the central part of the hot rolled steel sheet has an aspect ratio of prior austenite of 5 or more.

[0012] [Equation 1] 0.1≦X=(Nb / 9+Ti* / 5+V / 5) / (C / 12+N / 14)≦0.6

[0013] [Equation 2] 1.5≦T=Mn+2.8Mo+1.5Cr+500B≦5.0

[0014] [Formula 1] Ti*=Ti-3.42N-1.5S (However, in the above relational formula 1, Ti* is represented by the above [formula 1], and the content of each alloy composition in the above relational formulas 1 and 2 and formula 1 means wt %.)

[0015] The Rockwell hardness of the central portion may be 44 to 50 HrC.

[0016] The surface layer portion may have a Rockwell hardness of 38 to 46 HrC.

[0017] The above-mentioned hot-rolled steel sheet may have a value of [bending workability (R / t)÷(tensile strength−1000)]×1000 of 5 or less.

[0018] Another embodiment of the present invention is a composition comprising, in weight percent, 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 0.55%, Nb: 0.005 to 0.05%, 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%, A step of heating a slab containing 0.01%, B: 0.0005 to 0.005%, the balance Fe and other inevitable impurities, and satisfying the following relations 1 and 2, to 1150 to 1350 ° C; a step of completing rough rolling of the heated slab at a rough rolling temperature (RDT) of 880 to SCT + 170 ° C based on 1 / 2t (t: thickness of steel material) to obtain a bar; a step of completing rough rolling of the bar at a rough rolling temperature (RDT) of 780 to SCT + 170 ° C based on 1 / 2t (t: thickness of steel material). 30 the step of starting cooling of the hot-rolled steel sheet at a cooling start temperature (WCT) of 700°C to SCT+10°C and primarily cooling the hot-rolled steel sheet to a primary cooling stop temperature of Ms to Ms+50°C at a primary average cooling rate of 50 to 100°C / sec; and the step of secondarily cooling the primarily cooled hot-rolled steel sheet at a secondary average cooling rate of 1 to 40°C / sec to a coiling temperature (CT) of 70°C to Ms-50°C, followed by coiling; wherein the surface temperature of the bar at the end of the rough rolling (RST) is controlled to be 750 to RDT-40°C, and the surface temperature of the hot-rolled steel sheet at the end of the finish rolling (FST) is controlled to be 700 to FDT-40°C.

[0019] [Equation 1] 0.1≦X=(Nb / 9+Ti* / 5+V / 5) / (C / 12+N / 14)≦0.6

[0020] [Relationship 2] 1.5≦T=Mn+2.8Mo+1.5Cr+500B≦5.0

[0021] [Formula 1] Ti*=Ti-3.42N-1.5S

[0022] [Formula 2] SCT(℃)=741+134C-137Si+75.4Mn-21.4Cr+24.8Mo-1391Nb-13Ti+19330B

[0023] [Formula 3] Ms(℃)=430-380C-13.4Si-47.3Mn-16Cr-24.2Mo (However, in the above relational formula 1, Ti* is represented by the above [formula 1], and the content of each alloy composition in the above relational formulas 1 and 2 and formulas 1 to 3 means wt %.)

[0024] The surface temperature of the bar and the hot-rolled steel sheet can be controlled by a water injection device.

[0025] After the coiling, the method may further include pickling and oiling the coiled hot-rolled steel sheet. [Effects of the Invention]

[0026] According to one embodiment of the present invention, a hot-rolled steel sheet and a manufacturing method thereof can be provided.

[0027] According to a preferred embodiment of the present invention, it is possible to provide a hot-rolled steel sheet having excellent bend formability and hardness, and a method for manufacturing the same. [Brief explanation of the drawings]

[0028] [Figure 1] 1 is a graph showing the relationship between bending workability (R / t) and tensile strength for Examples 1 to 10 according to one embodiment of the present invention and Comparative Examples 1 to 12. [Figure 2] 1 is a photograph of Example 1 according to one embodiment of the present invention, observed with an electron microscope. [Figure 3] 1 is a photograph of a surface layer portion of Example 1 according to one embodiment of the present invention, observed with an electron microscope. [Figure 4]1 is a photograph of the center of Example 1 according to one embodiment of the present invention, observed with an electron microscope. DETAILED DESCRIPTION OF THE INVENTION

[0029] Hereinafter, a hot-rolled steel sheet according to one embodiment of the present invention will be described. First, the alloy composition will be described. The contents of the alloy composition described below are in weight percent unless otherwise specified.

[0030] C: 0.17 to 0.26% C is the most economical and effective element for strengthening steel and has a significant impact on hardness. Increasing its content increases hardening ability, facilitating the formation of hard phases such as bainite and martensite in the microstructure, thereby increasing tensile strength. Furthermore, C forms fine precipitates together with Ti and Nb, which have a high affinity for C, and this precipitation strengthening increases both yield strength and tensile strength. However, if the C content exceeds 0.26%, the hardness of martensite itself increases excessively, resulting in excessive strength increase and reduced bending workability, potentially making it difficult to ensure sufficient weldability. On the other hand, if the C content is less than 0.17%, it becomes difficult to achieve sufficient strengthening. Therefore, the C content is preferably in the range of 0.17 to 0.26%. The lower limit of the C content is more preferably 0.175%, even more preferably 0.18%, and most preferably 0.185%. The upper limit of the C content is more preferably 0.25%, further preferably 0.24%, and most preferably 0.23%.

[0031] Si: 0.01 to 0.5% Si is an element advantageous for improving formability because it deoxidizes molten steel, exerts a solid-solution strengthening effect, and delays the formation of coarse carbides. If the Si content is less than 0.01%, the solid-solution strengthening effect and the formability improvement effect cannot be sufficiently obtained. On the other hand, if the Si content exceeds 0.5%, it is difficult to remove the red scale formed on the surface of the steel sheet during hot rolling, which may significantly deteriorate the surface quality of the steel sheet. Furthermore, there are problems such as reduced ductility and weldability. Therefore, the Si content is preferably in the range of 0.01 to 0.5%. The lower limit of the Si content is more preferably 0.012%, even more preferably 0.015%, and most preferably 0.02%. The upper limit of the Si content is more preferably 0.4%, even more preferably 0.35%, and most preferably 0.3%.

[0032] Mn: 0.3 to 2.0% Like Si, Mn is an effective element for solid solution strengthening of steel, increasing the hardening ability of steel and facilitating the formation of hard phases, bainite and martensite, during cooling after hot rolling. If the Mn content is less than 0.3%, the effects of solid solution strengthening and the formation of bainite and martensite are insufficient. On the other hand, if the Mn content exceeds 2.0%, the grain boundaries become brittle, causing problems such as cold cracking. Furthermore, excessive strength increases, potentially making it difficult to ensure sufficient formability. During continuous casting, large segregations develop in the center of the slab during casting, leading to the formation of non-uniform microstructures in the thickness direction during cooling after hot rolling, resulting in poor bending workability. In particular, it becomes difficult to uniformly form a microstructure throughout the entire length and width of the hot-rolled steel sheet during cooling. Therefore, the Mn content is preferably in the range of 0.3 to 2.0%. The lower limit of the Mn content is more preferably 0.35%, even more preferably 0.4%, and most preferably 0.45%. The upper limit of the Mn content is more preferably 1.9%, further preferably 1.85%, and most preferably 1.8%.

[0033] Cr: 0.005 to 0.5% Cr strengthens the steel in solid solution and delays the ferrite phase transformation during cooling, contributing to the formation of martensite and bainite. If the Cr content is less than 0.005%, the effects of solid solution strengthening and the formation of martensite and bainite are insufficient. On the other hand, if the Cr content exceeds 0.5%, segregation at the center of the steel, like Mn, develops significantly, resulting in a non-uniform microstructure in the thickness direction and reduced bending workability. Therefore, the Cr content is preferably in the range of 0.005 to 0.5%. The lower limit of the Cr content is more preferably 0.007%, even more preferably 0.008%, and most preferably 0.01%. The upper limit of the Cr content is more preferably 0.4%, even more preferably 0.35%, and most preferably 0.3%.

[0034] Mo: 0.005 to 0.55% Mo increases the hardenability of steel and facilitates the formation of martensite and bainite. If the Mo content is less than 0.005%, the above-mentioned effects cannot be fully achieved. On the other hand, if the Mo content exceeds 0.55%, martensite is formed in the surface layer due to an excessive increase in hardenability, which rapidly reduces bending workability, is economically disadvantageous, and may make it difficult to ensure sufficient weldability. Therefore, the Mo content is preferably in the range of 0.005 to 0.55%. The lower limit of the Mo content is more preferably 0.01%, even more preferably 0.02%, and most preferably 0.03%. The upper limit of the Mo content is more preferably 0.52%, even more preferably 0.5%, and most preferably 0.45%.

[0035] Nb: 0.005 to 0.05% Nb, along with Ti and V, is a typical precipitation strengthening element. It precipitates as a precipitate during hot rolling and exerts a grain refinement effect by delaying recrystallization, thereby effectively improving the strength and impact toughness of steel. If the Nb content is less than 0.005%, the above-mentioned effect cannot be sufficiently obtained. On the other hand, if the Nb content exceeds 0.05%, coarse composite precipitates are formed during hot rolling, deteriorating bending workability. Therefore, the Nb content is preferably in the range of 0.005 to 0.05%. The lower limit of the Nb content is more preferably 0.007%, even more preferably 0.008%, and most preferably 0.01%. The upper limit of the Nb content is more preferably 0.04%, even more preferably 0.03%, and most preferably 0.02%.

[0036] Ti: 0.005 to 0.08% Ti, along with Nb and V, is a typical precipitation strengthening element and forms coarse TiN due to its strong affinity with nitrogen. TiN has the effect of suppressing grain growth during the heating process for hot rolling. Furthermore, Ti remaining after reacting with nitrogen dissolves in the steel and bonds with carbon to form TiC precipitates, making it a useful component for improving the strength of the steel. If the Ti content is less than 0.005%, the effects of suppressing grain growth and improving strength are not sufficiently achieved. On the other hand, if the Ti content exceeds 0.08%, coarse TiN is generated, and the precipitates become coarse, deteriorating bending workability during forming. Therefore, the Ti content is preferably in the range of 0.005 to 0.08%. The lower limit of the Ti content is more preferably 0.01%, even more preferably 0.015%, and most preferably 0.02%. The upper limit of the Ti content is more preferably 0.07%, even more preferably 0.06%, and most preferably 0.045%.

[0037] V: 0.005 to 0.2% V, along with Nb and Ti, is a typical precipitation strengthening element. Although it hardly precipitates during hot rolling, it forms precipitates after hot coiling, cooling, or tempering, thereby improving the strength of the steel. Therefore, it effectively improves strength without increasing deformation resistance and rolling load due to delayed recrystallization during hot rolling. If the V content is less than 0.005%, the strength improvement effect is insufficient. On the other hand, if the V content exceeds 0.2%, coarse precipitates are formed, deteriorating bending workability and being economically disadvantageous. Therefore, the V content is preferably in the range of 0.005 to 0.2%. The lower limit of the V content is more preferably 0.006%, even more preferably 0.008%, and most preferably 0.01%. The upper limit of the V content is more preferably 0.2%, even more preferably 0.1%, and most preferably 0.05%.

[0038] Al: 0.01 to 0.5% The Al is an element added mainly for deoxidation. If the Al content is less than 0.01%, the deoxidation effect is insufficient. On the other hand, if the Al content exceeds 0.5%, the Al combines with nitrogen to form excessive AlN, which makes corner cracks more likely to occur in the slab during continuous casting and defects due to the formation of inclusions more likely to occur. Therefore, the Al content is preferably in the range of 0.01 to 0.5%. The lower limit of the Al content is more preferably 0.015%, and even more preferably 0.02%. The upper limit of the Al content is more preferably 0.1%, even more preferably 0.08%, and most preferably 0.05%.

[0039] P: 0.003 to 0.05% Like Si, P simultaneously exerts the effects of solid solution strengthening and ferrite transformation acceleration. However, controlling the P content to less than 0.003% requires high production costs, which is economically disadvantageous and insufficient for achieving sufficient strength. On the other hand, if the P content exceeds 0.05%, embrittlement due to grain boundary segregation may occur, which makes it easier for fine cracks to occur during bending, significantly reducing ductility and impact resistance. 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%, even more preferably 0.007%, and most preferably 0.01%. The upper limit of the P content is more preferably 0.03%.

[0040] S: 0.001 to 0.01% S is an impurity present in steel, and if its content exceeds 0.01%, it bonds with Mn and other elements to form nonmetallic inclusions, which can easily cause microcracks during bending of the steel and significantly reduce impact resistance. In the present invention, the lower limit of the S content is not particularly limited. However, controlling the S content to less than 0.001% requires a long steelmaking operation time and reduces productivity. Taking this into consideration, the lower limit of the S content can be set to 0.001%. Therefore, the S content is preferably in the range of 0.001 to 0.01%. The lower limit of the S content is more preferably 0.002%. The upper limit of the S content is more preferably 0.008%, even more preferably 0.006%, and most preferably 0.005%.

[0041] N: 0.001 to 0.01% N, along with C, is a typical solid-solution strengthening element and forms coarse precipitates together with Ti, Al, etc. If the N content is less than 0.001%, not only are the effects of solid-solution strengthening and precipitate formation insufficient, but controlling the N content to less than 0.001% requires a long time during steelmaking operation, resulting in reduced productivity. On the other hand, although N generally has a better solid-solution strengthening effect than carbon, if the N content exceeds 0.01%, there is a problem of a significant decrease in toughness. Therefore, the N content is preferably in the range of 0.001 to 0.01%. The lower limit of the N content is more preferably 0.002%, and even more preferably 0.003%. The upper limit of the N content is more preferably 0.008%, even more preferably 0.007%, and most preferably 0.006%.

[0042] B: 0.0005 to 0.005% When present in a solid solution state in steel, B segregates mainly at grain boundaries, stabilizing the grain boundaries and improving the brittleness of steel. It also stabilizes solute N and suppresses the formation of coarse AlN nitrides. It also delays the ferrite phase transformation, effectively promoting the formation of hard phases such as bainite and martensite. If the B content is less than 0.0005%, the effects of improving brittleness, suppressing the formation of coarse AlN nitrides, and promoting the formation of bainite and martensite are not fully achieved. On the other hand, if the B content exceeds 0.005%, the aforementioned effects are no longer enhanced, and there is a drawback in that ductility decreases and formability deteriorates. Therefore, the B content is preferably in the range of 0.0005 to 0.005%. The lower limit of the B content is more preferably 0.0006%, even more preferably 0.0008%, and most preferably 0.001%. The upper limit of the B content is more preferably 0.004%, even more preferably 0.003%.

[0043] The remaining component is iron (Fe). However, in a normal manufacturing process, unintentional impurities may be inevitably mixed in from the raw materials or the surrounding environment, and this cannot be excluded. Since these impurities are known to anyone skilled in the normal manufacturing process, not all of the contents thereof are specifically mentioned in this specification.

[0044] On the other hand, it is preferable that the hot-rolled steel sheet of the present invention satisfies the above-mentioned alloy composition and also satisfies the following relational expressions 1 and 2.

[0045] [Equation 1] 0.1≦X=(Nb / 9+Ti* / 5+V / 5) / (C / 12+N / 14)≦0.6

[0046] [Formula 1] Ti*=Ti-3.42N-1.5S

[0047] The above relational expression 1 is intended to balance precipitation and hardenability to improve bendability. If the X value is less than 0.1, grain growth during reheating is facilitated, recrystallization becomes nonuniform during hot rolling, resulting in the formation of locally coarse grains, resulting in an excessive amount of solute C and solute N, and a tendency for the hard phase to have a high hardness, which may result in poor bendability. If the X value exceeds 0.6, the formation of precipitates increases excessively, and furthermore, the untransformed phase lacks solute C and solute N atoms during cooling of the hot-rolled steel sheet, making it difficult for the hard phase to form stably, resulting in brittle grain boundaries and poor bendability. Therefore, the X value is preferably in the range of 0.1 to 0.6. The lower limit of the X value is more preferably 0.13, even more preferably 0.15, and most preferably 0.18. The upper limit of the X value is more preferably 0.58, even more preferably 0.56, and most preferably 0.55.

[0048] [Equation 2] 1.5≦T=Mn+2.8Mo+1.5Cr+500B≦5.0

[0049] The above-mentioned relational expression 2 represents a factor of the combination of alloying elements that can maintain the formation of the hard phases bainite and martensite at an appropriate level in the microstructure of the steel of the present invention. If the T value is less than 1.5, the hard phases cannot be sufficiently secured, making it difficult to obtain the desired hardness value. On the other hand, as the T value increases, the formation of the hard phases bainite, martensite, and MA phase increases, and the hardness value of each hard phase also increases. Therefore, a larger T value is advantageous for ensuring strength and hardness. However, if the T value exceeds 5.0, there are problems such as deterioration of bending workability and increased material variation across the entire length and width of the hot-rolled steel sheet. Therefore, the T value is preferably in the range of 1.5 to 5.0. The lower limit of the T value is more preferably 1.7, even more preferably 2.0, and most preferably 2.5. The upper limit of the T value is more preferably 4.9, even more preferably 4.7, and most preferably 4.5.

[0050] The microstructure of the hot-rolled steel sheet of the present invention preferably contains, in area percentages, 90% or more of martensite and autotempered martensite combined and 10% or less of at least one of pearlite and bainite in the center, and 90% or more of bainite and 10% or less of at least one of ferrite, martensite, and pearlite in the surface layer. If the combined area percentage of martensite and autotempered martensite in the microstructure of the center is less than 90% or if the combined area percentage of at least one of pearlite and bainite exceeds 10%, the hardness of the center becomes excessively low, making it difficult to achieve the desired high hardness. The combined area percentage of martensite and autotempered martensite in the microstructure of the center is more preferably 92% or more, even more preferably 94% or more, and most preferably 95% or more. The combined area percentage of at least one of ferrite, martensite, and pearlite in the microstructure of the center is more preferably 8% or less, even more preferably 6% or less, and most preferably 5% or less. If the fraction of bainite in the microstructure of the surface layer portion is less than 90%, or if at least one of ferrite, martensite, and pearlite exceeds 10%, there is a drawback in that the bendability of the surface layer portion is reduced, resulting in a reduction in overall bendability. The fraction of bainite in the microstructure of the surface layer portion is more preferably 92% or more, even more preferably 94% or more, and most preferably 95% or more. The fraction of at least one of ferrite, martensite, and pearlite in the microstructure of the surface layer portion is more preferably 8% or less, even more preferably 6% or less, and most preferably 5% or less.

[0051] The average thickness of the surface layer portion is preferably 30 to 200 μm. If the average thickness of the surface layer portion is less than 30 μm, a sufficient soft surface layer portion cannot be ensured, which may result in deterioration of bending workability. In the present invention, there is no particular limitation on the upper limit of the average thickness of the surface layer portion, but it is difficult for the average thickness of the surface layer portion to exceed 200 μm in terms of the manufacturing process. The lower limit of the average thickness of the surface layer portion is more preferably 32 μm, even more preferably 35 μm, and most preferably 40 μm. The upper limit of the average thickness of the surface layer portion is more preferably 150 μm, even more preferably 120 μm, and most preferably 100 μm. That is, in the present invention, the surface layer portion refers to the region [(from the surface of the steel sheet to 30 μm in the thickness direction) to (from the surface of the steel sheet to 200 μm in the thickness direction)], and the central portion refers to the region outside the surface layer portion. However, although an oxide layer may be formed on the surface of the hot-rolled steel sheet of the present invention, the surface layer portion does not include the oxide layer, and therefore the surface layer portion can be a region from just below the oxide layer of the steel sheet to a depth of 30 to 200 μm in the thickness direction.

[0052] On the other hand, the present inventors have confirmed that the average dislocation density of the surface layer (hereinafter also referred to as "GND (Geometrical Necessary Dislocation)") is an important factor in the balance between the strength and bending workability of a steel material. More specifically, when the average dislocation density of the surface layer is 1.7 × 10 14 ~3.0×10 14 m -2 It is preferable that the average dislocation density of the surface layer portion is 1.7×10 14 m -2 If the bending strength is less than 3.0 × 10, deformation is concentrated in the surface layer, its continuity decreases, and bending workability deteriorates. 14 m -2 If the average dislocation density of the surface layer exceeds 1.7×10, the surface layer is not softened and becomes brittle to bending deformation. 14 ~3.0×10 14 m -2 The lower limit of the average dislocation density of the surface layer portion is preferably 1.8×10 14 m -2More preferably, it is 1.9×10 14 m -2 More preferably, it is 2.0 × 10 14 m -2 The upper limit of the average dislocation density in the surface layer is 2.9 × 10 14 m -2 More preferably, it is 2.8×10 14 m -2 It is most preferable that the average dislocation density is 0.05. Meanwhile, the average dislocation density can be calculated using kernel average misorientation (KAM) data measured by EBSD as shown in Equation 4 below. For convenience, this calculation can be performed using OIM analysis software, which is software for analyzing the EBSD measurement results. TM The EBSD measurement can be performed using a cross section parallel to the rolling direction at a position ¼ of the thickness of the steel sheet as a reference.

[0053] [Formula 4] GND(m -2 )=2θ / ub (In the above formula 4, θ is the average misorientation (KAM values), u is the unit length (step size in the EBSD measurement), and b is the burgers vector.)

[0054] The aspect ratio of the prior austenite in the central portion is preferably 5 or more. In this way, by increasing the aspect ratio of the prior austenite, it is possible to achieve grain refinement in the minor axis direction of the crystal grains and improve bending workability. If the aspect ratio of the prior austenite in the central portion is less than 5, it may be difficult to fully achieve the effect of improving bending workability. The aspect ratio of the prior austenite in the central portion is more preferably 6 or more, even more preferably 7 or more, and most preferably 8 or more. On the other hand, in the present invention, the larger the aspect ratio of the prior austenite in the central portion, the more advantageous effects are exhibited, so there is no particular upper limit. However, it is difficult for the aspect ratio of the prior austenite in the central portion to exceed 30 due to the manufacturing process.

[0055] The hot-rolled steel sheet of the present invention provided as described above may have a Rockwell hardness of 44 to 50 HrC at the center portion and a Rockwell hardness of 38 to 46 HrC at the surface layer portion, and the hot-rolled steel sheet may have a [bending workability (R / t) ÷ (tensile strength - 1000)] × 1000 of 5 or less. The Rockwell hardness may be measured by measuring five points on the surface using a Rockwell hardness tester (C scale) in accordance with ASTM-E18-22 and calculating the average value.

[0056] A method for manufacturing a hot-rolled steel sheet according to one embodiment of the present invention will be described below.

[0057] First, a slab satisfying the above-described alloy composition and Relational Expressions 1 and 2 is heated to 1150 to 1350°C. If the slab heating temperature is lower than 1150°C, precipitates are not sufficiently redissolved, resulting in reduced precipitate formation in processes after hot rolling, leaving coarse TiN particles. Furthermore, the slab is not sufficiently heated, making it difficult to maintain a constant temperature during hot rolling. On the other hand, if the slab heating temperature exceeds 1350°C, abnormal grain growth of austenite grains occurs, resulting in a decrease in strength. Therefore, the slab heating temperature is preferably in the range of 1150 to 1350°C. The lower limit of the slab heating temperature is more preferably 1155°C, and even more preferably 1160°C. The upper limit of the slab heating temperature is more preferably 1340°C, even more preferably 1330°C, and most preferably 1320°C.

[0058] The heated slab is then rough-rolled at a rough rolling temperature (RDT) of 880 to SCT+170°C based on 1 / 2t (t: steel thickness) to obtain a bar. If the rough rolling temperature is less than 880°C, an excessively large rolling load may cause problems in equipment operation, while if it exceeds SCT+170°C, it is difficult to sufficiently elongate the prior austenite, making it difficult to improve bending workability. Therefore, the rough rolling temperature is set to 880 to SCT+170°C. 170 °C range. The lower limit of the rough rolling temperature is more preferably 890°C, even more preferably 900°C, and most preferably 910°C. The upper limit of the rough rolling temperature is more preferably SCT + 165°C, even more preferably SCT + 160°C, and most preferably SCT + 155°C. The SCT can be calculated by Equation 2 shown below. Meanwhile, the above-mentioned 1 / 2t refers to a point halfway from the surface of the steel sheet in the thickness direction.

[0059] [Formula 2] SCT(℃)=741+134C-137Si+75.4Mn-21.4Cr+24.8Mo-1391Nb-13Ti+19330B

[0060] In the present invention, the surface temperature (RST) of the bar at the end of the rough rolling is preferably controlled to 750 to RDT - 40°C. This is because by controlling the surface temperature of the steel sheet lower than the temperature at the center, the temperature difference at the surface during the cooling stage after the rolling process is smaller than that at the center, thereby forming bainite instead of martensite. If the surface temperature of the bar at the end of the rough rolling is less than 750°C, an excessively large rolling load may cause problems in equipment operation, while if it exceeds RDT - 40°C, there is a disadvantage in that bainite cannot be sufficiently formed in the surface layer. Therefore, the surface temperature of the bar at the end of the rough rolling is preferably in the range of 750 to RDT - 40°C. The lower limit of the surface temperature of the bar at the end of the rough rolling is more preferably 765°C, even more preferably 780°C, and most preferably 800°C. The upper limit of the surface temperature of the bar at the end of the rough rolling is more preferably RDT - 42°C, even more preferably RDT - 45°C, and most preferably RDT - 50°C. In the present invention, there is no particular limitation on the method for controlling the surface temperature of the bar at the end of the rough rolling, and for example, a water injection device such as a descaler can be used.

[0061] The bar is then finish-rolled at a finish rolling temperature (FDT) of 780 to SCT + 30°C based on 1 / 2t (t: steel thickness) to obtain a hot-rolled steel sheet. If the finish rolling temperature is less than 780°C, an excessively large rolling load may cause problems in equipment operation, while if it exceeds SCT + 30°C, it is difficult to sufficiently elongate the prior austenite, making it difficult to improve bending workability. Therefore, the finish rolling temperature is preferably in the range of 780 to SCT + 30°C. The lower limit of the finish rolling temperature is more preferably 790°C, even more preferably 800°C, and most preferably 810°C. The upper limit of the finish rolling temperature is more preferably SCT + 25°C, even more preferably SCT + 22°C, and most preferably SCT + 20°C.

[0062] In the present invention, the surface temperature (FST) of the hot rolled steel sheet at the end of the above-mentioned finish rolling is preferably controlled to be 700 to FDT - 40°C. This is because, as described above, by controlling the surface temperature of the steel sheet to be lower than the temperature at the center, the temperature difference at the surface during the cooling stage after the rolling process is smaller than that at the center, thereby forming bainite instead of martensite. If the surface temperature of the bar at the end of the above-mentioned finish rolling is less than 700°C, an excessively large rolling load may cause problems in equipment operation, and if it exceeds FDT - 40°C, there is a disadvantage that bainite cannot be sufficiently formed in the surface layer. Therefore, hot rolled steel plate The surface temperature of the bar at the end of the finish rolling is preferably in the range of 700 to FDT - 40°C. The lower limit of the surface temperature of the bar at the end of the finish rolling is more preferably 710°C, even more preferably 730°C, and most preferably 750°C. The upper limit of the surface temperature of the bar at the end of the finish rolling is more preferably FDT - 42°C, even more preferably FDT - 43°C, and most preferably FDT - 45°C. In the present invention, there are no particular limitations on the method for controlling the surface temperature of the bar at the end of the finish rolling, and for example, a water injection device such as a descaler can be used.

[0063] The hot-rolled steel sheet is then cooled at a cooling start temperature (WCT) of 700°C to SCT+10°C, and then primarily cooled at a primary average cooling rate of 50 to 100°C / sec to a primary cooling stop temperature of Ms to Ms+50°C. If the primary cooling start temperature is less than 700°C, high-temperature phases such as ferrite and pearlite are formed rather than low-temperature phases such as bainite and martensite, resulting in a significant decrease in strength and hardness. If the primary cooling start temperature is greater than SCT+10°C, the driving force for the formation of martensite rather than bainite in the surface layer becomes stronger, resulting in insufficient bainite formation in the surface layer or an insufficient thickness of the surface layer. Therefore, the primary cooling start temperature is preferably in the range of 700°C to SCT+10°C. The lower limit of the primary cooling start temperature is more preferably 710°C, even more preferably 730°C, and most preferably 750°C. The upper limit of the primary cooling start temperature is more preferably SCT + 5°C, even more preferably SCT + 2°C, and most preferably SCT + 10°C. If the primary cooling stop temperature is lower than Ms, high-temperature phases such as ferrite or pearlite are formed before low-temperature phases such as martensite or bainite are formed, resulting in a significant decrease in the strength and hardness of the steel sheet or in a non-uniform material. If the temperature exceeds Ms + 50°C, martensite is not sufficiently formed in the center, resulting in a decrease in hardness. Therefore, the primary cooling stop temperature is preferably in the range of Ms to Ms + 50°C. The lower limit of the primary cooling stop temperature is more preferably Ms + 5°C, even more preferably Ms + 8°C, and even more preferably Ms + 10°C. The upper limit of the primary cooling stop temperature is more preferably Ms + 45°C, even more preferably Ms + 40°C, and most preferably Ms + 30°C. If the primary average cooling rate is less than 50°C / sec, the formation of martensite and bainite may become uneven, and the driving force for forming the low-temperature phase may be low, making it difficult to obtain sufficient strength and hardness.If the primary average cooling rate exceeds 100°C / sec, excessively fast phase transformation may cause a sudden change in the volume of the steel plate, increasing the risk of accidents such as being caught in the equipment.Therefore, the primary average cooling rate is preferably in the range of 50 to 100°C / sec. The lower limit of the primary average cooling rate is more preferably 52°C / sec, even more preferably 55°C / sec, and most preferably 60°C / sec. The upper limit of the primary average cooling rate is more preferably 97°C / sec, even more preferably 93°C / sec, and most preferably 90°C / sec.

[0064] The primarily cooled hot-rolled steel sheet is then secondarily cooled at a secondary average cooling rate of 1 to 40°C / sec to a coiling temperature (CT) of 70°C to Ms-50°C, and then coiled. The reason for making the secondary average cooling rate lower than the primary average cooling rate is that, once sufficient martensite has already been formed in the center, any faster cooling would only deteriorate the shape quality of the steel sheet without providing any additional benefits, and therefore, from the perspective of productivity, it is advantageous to switch to slower cooling. If the secondary average cooling rate is less than 1°C / sec, the cooling section becomes excessively long, which can make equipment operation difficult. If it exceeds 40°C / sec, it becomes difficult to achieve uniform cooling. Therefore, the secondary average cooling rate is preferably in the range of 1 to 40°C / sec. The lower limit of the secondary average cooling rate is more preferably 2°C / sec, even more preferably 3°C / sec, and most preferably 5°C / sec. The upper limit of the secondary average cooling rate is more preferably 39°C / sec, even more preferably 37°C / sec, and most preferably 35°C / sec. Meanwhile, in order to overcome poor workability, wear-resistant steels with high hardness are generally produced by heat treatment to convert the microstructure to tempered martensite. However, when heat treatment is performed, the surface layer is decarburized, and the surface hardness is excessively reduced to the level of ferrite, resulting in a significant decrease in the functionality of the wear-resistant steel. To prevent this, it is preferable to omit heat treatment in the present invention. However, if the heat treatment is omitted, the strength of the steel sheet becomes excessively high, resulting in poor bending workability. Therefore, it is preferable to control the coiling temperature to 70°C to Ms-50°C in order to impart an auto-tempering (self-tempering) effect to the bainite in the surface layer and the martensite in the center. If the coiling temperature is less than 70°C, the autotempering effect is not large, making it difficult to improve bending workability, and if it exceeds Ms-50°C, the stress formed by the phase transformation inside the low-temperature phase is insufficient, making it possible that high hardness cannot be obtained. Therefore, the coiling temperature is preferably in the range of 70°C to Ms-50°C.The lower limit of the winding temperature is more preferably 80°C, even more preferably 90°C, and most preferably 100°C. The upper limit of the winding temperature is more preferably Ms-60°C, even more preferably Ms-80°C, and most preferably Ms-100°C. Ms can be calculated using Equation 3 as follows:

[0065] [Formula 3] Ms(℃)=430-380C-13.4Si-47.3Mn-16Cr-24.2Mo

[0066] After the coiling, the method may further include pickling and oiling the coiled hot-rolled steel sheet. The pickling and oiling steps are not particularly limited in the present invention, and any method commonly used in the art may be used. [Example]

[0067] The present invention will be described in more detail with reference to the following examples. However, it should be noted that the following examples are intended to illustrate and explain the present invention in more detail, and are not intended to limit the scope of the present invention. The scope of the present invention is determined by the matters described in the claims and matters that can be reasonably inferred therefrom.

[0068] (Example) Slabs having the alloy compositions shown in Tables 1 and 2 below were prepared, and then hot-rolled steel sheets were produced under the conditions shown in Tables 3 and 4 below. The bar temperatures at the end of rough rolling and the hot-rolled steel sheet temperatures at the end of finish rolling shown in Tables 3 and 4 below were based on 1 / 2t (t: steel thickness), and the slab heating temperature, cooling start temperature, and coiling temperature were also based on 1 / 2t (t: steel thickness).

[0069] The microstructure and mechanical properties of the hot-rolled steel sheets thus produced were measured, and the results are shown in Table 5.

[0070] The type and fraction of microstructure, the aspect ratio of the prior austenite in the center, and the average thickness of the surface layer were measured using an electron microscope after specimens were taken from the surface layer and center (1 / 4 of the thickness of the steel sheet) of the hot-rolled steel sheet. The aspect ratio of the prior austenite in the center was measured in the transverse direction (TD), which is the side of the steel sheet, so that the horizontal direction in the image was the rolling direction (RD) and the vertical direction was the normal direction (ND).

[0071] The average dislocation density (GND) was measured using electron backscattered diffraction (EBSD, JEOL JSM-7001F) at a position 1 / 4 of the thickness of the hot-rolled steel sheet, based on a cross section parallel to the rolling direction, and then analyzed by OIM analysis. TM Calculations were performed using EDAX.

[0072] The hardness was measured at five points on the surface and center (at 1 / 4 of the thickness of the steel plate) of the hot-rolled steel plate using a Rockwell hardness tester (C scale) in accordance with ASTM-E18-22, and the average value was calculated.

[0073] Bending workability (R / t) refers to the ratio of bending radius (R) to steel sheet thickness (t), and is expressed as the minimum bending radius at which cracks do not occur on the surface even after a 90-degree bending test of the steel sheet. In this case, the bending test specimen was processed to be long in a direction perpendicular to the rolling direction, and the bending test was performed so that the bend line of the bending test specimen was parallel to the rolling direction.

[0074] The tensile strength and yield strength were measured using a tensile testing machine after a JIS No. 5 test piece was taken from the center (at a position 1 / 4 of the thickness of the steel plate).

[0075] [Table 1]

[0076] [Table 2]

[0077] [Table 3]

[0078] [Table 4]

[0079] [Table 5]

[0080] As can be seen from Tables 1 to 5 above, Examples 1 to 10, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, satisfy the microstructure conditions of the present invention and also have excellent mechanical properties.

[0081] In Comparative Example 1, the driving force for forming precipitates was excessive because the value of X in Relational Formula 1 was too high, and as a result, carbon and nitrogen could not fully play their role in improving strength in martensite and bainite, resulting in insufficient hardness. Also, it can be seen that the GND did not reach the range of the present invention.

[0082] In Comparative Example 2, the T value in Relational Formula 2 was too high, so bainite or martensite could not be properly formed, and the entire steel sheet was composed almost entirely of martensite without a surface layer, resulting in high surface hardness and poor bending workability. In addition, it was confirmed that the GND was outside the range of the present invention.

[0083] In Comparative Example 3, the X value in Relational Formula 1 was too low, so there was almost no driving force for forming precipitates, and as a result, the role of carbon and nitrogen in improving strength in martensite and bainite became excessive, resulting in deterioration of bending workability. Also, it was confirmed that the GND was outside the range of the present invention.

[0084] In Comparative Example 4, the T value in Relational Formula 2 was too low, so bainite or martensite could not be properly formed, resulting in an excessively low hardness. It can also be seen that the GND did not reach the range of the present invention.

[0085] Although Comparative Example 5 satisfied the alloy composition of the present invention, the cooling start temperature was high, which caused the surface layer to cool rapidly, resulting in a stronger driving force for the formation of martensite rather than bainite, and the thickness of the surface layer was insufficient, resulting in a deterioration in bending workability.

[0086] Although Comparative Example 6 satisfies the alloy composition of the present invention, the high finish rolling temperature resulted in a small aspect ratio of prior austenite in the martensite at the center, at 3, which resulted in a deterioration in bending workability compared to strength.

[0087] Although Comparative Examples 7 and 8 satisfied the alloy composition of the present invention, the temperature difference between RDT and RST or FDT and FST was not sufficient, which reduced the driving force for forming bainite in the surface layer and prevented the formation of a surface layer with sufficient thickness. This resulted in a deterioration in bending workability.

[0088] In Comparative Examples 9 and 10, the carbon content was too high, and therefore the hardness of both the surface layer and the center was excessively high, resulting in poor bending workability, even though all of the manufacturing conditions of the present invention were met. It was also confirmed that the GND was outside the range of the present invention.

[0089] In Comparative Example 11, although the alloy composition of the present invention was satisfied, the coiling temperature was too low, so the auto-tempering effect was not obtained, and it was found that the bending workability was deteriorated compared to the strength.

[0090] In Comparative Example 12, although the alloy composition satisfies the requirements of the present invention, the coiling temperature is too high, so the driving force for forming martensite and bainite is weak, resulting in low hardness and strength. It is also clear that the GND does not reach the range of the present invention.

[0091] Fig. 1 is a graph showing the relationship between bending workability (R / t) and tensile strength for invention examples 1 to 10 and comparative examples 1 to 12. As shown in Fig. 1, in the case of comparative examples 1 to 12, there is a tendency for bending workability (R / t) to deteriorate overall as the tensile strength increases, but in the case of invention examples 1 to 10, the bending workability is relatively good compared to the tensile strength, and the ratio [bending workability (R / t) ÷ (tensile strength - 1000)] × 1000 is 5 or less, which confirms that the bending workability is excellent.

[0092] Fig. 2 is a photograph of Example 1 observed under an electron microscope. Fig. 3 is a photograph of the surface layer of Example 1 observed under an electron microscope. Fig. 4 is a photograph of the center portion of Example 1 observed under an electron microscope. As can be seen from Figs. 2 to 4, in the case of Example 1, it can be confirmed that the microstructures of the surface layer and center portion that the present invention aims to achieve are formed.

Claims

1. 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 0.55%, Nb: 0.005 to 0.05%, 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 balance being Fe and other unavoidable impurities, The following relations 1 and 2 are satisfied: The microstructure includes, in area percentage, a center portion including 90% or more of martensite and autotempered martensite in total, and 10% or less of at least one of pearlite and bainite, and a surface layer portion including 90% or more of bainite and 10% or less of at least one of ferrite, martensite, and pearlite, The average thickness of the surface layer portion is 30 to 200 μm, The average dislocation density of the surface layer is 1.7×10 １４ ~3.0 x 10 １４ m －２ and The hot-rolled steel sheet has a central portion in which the aspect ratio of prior austenite is 5 or more. [Relationship 1] 0.1≦X=(Nb / 9+Ti* / 5+V / 5) / (C / 12+N / 14)≦0.6 [Relationship 2] 1.5≦T=Mn+2.8Mo+1.5Cr+500B≦5.0 [Formula 1] Ti*=Ti-3.42N-1.5S (However, in the above-mentioned relational formula 1, Ti* is represented by the above-mentioned [formula 1], and the contents of each alloy composition in the above-mentioned relational formulas 1 and 2 and formula 1 mean wt %.)

2. The hot-rolled steel sheet according to claim 1, wherein the center portion has a Rockwell hardness of 44 to 50 HrC.

3. The hot-rolled steel sheet according to claim 1, wherein the surface layer portion has a Rockwell hardness of 38 to 46 HrC.

4. The hot-rolled steel sheet according to claim 1, wherein the heat-rolled steel sheet has a value of [bending workability (R / t) ÷ (tensile strength − 1000)] × 1000 of 5 or less.

5. a step of heating a slab containing, in weight percent, 0.17 to 0.26% C, 0.01 to 0.5% Si, 0.3 to 2.0% Mn, 0.005 to 0.5% Cr, 0.005 to 0.55% Mo, 0.005 to 0.05% Nb, 0.005 to 0.08% Ti, 0.005 to 0.2% V, 0.01 to 0.5% Al, 0.003 to 0.05% P, 0.001 to 0.01% S, 0.001 to 0.01% N, 0.0005 to 0.005% B, the balance being Fe and other inevitable impurities, and satisfying the following relations 1 and 2 to 1150 to 1350°C; Rough rolling the heated slab at a rough rolling temperature (RDT) of 880 to SCT+170°C based on 1 / 2t (t: thickness of steel material) to obtain a bar; Finishing the bar at a finish rolling temperature (FDT) of 780 to SCT+50°C based on 1 / 2t (t: thickness of steel material) to obtain a hot-rolled steel sheet; cooling the hot-rolled steel sheet from a cooling start temperature (WCT) of 700°C to SCT+10°C, and performing primary cooling at a primary average cooling rate of 50 to 100°C / sec to a primary cooling stop temperature of Ms to Ms+50°C; and Secondarily cooling the primarily cooled hot-rolled steel sheet to a coiling temperature (CT) of 70°C to Ms-50°C at a secondary average cooling rate of 1 to 40°C / sec, and then coiling the steel sheet; The surface temperature of the bar at the end of the rough rolling (RST) is 750 to RDT-40 ° C., and the surface temperature of the hot-rolled steel plate at the end of the finish rolling (FST) is controlled to be 700 to FDT-40 ° C. A method for manufacturing a hot-rolled steel plate. [Relationship 1] 0.1≦X=(Nb / 9+Ti* / 5+V / 5) / (C / 12+N / 14)≦0.6 [Relationship 2] 1.5≦T=Mn+2.8Mo+1.5Cr+500B≦5.0 [Formula 1] Ti*=Ti-3.42N-1.5S [Formula 2] SCT(℃)=741+134C-137Si+75.4Mn-21.4Cr+24.8Mo-1391Nb-13Ti+19330B [Formula 3] Ms(℃)=430-380C-13.4Si-47.3Mn-16Cr-24.2Mo (However, in the above-mentioned relational formula 1, Ti* is represented by the above-mentioned [formula 1], and the contents of each alloy composition in the above-mentioned relational formulas 1 and 2 and formulas 1 to 3 mean wt %.)

6. The method for producing a hot-rolled steel sheet according to claim 5, wherein the surface temperatures of the bar and the hot-rolled steel sheet are controlled by a water injection device.

7. The method for manufacturing a hot-rolled steel sheet according to claim 5, further comprising pickling and oiling the coiled hot-rolled steel sheet after the coiling.

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