EP4621092A1

EP4621092A1 - Hot rolled steel sheet and manufacturing method therefor

Info

Publication number: EP4621092A1
Application number: EP23892002.9A
Authority: EP
Inventors: Il-Cheol YI; Sung-Il Kim; Je-Woong LEE
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2022-11-18
Filing date: 2023-11-15
Publication date: 2025-09-24
Also published as: WO2024106936A1; EP4621092A4; JP2026508718A; KR20240075040A; CN120187885A

Abstract

The present invention relates to a hot rolled steel sheet and a manufacturing method therefor. The objective of one aspect of the present invention is to provide a hot rolled steel sheet having excellent bending workability, strength and hardness, and a manufacturing method therefor.

Description

Technical Field

The present disclosure relates to a hot-rolled steel sheet and a method for manufacturing the same.

Background Art

In the prior art, a wear-resistant hot-rolled steel sheet having high hardness has high strength and hardness by mainly using a martensite-series microstructure, and such high hardness leads to high wear resistance, so the wear-resistant hot-rolled steel sheet has been used as a component requiring wear resistance. However, it has the disadvantage of being used with only minimal processing applied because processing is limited due to poor bending workability due to the high strength. This also acts as a limiting factor similarly in a hot-rolled steel material having high strength using martensite as a main phase, and various technologies have been proposed to overcome this.
In patent document 1, by mainly controlling alloying components in steel such as C, Si, Mn, or the like, and minimizing rolling at an austenite non-recrystallization zone temperature during hot rolling to reduce an aspect ratio of old austenite particles and simultaneously suppressing a texture which strengthens anisotropy, it was intended to improve bendability.
In patent document 2, a high-strength steel sheet was comprised of tempered martensite in a central portion and ferrite and pearlite in a surface layer portion as main phases in a thickness direction of the steel sheet, so that it was intended to improve bendability.
However, Patent Document 1 discloses a technology which requires high-temperature rolling and requires control of a texture of steel, which is difficult to control, and thus has difficulties in the manufacturing process. Patent Document 1 has the disadvantage that it is difficult to be used as a wear-resistant steel since it is difficult to secure uniform hardness as a bainite phase is included in the central portion, as a main phase in addition to martensite that can secure high hardness.
In Patent Document 2, as the surface layer portion is comprised of ferrite and pearlite, which are excessively soft structures, compared to the central portion, the surface layer portion has low hardness, it has low hardness, and at the same time, deformation is concentrated in the surface layer portion during bending, so that Patent Document 2 has the disadvantage that it is unsuitable to be used as a wear-resistant steel.
In addition, alloying components such as Si, Mn, Mo, Cr, Cu, Ni, or the like which are mainly used to manufacture the steels having high hardness described above, are effective in improving hardness and formability, but if a large amount of alloying components are added to improve properties, segregation of alloying components and unevenness of microstructure occur, resulting in poor bending workability. In particular, a steel with high hardenability is sensitive to a change in microstructures when cooled, so a low-temperature transformation structure is formed unevenly, making it difficult to obtain higher bending workability.

[Prior art Document]

(Patent document 1) Japanese Patent Publication No. 2013- 117068
(Patent document 2) Korean Patent Publication No. 10-2021-0088646

Summary of Invention

Technical Problem

An aspect of the present disclosure is to provide a hot-rolled steel sheet and a method for manufacturing the same.
A preferred aspect of the present disclosure is to provide a hot-rolled steel sheet having excellent bending workability, strength, and hardness and a method for manufacturing the same.

Solution to Problem

According to an aspect of the present disclosure, provided is a hot-rolled steel sheet, the hot-rolled steel sheet including by weight%: 0.17 to 0.26% of C, 0.01 to 0.5% of Si, 0.3 to 2.0% of Mn, 0.005 to 0.5% of Cr, 0.005 to 0.55% of Mo, 0.005 to 0.05% of Nb, 0.005 to 0.08% of Ti, 0.005 to 0.2% of V, 0.01 to 0.5% of Al, 0.003 to 0.05% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.0005 to 0.005% of B, with a remainder of Fe and other inevitable impurities, and satisfying the following Relational Expressions 1 and 2, wherein a microstructure has a central portion including by area%, 90% or more of the sum of martensite and auto-tempered martensite and 10% or less of at least one of pearlite and bainite, and a surface layer portion including by area%, 90% or more of bainite and 10% or less of at least one of ferrite, martensite, and pearlite, wherein the surface layer portion has an average thickness of 30 to 200 µm, wherein the surface layer portion has an average potential density of 1.7×10¹⁴ to 3.0×10¹⁴m^-2, wherein the central portion has an aspect ratio of old austenite of 5 or more. 0.1 ≤ X = (Nb/9+Ti*/5+V/5)/(C/12+N/14) ≤ 0.6 1.5 ≤ T = Mn+2.8Mo+1.5Cr+500B ≤ 5.0 $Ti * = Ti - 3.42 N - 1.5 S$
in the Relational Expression 1, Ti* is represented by the [Equation 1], and a content of each alloy composition in the Relational Expressions 1 and 2 and Equation 1 refers to % by weight.
The central portion may have Rockwell hardness of 44 to 50 HrC.
The surface layer portion may have Rockwell hardness of 38 to 46 HrC.
The hot-rolled steel sheet may have a value of [bending workability (R/t) ÷ (tensile strength - 1000)] × 1000 of 5 or less.
According to another aspect of the present disclosure, provided is a method of manufacturing a hot-rolled steel sheet, the method including: heating a slab including by weight%, 0.17 to 0.26% of C, 0.01 to 0.5% of Si, 0.3 to 2.0% of Mn, 0.005 to 0.5% of Cr, 0.005 to 0.55% of Mo, 0.005 to 0.05% of Nb, 0.005 to 0.08% of Ti, 0.005 to 0.2% of V, 0.01 to 0.5% of Al, 0.003 to 0.05% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.0005 to 0.005% of B, with a remainder of Fe and other inevitable impurities, and satisfying the following Relational Expressions 1 and 2, to a temperature within a range of 1150 to 1350°C; completing rough rolling of the heated slab at a rough rolling temperature (RDT) of 880 to SCT+170°C based on 1/2t, where t is a thickness of a steel material, to obtain a bar; completing finishing rolling of the bar at a finishing rolling temperature (FDT) of 780 to SCT+170°C based on 1/2t, where t is a thickness of a steel material, to obtain a hot-rolled steel sheet; initiating cooling of the hot-rolled steel sheet at a cooling start temperature (WCT) of 700°C to SCT+10°C and primarily cooling the same 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 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 same, wherein a surface temperature (RST) of the bar at the end of the rough rolling is controlled to be 750 to RDT-40°C, and a surface temperature (FST) of the hot-rolled steel sheet at the end of the finishing rolling is controlled to be 700 to FDT-40°C. 0.1 ≤ X = (Nb/9+Ti*/5+V/5)/(C/12+N/14) ≤ 0.6 1.5 ≤ T = Mn+2.8Mo+1.5Cr+500B ≤ 5.0 $Ti * = Ti - 3.42 N - 1.5 S$ SCT(°C) = 741+134C-137Si+75.4Mn- 21.4Cr+24.8Mo-1391Nb-13Ti+19330B $Ms (° C) = 430 - 380 C - 13.4 Si - 47.3 Mn - 16 Cr - 24.2 Mo$
in the Relational Expression 1, Ti* is represented by the [Equation 1], and a content of each alloy composition in the Relational Expressions 1 and 2 and Equations 1 to 3 refers to % by weight.
The surface temperature of the bar and the surface temperature of the hot-rolled steel sheet may be controlled by a water spraying device.
After the coiling, pickling and oiling the coiled hot-rolled steel sheet may be further included.

Advantageous Effects of Invention

According to an aspect of the present disclosure, a hot-rolled steel sheet and a method for manufacturing the same may be provided.
According to a preferred aspect of the present disclosure, a hot-rolled steel sheet having excellent bending formability and hardness and a method for manufacturing the same may be provided.

Brief description of drawings

FIG. 1 is a graph illustrating a relationship between bending workability (R/t) according to tensile strengths of Inventive Examples 1 to 10 and Comparative Examples 1 to 12 according to an embodiment of the present disclosure.
FIG. 2 is a photograph of Inventive Example 1 according to an embodiment of the present disclosure observed using an electron microscope.
FIG. 3 is a photograph of a surface layer portion of Inventive Example 1 according to an embodiment of the present disclosure observed using an electron microscope.
FIG. 4 is a photograph of a central portion of Inventive Example 1 according to an embodiment of the present disclosure observed using an electron microscope.

Best Mode for Invention

Hereinafter, a hot-rolled steel sheet according to an embodiment of the present disclosure will be described. A content of an alloy composition described below refers to % by weight, unless otherwise specified.

Carbon (C): 0.17 to 0.26%

Carbon (C) is an element which is the most economical and effective for strengthening steel, and has a significant influence on a hardness value. As a content of added C increases, hardenability increases, making it easier to form hard phases such as bainite, martensite, or the like, in a microstructure, which increases a tensile strength. In addition, C forms fine precipitates together with Ti and Nb, which have a high affinity with C, and both a yield strength and tensile strength increase through precipitation strengthening. However, when the content of C exceeds 0.26%, there is a problem in that hardness of martensite itself increases excessively, resulting in excessive increase in strength and reduced bending workability, and it may be difficult to secure sufficient weldability. Meanwhile, when the content of C is less than 0.17%, it is difficult to obtain a sufficient strengthening effect. Therefore, it is preferable that the content of C be within the range of 0.17 to 0.26%. A lower limit of the content of C is more preferably 0.175%, even more preferably 0.18%, and most preferably 0.185%. An upper limit of the content of C is more preferably 0.25%, even more preferably 0.24%, and most preferably 0.23%.

Silicon (Si): 0.01 to 0.5%

Silicon (Si) is an element which is advantageous in deoxidizing molten steel, exhibiting a solid solution strengthening effect, and improving formability by delaying the formation of coarse carbides. When the content of Si is less than 0.01%, the effects of solid solution strengthening and the formation of bainite and martensite may not be sufficiently obtained. On the other hand, when the content of Si exceeds 0.5%, it is not easy to remove red scales formed on a surface of the steel sheet during hot rolling, and as a result thereof, the surface quality of the steel sheet may become very poor. In addition, there is a problem in that the ductility and weldability are also reduced. Therefore, it is preferable that the content of Si be within a range of 0.01 to 0.5%. A lower limit of the content of Si is more preferably 0.012%, even more preferably 0.015%, and most preferably 0.02%. An upper limit of the content of Si is more preferably 0.4%, even more preferably 0.35%, and most preferably 0.3%.

Manganese (Mn): 0.3 to 2.0%

Manganese (Mn), like Si, is an element which is effective in strengthening steel by solid solution, and increases the hardenability of steel, facilitates the formation of bainite and martensite, which are hard phases, during cooling after hot rolling. When the content of Mn is less than 0.3%, the effects of solid solution strengthening and the formation of bainite and martensite may not be sufficiently obtained. On the other hand, when the content of Mn exceeds 2.0%, grain boundaries becomes weak, which causes a problem such as low-temperature cracking, or the like. In addition, the strength may increase excessively, making it difficult to secure sufficient formability, and in a continuous casting process, a segregation zone is greatly developed in a central portion of the thickness during slab casting, and when cooling is performed after hot rolling, a microstructure is formed unevenly in the thickness direction, resulting in poor bending workability. In particular, it is also difficult to uniformly manufacture the microstructure during cooling in the overall length and overall width of the hot-rolled steel sheet. Therefore, it is preferable that the content of Mn be within the range of 0.3 to 2.0%. A lower limit of the content of Mn is more preferably 0.35%, even more preferably 0.4%, and most preferably 0.45%. An upper limit of the content of Mn is more preferably 1.9%, even more preferably 1.85%, and most preferably 1.8%.

Chromium (Cr): 0.005 to 0.5%

Chromium (Cr) strengthens steel by solid solution and delays ferrite phase transformation during cooling, thereby serving to help the formation of martensite and bainite. When the content of Cr is less than 0.005%, the effects of solid solution strengthening and the formation of martensite and bainite may not be sufficiently obtained. On the other hand, when the content of Cr exceeds 0.5%, similarly to Mn, a segregation zone in a central portion of the thickness is greatly developed and a microstructure in the thickness direction is formed unevenly, which reduces the bending workability. Therefore, it is preferable that the content of Cr be within the range of 0.005 to 0.5%. A lower limit of the content of Cr is more preferably 0.007%, even more preferably 0.008%, and most preferably 0.01%. An upper limit of the content of Cr is more preferably 0.4%, even more preferably 0.35%, and most preferably 0.3%.

Molybdenum (Mo): 0.005 to 0.55%

Molybdenum (Mo) increases the hardenability of steel to facilitate the formation of martensite and bainite. When the content of Mo is less than 0.005%, the above-described effect may not be sufficiently obtained. On the other hand, when the content of Mo exceeds 0.55%, martensite is formed in the surface layer portion due to an excessive increase in hardenability, which significantly deteriorates bending workability, is economically disadvantageous, and it may be difficult to secure sufficient weldability. Therefore, it is preferable that the content of Mo be within the range of 0.005 to 0.55%. A lower limit of the content of Mo is more preferably 0.01%, even more preferably 0.02%, and most preferably 0.03%. An upper limit of the content of Mo is more preferably 0.52%, even more preferably 0.5%, and most preferably 0.45%.

Niobium (Nb): 0.005 to 0.05%

Niobium (Nb) is a representative precipitation strengthening element, together with Ti and V, and is precipitated 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. When the content of Nb is less than 0.005%, the above-described effect may not be sufficiently obtained. On the other hand, when the content of Nb exceeds 0.05%, coarse composite precipitates are formed during hot rolling, which deteriorates bending workability. Therefore, it is preferable that the content of Nb be within the range of 0.005 to 0.05%. A lower limit of the content of Nb is more preferably 0.007%, even more preferably 0.008%, and most preferably 0.01%. An upper limit of the content of Nb is more preferably 0.04%, even more preferably 0.03%, and most preferably 0.02%.

Titanium (Ti): 0.005 to 0.08%

Titanium (Ti) is a representative precipitation strengthening element together with Nb and V, and forms coarse TiN through strong affinity with nitrogen. The TiN has the effect of suppressing grain growth during a heating process for hot rolling. In addition, Ti remaining after reacting Ni is dissolved in steel, and is combined with carbon to form TiC precipitates, which is a useful component for improving the strength of steel. When the content of Ti is less than 0.005%, the effects of suppressing grain growth and improving strength may not be sufficiently obtained. On the other hand, when the content of Ti exceeds 0.08%, coarse TiN occurs and the precipitates become coarse, which deteriorates the bending workability during forming. Therefore, it is preferable that the content of Ti be within the range of 0.005 to 0.08%. A lower limit of the content of Ti is more preferably 0.01%, even more preferably 0.015%, and most preferably 0.02%. An upper limit of the content of Ti is more preferably 0.07%, even more preferably 0.06%, and most preferably 0.045%.

Vanadium (V): 0.005 to 0.2%

Vanadium (V) is a representative precipitation strengthening element together with Nb and Ti, and V hardly precipitates during hot rolling, but V forms precipitates after high-temperature coiling, cooling, or tempering, thereby improving the strength of steel. Therefore, V is effective in improving the additional strength without increasing deformation resistance and a rolling load due to recrystallization delay during hot rolling. When the content of V is less than 0.005%, the effect of improving strength may not be sufficiently obtained. On the other hand, when the content of V exceeds 0.2%, coarse precipitates are formed, which deteriorates bending workability and is also economically disadvantageous. Therefore, it is preferable that the content of V be within the range of 0.005 to 0.2%. A lower limit of the content of V is more preferably 0.006%, even more preferably 0.008%, and most preferably 0.01%. An upper limit of the content of V is more preferably 0.2%, even more preferably 0.1%, and most preferably 0.05%.

Aluminum (Al): 0.01 to 0.5%

Aluminum (Al) is an element added mainly for deoxidation. When the content of Al is less than 0.01%, the deoxidation effect may not be sufficiently obtained. On the other hand, when the content of Al exceeds 0.5%, Al combines with nitrogen to form excessive AlN, which makes it easy for corner cracks to occur in a slab during continuous casting, and defects due to the formation of inclusions are likely to occur. Therefore, it is preferable that the content of Al be within the range of 0.01 to 0.5%. A lower limit of the content of Al is more preferably 0.015%, and even more preferably 0.02%. An upper limit of the content of Al is more preferably 0.1%, and even more preferably 0.08%, and even more preferably 0.05%.

Phosphorus (P): 0.003 to 0.05%

Phosphorus (P), like Si, has the effects of both solid solution strengthening and ferrite transformation promotion. However, in order to control the content of P to be less than 0.003%, it requires a lot of manufacturing costs, which is economically disadvantageous and is insufficient to obtain the strength. On the other hand, when the content of P exceeds 0.05%, brittleness due to grain boundary segregation may occur, microcracks are likely to occur during bending, and ductility and impact resistance properties are significantly reduced. Therefore, it is preferable that the content of P be within the range of 0.003 to 0.05%. A lower limit of the content of P is more preferably 0.005%, even more preferably 0.007%, and most preferably 0.01%. An upper limit of the content of P is more preferably 0.03%.

Sulfur (S): 0.001 to 0.01%

Sulfur (S) is an impurity present in steel, and when the content of S exceeds 0.01%, S combines with Mn, or the like to form non-metallic inclusions, and accordingly, S has the problem in that microcracks are likely to occur during bending of steel and impact resistance is significantly reduced. In the present disclosure, a lower limit of the content of S is not particularly limited, but in order to control a lower limit of the content of S to be less than 0.001%, it takes a lot of time during steelmaking, which reduces productivity, so considering the same, the lower limit of the content of S may be limited to 0.001%. Therefore, it is preferable that the content of S be within the range of 0.001 to 0.01%. The lower limit of the content of S is more preferably 0.002%. An upper limit of the content of S is more preferably 0.008%, even more preferably 0.006%, and most preferably 0.005%.

Nitrogen (N): 0.001 to 0.01%

Nitrogen (N) is a representative solid solution strengthening element together with C, and forms coarse precipitates together with Ti, Al, or the like. When the content of N is less than 0.001%, it is difficult to sufficiently obtain the effects of solid solution strengthening and precipitate formation, and in order to control the content of N to be less than 0.001%, it takes a lot of time during steelmaking, which reduces productivity. Meanwhile, although the solid solution strengthening effect of N is generally superior to that of carbon, there is a problem in that the toughness is significantly reduced when the content of N exceeds 0.01%. Therefore, it is preferable that the content of N be within the range of 0.001 to 0.01%. A lower limit of the content of N is more preferably 0.002%, and even more preferably 0.003%. An upper limit of the content of N is more preferably 0.008%, and even more preferably 0.007%, and even more preferably 0.006%.

Boron (B): 0.0005 to 0.005%

When boron (B) is present in steel, in a solid-solution state, B is mainly segregated at grain boundaries and has the effect of improving the brittleness of steel by stabilizing grain boundaries, and plays a role in suppressing the formation of coarse AlN nitrides by stabilizing solid solution N. In addition, B delays ferrite phase transformation and is effective in the formation of bainite and martensite, which are hard phases. When the content of B is less than 0.0005%, the effects of improving brittleness, suppressing the formation of coarse AlN nitrides, and forming bainite and martensite may not be sufficiently obtained. On the other hand, when the content of B exceeds 0.005%, the above-described effects no longer increase, and there is a disadvantage in that the ductility decreases and the formability deteriorates. Therefore, it is preferable that the content of B be within the range of 0.0005 to 0.005%. A lower limit of the content of B is more preferably 0.0006%, even more preferably 0.0008%, and most preferably 0.001%. An upper limit of the content of B is more preferably 0.004%, and even more preferably 0.003%.
The remaining component of the present disclosure is iron (Fe). However, since in the common manufacturing process, unintended impurities may be inevitably incorporated from raw materials or the surrounding environment, the component may not be excluded. Since these impurities are known to any person skilled in the common manufacturing process, the entire contents thereof are not particularly mentioned in the present specification.
Meanwhile, it is preferable that the hot-rolled steel sheet of the present disclosure satisfies the alloy composition described above, and satisfies the following Relational Expressions 1 and 2. 0.1 ≤ X = (Nb/9+Ti*/5+V/5)/(C/12+N/14) ≤ 0.6 $Ti * = Ti - 3.42 N - 1.5 S$
The Relational Expression 1 is intended to balance a precipitation phenomenon and hardenability to be advantageous for bendability. When the X value is less than 0.1, grain growth is facilitated during reheating, and recrystallization becomes uneven during hot rolling, so coarse grains are formed locally, and solid solution C and solid solution N become excessively high, which tends to increase a hardness value of the hard phase, which may ultimately result in poor bending workability. When the X value exceeds 0.6, the formation of precipitates increases excessively, and also, when the hot-rolled steel sheet is cooled, the solid solution C and solid solution N atoms in an untransformed phase are insufficient, making it difficult to stably form a hard phase and grain boundaries become weak, resulting in poor bending workability. Therefore, it is preferable that the X value be within the range of 0.1 to 0.6. A lower limit of the X value is more preferably 0.13, even more preferably 0.15, and most preferably 0.18. An upper limit of the X value is more preferably 0.58, even more preferably 0.56, and most preferably 0.55. 1.5 ≤ T = Mn+2.8Mo+1.5Cr+500B ≤ 5.0
The Relational Expression 2 is factorization of a combination of alloying elements that can maintain the formation of bainite and martensite, which are hard phases, in the microstructure of the steel of the present disclosure, at an appropriate level. When the T value is less than 1.5, it is difficult to obtain the desired hardness value because the hard phase is not sufficiently secured. Meanwhile, the larger the T value, the more the formation of hard phases such as bainite, martensite, and MA phase increases, and the hardness value of each of the hard phases also increases. Therefore, the larger the T value, the more advantageous it is for securing strength and hardness. However, when the T value exceeds 5.0, there is a problem that the bending workability deteriorates and the material deviation increases in the overall length and overall width of the hot-rolled steel sheet. Therefore, it is preferable that the T value be within the range of 1.5 to 5.0. A lower limit of the T value is more preferably 1.7, even more preferably 2.0, and most preferably 2.5. An upper limit of the T value is more preferably 4.9, even more preferably 4.7, and most preferably 4.5.
It is preferable that the microstructure of the hot-rolled steel sheet of the present disclosure includes a central portion including by area %, 90% or more of the sum of martensite and auto-tempered martensite and 10% or less of at least one of pearlite and bainite, and a surface layer portion including by area %, 90% or more of bainite and 10% or less of at least one of ferrite, martensite, and pearlite. When a total fraction of martensite and auto-tempered martensite in the microstructure of the central portion is less than 90% or a total fraction of at least one of pearlite and bainite exceeds 10%, there is a disadvantage in that hardness of the central portion becomes excessively low, making it difficult to obtain the desired high hardness. The total fraction of martensite and auto-tempered martensite in the microstructure of the central portion is more preferably 92% or more, even more preferably 94% or more, and most preferably 95% or more. The total fraction of at least one of ferrite, martensite, and pearlite in the microstructure of the central portion is more preferably 8% or less, even more preferably 6% or less, and most preferably 5% or less. When the fraction of bainite in the microstructure of the surface layer portion is less than 90% or the fraction of at least one of ferrite, martensite, and pearlite in the microstructure of the surface layer portion exceeds 10%, there is a disadvantage in that the bendability of the surface layer portion is reduced, thereby lowering 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 less. 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.
An average thickness of the surface layer portion is preferably 30 to 200 µm. When the average thickness of the surface layer portion is less than 30 µm, a soft surface layer portion may not be sufficiently secured, resulting in poor bending workability. In the present disclosure, an upper limit of the average thickness of the surface layer portion is not particularly limited, but it is not easy to exceed 200 µm in the manufacturing process. A 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 disclosure, the surface layer portion means a region [(from a surface of the steel sheet to a point 30 µm deep in a thickness direction) to (from a surface of the steel sheet to a point 200 µm deep in the thickness direction)], and the central portion means a region outside the surface layer portion. However, although the hot-rolled steel sheet of the present disclosure may have an oxide layer formed on the surface, the surface layer portion does not include the oxide layer, and therefore, the surface layer portion may be a region from directly beneath the oxide layer of the steel sheet to 30 to 200 µm in the thickness direction.
Meanwhile, the present inventors have confirmed that an average dislocation density, hereinafter referred to as 'Geometrical Necessary Dislocation (GND)', of the surface layer portion is an important factor in the balance between the strength and bending workability of a steel material. More specifically, the average potential density of the surface layer is preferably 1.7×10¹⁴ to 3.0×10¹⁴m^-2. When the average dislocation density of the surface layer portion is less than 1.7×10¹⁴m^-2, deformation is concentrated in the surface layer portion and the continuity is reduced, resulting in poor bending workability. When the average dislocation density of the surface layer portion exceeds 3.0×10¹⁴m^-2, the surface layer portion is not softened and becomes vulnerable to bending deformation. Therefore, it is preferable that the average potential density of the surface layer be within the range of 1.7×10¹⁴ to 3.0×10¹⁴m^-2. A lower limit of the average dislocation density of the surface layer portion is more preferably 1.8×10¹⁴m^-2, even more preferably 1.9×10¹⁴m^-2, and most preferably 2.0×10¹⁴m^-2. An upper limit of the average dislocation density of the surface layer portion is more preferably 2.9×10¹⁴m^-2, and most preferably 2.8×10¹⁴m^-2. Meanwhile, the average dislocation density may be calculated using kernel average misorientation (KAM) data measured by EBSD as expressed in the following Equation 1. For convenience, such calculation may be performed using software such as OIM analysis^™(EDAX), which analyzes the EBSD measurement results. In addition, the EBSD measurement may be performed based on a cross-section of the steel sheet, parallel to the rolling direction at 1/4 a thickness of the steel sheet. $GND (m^{- 2}) = 2 θ / ub$
in the Expression 4, θ is average misorientation (KAM values), u is a unit length (step size in the EBSD measurement), and b is a burgers vector.
It is preferable that the central portion has an aspect ratio of old austenite of 5 or more. By increasing the aspect ratio of old austenite as described above, grain refinement in a short-axis direction can be achieved, thereby improving bending workability. When the aspect ratio of old austenite of the central portion is less than 5, it may be difficult to sufficiently obtain the effect for improving bending workability. The aspect ratio of old austenite of the central portion is more preferably 6 or more, even more preferably 7 or more, and most preferably 8 or more. Meanwhile, in the present disclosure, since the larger the aspect ratio of old austenite of the central portion, the more advantageous the effect is, the upper limit thereof is not particularly limited. However, it is difficult for the aspect ratio of old austenite of the central portion to exceed 30 in the manufacturing process.
As described above, a hot-rolled steel sheet of the present disclosure provided may have Rockwell hardness of 44 to 50 HrC of the central portion, and have Rockwell hardness of 38 to 46 HrC of the surface layer portion. The hot-rolled steel sheet may have a value of [bending workability (R/t) ÷ (tensile strength - 1000)] × 1000 of 5 or less. The Rockwell hardness may be measured by using a Rockwell hardness tester (C scale) to measure the surface at five points in accordance with ASTM-E18-22 and calculate an average value thereof.
Hereinafter, a method for manufacturing a hot-rolled steel sheet according to an embodiment of the present disclosure will be described.
First, a slab satisfying the above-described alloy composition and Relational Expressions 1 and 2 is heated to a temperature within a range of 1150 to 1350°C. When the slab heating temperature is lower than 1150°C, precipitates are not sufficiently redissolved, so the formation of precipitates is reduced in a process after hot rolling, coarse TiN remains, and the slab is not sufficiently heated, making it difficult to control a temperature of a steel sheet at a constant level during hot rolling. On the other hand, when the slab heating temperature exceeds 1350°C, the strength is reduced due to abnormal grain growth of austenite crystal grains. Therefore, it is preferable that the slab heating temperature be within the range of 1150 to 1350°C. A lower limit of the slab heating temperature is more preferably 1155°C, and even more preferably 1160°C. An upper limit of the slab heating temperature is more preferably 1340°C, even more preferably 1330°C, and most preferably 1320°C.
Thereafter, the heated slab is subjected to completing rough rolling at a rough rolling temperature (RDT) of 880 to SCT+170°C based on 1/2t, where t is a thickness of a steel material, to obtain a bar. When the rough rolling temperature is lower than 880°C, a problem may occur in equipment operation due to an excessive high rolling load, and when the rough rolling temperature exceeds SCT+170°C, it is difficult to sufficiently elongate old austenite, making it difficult to improve bending workability. Therefore, it is preferable that the rough rolling temperature be within the range of 880 to SCT+150°C. A lower limit of the rough rolling temperature is more preferably 890°C, even more preferably 900°C, and most preferably 910°C. An upper limit of the rolling temperature is more preferably SCT+165°C, even more preferably SCT+160°C, and most preferably SCT+155°C. The SCT may be obtained by Equation 2, which is represented as follows. Meanwhile, the 1/2t, which is mentioned above, refers to a point of 1/2 in the thickness direction from a surface of the steel sheet. SCT(°C) = 741+134C-137Si+75.4Mn- 21.4Cr+24.8Mo-1391Nb-13Ti+19330B
In the present disclosure, it is preferable to control a surface temperature (RST) of a bar at the end of the rough rolling to be 750 to RDT-40°C. This is to control a surface temperature of the steel sheet to be lower than a temperature of a central portion thereof, so that a difference in temperature of a surface layer portion thereof during a cooling operation after the rolling process, is smaller than that in the central portion, thereby forming bainite instead of martensite. When the surface temperature of the bar at the end of the rough rolling is lower than 750°C, a problem may occur in equipment operation due to an excessively large rolling load. When the surface temperature of the bar exceeds RDT-40°C, there is a disadvantage in that bainite may not be sufficiently formed in the surface layer portion. Therefore, it is preferable that the surface temperature of the bar at the end of the rough rolling be within the range of 750 to RDT-40°C. A 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. An 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 disclosure, a method for controlling the surface temperature of the bar at the end of the rough rolling is not particularly limited, and for example, a water spraying device such as a descaler, or the like may be used.
Thereafter, the bar is subjected to completing finishing rolling at a finishing rolling temperature (FDT) of 780 to SCT+30°C based on 1/2t, where t is a thickness of a steel material, to obtain a hot-rolled steel sheet. When the finishing rolling temperature is lower than 780°C, a problem may occur in equipment operation due to an excessively large rolling load, and when the finishing rolling temperature exceeds SCT+30°C, it is difficult to sufficiently elongate old austenite, which has the disadvantage of making it difficult to improve bending workability. Therefore, it is preferable that the finishing rolling temperature be within the range of 780 to SCT+30°C. A lower limit of the finishing rolling temperature is more preferably 790°C, even more preferably 800°C, and most preferably 810°C. An upper limit of the finishing rolling temperature is more preferably SCT+25°C, even more preferably SCT+22°C, and most preferably SCT+20°C.
In the present disclosure, it is preferable to control a surface temperature (FST) of the hot-rolled steel sheet at the end of the finishing rolling to be 700 to FDT-40°C. As mentioned above, this is to control the surface temperature of the steel sheet to be lower than a temperature in the central portion of the steel sheet, so that a difference in temperature of the surface layer portion during a cooling operation after the rolling process, is smaller than that in the central portion, thereby forming bainite instead of martensite. When the surface temperature of the bar at the end of the finishing rolling is lower than 700°C, a problem may occur in equipment operation due to an excessively large rolling load, and when the surface temperature of the bar exceeds FDT-40°C, there is a disadvantage in that bainite may not be sufficiently formed in the surface layer portion. Therefore, it is preferable that the surface temperature of the bar at the end of the finishing rolling be within the range of 700 to FDT-40°C. A lower limit of the surface temperature of the bar at the end of the finishing rolling is more preferably 710°C, even more preferably 730°C, and most preferably 750°C. An upper limit of the surface temperature of the bar at the end of the finishing rolling is more preferably FDT-42°C, even more preferably FDT-43°C, and most preferably FDT-45°C. In the present disclosure, a method for controlling the surface temperature of the bar at the end of the finishing rolling is not particularly limited, and for example, a water spraying device such as a descaler, or the like may be used.
Thereafter, cooling of the hot-rolled steel sheet is initiated at a cooling start temperature (WCT) of 700°C to SCT+10°C and primarily cooled to a primary cooling stop temperature of Ms to Ms+50°C at a primary average cooling rate of 50 to 100°C/sec. When the primary cooling start temperature is lower than 700°C, a high-temperature phase such as ferrite or pearlite is formed instead of a low-temperature phase such as bainite or martensite, which has the disadvantage of significantly reducing strength and hardness. When the primary cooling start temperature exceeds SCT+10°C, the driving force for forming martensite rather than bainite in the surface layer portion is increased, so that bainite is not sufficiently formed in the surface layer portion or a thickness of the surface layer portion is insufficient. Therefore, it is preferable that the primary cooling start temperature be within the range of 700°C to SCT+10°C. A lower limit of the primary cooling start temperature is more preferably 710°C, even more preferably 730°C, and most preferably 750°C. An upper limit of the primary cooling start temperature is more preferably SCT+5°C, even more preferably SCT+2°C, and most preferably SCT°C. When the primary cooling stop temperature is lower than Ms, a high-temperature phase such as ferrite or pearlite is formed before a low-temperature phase such as martensite or bainite is formed, which has the disadvantage of significantly reducing the strength and hardness of the steel sheet or making the material uneven. When the primary cooling stop temperature exceeds Ms+50°C, martensite is not sufficiently formed in the central portion, so the hardness is decreased. Therefore, it is preferable that the primary cooling stop temperature be within the range of Ms to Ms+50°C. A lower limit of the primary cooling stop temperature is more preferably Ms+5°C, even more preferably Ms+8°C, and more preferably Ms+10°C. An 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. When 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. When the primary average cooling rate exceeds 100°C/sec, the risk of accidents such as plate entrapment within the equipment may increase due to a rapid volume change in the steel sheet caused by excessively rapid phase transformation. Therefore, it is preferable that the primary average cooling rate be within the range of 50 to 100°C/sec. A 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. An 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.
Thereafter, the primarily-cooled hot-rolled steel sheet is secondarily cooled at a second 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 why the secondary average cooling rate is lower than the primary average cooling rate is that once sufficient martensite has been formed in the central portion, further high-speed cooling will only deteriorate the shape quality of the steel sheet and will have no further advantages, so it is advantageous that high-speed cooling is switched to low-speed cooling in terms of productivity. When the secondary average cooling rate is less than 1°C/sec, there is a disadvantage that a cooling section becomes excessively long, there is a disadvantage of making equipment operation difficult, and when the secondary average cooling rate exceeds 40°C/sec, it is difficult that uniform cooling is performed. Therefore, it is preferable that the secondary average cooling rate has a range of 1 to 40°C/sec. A 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. An 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. On the other hand, a wear-resistant steel with high hardness is generally manufactured by performing a heat treatment to make the microstructure tempered martensite in order to overcome poor workability. However, when a heat treatment is performed, the surface layer portion is decarburized and the surface hardness is excessively lowered to the level of ferrite, which has the disadvantage of greatly reducing the function as a wear-resistant steel material. To prevent this, it is preferable to omit the heat treatment in the present disclosure. However, when the heat treatment is omitted, the strength of the steel sheet is excessively high and the bending workability is inferior, so it is preferable to control a coiling temperature to 70°C to Ms-50°C in order to add an auto-tempering (self-tempering) effect to the bainite of the surface layer portion and the martensite of the central portion. When the coiling temperature is lower than 70°C, the auto-tempering effect is not significant, there is a disadvantage in that it is difficult to obtain the effect of improving bending workability. When the coiling temperature exceeds Ms-50°C, a stress formed by phase transformation inside a low-temperature phase may not be sufficient, and thus high hardness may not be obtained. Therefore, it is preferable that the coiling temperature be within the range of 70°C to Ms-50°C. A lower limit of the coiling temperature is more preferably 80°C, even more preferably 90°C, and most preferably 100°C. An upper limit of the coiling temperature is more preferably Ms-60°C, even more preferably Ms-80°C, and most preferably Ms-100°C. The SCT may be obtained by Equation 3, which is expressed as follows. Ms(°C) = 430-380C-13.4Si-47.3Mn-16Cr- 24.2Mo
After the coiling, pickling and oiling the coiled hot-rolled steel sheet may be further included. In the present disclosure, the pickling and oiling processes are not specifically limited, and all methods commonly used in the relevant technical field may be used.

Mode for Invention

Hereinafter, the present disclosure will be specifically described through the following Examples. However, it should be noted that the following examples are only for describing the present disclosure by illustration, and not intended to limit the right scope of the present disclosure. The reason is that the right scope of the present disclosure is determined by the matters described in the claims and reasonably inferred therefrom.

(Example)

A slab having the alloy composition illustrated in Tables 1 and 2 below was prepared and then a hot-rolled steel sheet was manufactured under the conditions illustrated in Tables 3 and 4 below. Meanwhile, a temperature of a bar at the end of rough rolling and a temperature of the hot-rolled steel sheet at the end of finishing rolling illustrated in Tables 3 and 4 below were based on 1/2t, where t is a thickness of a steel material, and a slab heating temperature, a cooling start temperature, and a coiling temperature were based on 1/2t, where t is a thickness of a steel material.
A microstructure and mechanical properties of the hot-rolled steel sheet manufactured in this manner were measured, and the results thereof were shown in Table 5.
A type and fraction of a microstructure, an aspect ratio of old austenite of a central portion, an average thickness of a surface layer portion were measured using an electron microscope after collecting specimens from the surface layer portion and the central portion of a hot-rolled steel sheet, at 1/4 a thickness of the steel sheet. In this case, the aspect ratio of old austenite of the central portion was measured in a transverse direction (TD), which is a side surface of the steel sheet, and a horizontal direction in the image was set to be a rolling direction (RD), and a vertical direction in the image was set to be a normal direction (ND).
An average potential density (GND) was calculated using OIM analysis^™(EDAX)after measuring electron back scattered diffraction (EBSD, (JEOL JSM-7001F)) based on a cross-section of the hot-rolled steel sheet parallel to a rolling direction at the position of 1/4 the thickness of the hot-rolled steel sheet.
Hardness was measured by measuring a surface and a central portion, at 1/4 the thickness of the steel sheet, of a hot-rolled steel sheet at 5 points on the surface using a Rockwell hardness tester (C scale) in accordance with ASTM-E18-22, and an average value thereof was calculated.
Bending workability (R/t) refers to a ratio of a bending radius (R) and a thickness (t) of a steel sheet, and is represented based on a minimum bending radius at which no cracks occur on the surface after a 90-degree bending test of the steel sheet. In this case, a bending specimen was processed to be long in a direction perpendicular to a rolling direction, and a bending test was performed so that a bending line of the bending specimen was parallel to the rolling direction.

Tensile strength and yield strength were measured using a tensile tester after taking a JIS No. 5 standard specimen from the central portion, which is 1/4 of the thickness of the steel sheet.

[Table 1]

Steel type No.	Alloy composition (weight %)
Steel type No.	C	Si	Mn	Cr	Mo	Nb	Ti	V
Inventi ve Steel 1	0.2	0.3	1.5	0.2	0.1	0.02	0.03	0.01
Inventi ve Steel 2	0.19	0.2	0.9	0.3	0.3	0.015	0.02	0.02
Inventi ve Steel 3	0.22	0.03	1	0.05	0.2	0.005	0.02	0.02
Inventi ve Steel 4	0.18	0.05	1.2	0.01	0.5	0.02	0.04	0.01
Inventi ve Steel 5	0.24	0.02	0.5	0.4	0.4	0.01	0.03	0.02
Inventi ve Steel 6	0.25	0.02	0.8	0.2	0.2	0.01	0.03	0.01
Inventi ve Steel 7	0.23	0.02	1.8	0.1	0.03	0.01	0.025	0.02
Inventi ve Steel 8	0.21	0.05	1.3	0.1	0.5	0.01	0.025	0.01
Inventi ve Steel 9	0.25	0.1	1.2	0.3	0.3	0.01	0.045	0.03
Inventi ve Steel 10	0.21	0.01	1.35	0.01	0.22	0.005	0.03	0.01
Compara tive Example 1	0.22	0.03	1	0.005	0.2	0.1	0.02	0.01
Compara tive Example 2	0.24	0.2	2.3	0.7	0.6	0.005	0.04	0.01
Compara tive Example 3	0.21	0.05	1.2	0.005	0.1	0.01	0.01	0.005
Compara tive Example 4	0.16	0.05	0.5	0.005	0.1	0.01	0.03	0.01
Inventi ve Steel 11	0.21	0.05	1.2	0.1	0.5	0.01	0.025	0.01
Inventi ve Steel 12	0.21	0.02	1.7	0.4	0.4	0.015	0.02	0.02
Inventi ve Steel 13	0.25	0.03	0.8	0.25	0.2	0.018	0.02	0.01
Inventi ve Steel 14	0.2	0.04	1.4	0.01	0.3	0.011	0.02	0.02
Compara tive Example 5	0.28	0.02	1.8	0.5	0.1	0.015	0.03	0.01
Compara tive Example 6	0.3	0.06	1	0.4	0.4	0.02	0.04	0.03
Inventi ve Steel 15	0.19	0.1	1.6	0.3	0.1	0.016	0.02	0.01
Inventi ve Steel 16	0.19	0.2	1.7	0.2	0.4	0.016	0.03	0.01

[Table 2]

Steel type No.	Alloy composition (weight %)
Steel type No.	Al	P	S	N	B	X	T
Inventive Steel 1	0.03	0.01	0.002	0.004	0.0015	0.41	2.83
Inventive Steel 2	0.03	0.01	0.002	0.004	0.002	0.39	3.19
Inventive Steel 3	0.03	0.01	0.002	0.005	0.002	0.24	2.64
Inventive Steel 4	0.03	0.01	0.005	0.006	0.002	0.43	3.62
Inventive Steel 5	0.02	0.02	0.003	0.004	0.002	0.37	3.22
Inventive Steel 6	0.02	0.005	0.003	0.003	0.002	0.29	2.66
Inventive Steel 7	0.03	0.015	0.004	0.005	0.0025	0.28	3.28
Inventive Steel 8	0.03	0.015	0.004	0.005	0.003	0.20	4.35
Inventive Steel 9	0.02	0.01	0.004	0.005	0.002	0.54	3.49
Inventive Steel 10	0.04	0.01	0.002	0.005	0.002	0.25	2.98
Comparati ve Example 1	0.03	0.01	0.002	0.005	0.002	0.70	2.57
Comparati ve Example 2	0.03	0.015	0.005	0.006	0.002	0.24	6.03
Comparati ve Example 3	0.03	0.01	0.008	0.006	0.002	-0.13	2.49
Comparati ve Example 4	0.04	0.01	0.004	0.006	0.001	0.28	1.29
Inventive Steel 11	0.04	0.015	0.003	0.004	0.0025	0.25	4.00
Inventive Steel 12	0.04	0.015	0.002	0.004	0.001	0.36	3.92
Inventive Steel 13	0.04	0.01	0.003	0.003	0.0028	0.24	3.14
Inventive Steel 14	0.02	0.01	0.001	0.003	0.001	0.41	2.76
Comparati ve Example 5	0.03	0.015	0.001	0.005	0.0025	0.25	4.08
Comparati ve Example 6	0.03	0.02	0.002	0.005	0.0015	0.48	3.47
Inventive Steel 15	0.02	0.01	0.003	0.004	0.002	0.26	3.33
Inventive Steel 16	0.04	0.015	0.005	0.003	0.003	0.39	4.62
X = (Nb/9+Ti*/5+V/5)/(C/12+N/14)
Ti* = Ti-3.42N-1.5S
T = Mn+2.8Mo+1.5Cr+500B

[Table 3]

Division	Steel type No.	Slab heating temperatu re (°C)	Temperatu re of bar at the end of rough rolling	Surface temperatu re of bar at the end of rough rolling (RST) (°C)	Temperatu re of hot-rolled steel sheet at the end of finishing rolling (FDT) (°C)	Surface temperatu re of hot-rolled steel sheet at the end of finishing rolling (F ST) (°C)	Cooling start temperatu re (WCT) (°C)
Inventive Example 1	Inventive Steel 1	1195	989	923	846	780	820
Inventive Example 2	Inventive Steel 2	1160	957	883	846	778	814
Inventive Example 3	Inventive Steel 3	1220	1009	955	879	826	857
Inventive Example 4	Inventive Steel 4	1215	1002	927	891	817	871
Inventive Example 5	Inventive Steel 5	1170	966	897	845	779	824
Inventive Example 6	Inventive Steel 6	1210	991	919	867	800	840
Inventive Example 7	Inventive Steel 7	1290	1087	1015	939	874	934
Inventive Example 8	Inventive Steel 8	1270	1049	970	944	865	911
Inventive Example 9	Inventive Steel 9	1220	1005	954	877	828	864
Inventive Example 10	Inventive Steel 10	1270	1049	986	917	859	894
Comparati ve Example 1	Comparati ve Steel 1	1160	891	828	782	729	740
Comparati ve Example 2	Comparati ve Steel 2	1290	1076	1020	964	909	950
Comparati ve Example 3	Comparati ve Steel 3	1230	1025	946	910	834	870
Comparati ve Example 4	Comparati ve Steel 4	1160	934	865	829	769	793
Comparati ve Example 5	Inventive Steel 11	1250	1041	988	911	860	905
Comparati ve Example 6	Inventive Steel 12	1230	1018	954	983	920	890
Comparati ve Example 7	Inventive Steel 13	1230	1008	988	875	850	849
Comparati ve Example 8	Inventive Steel 14	1230	1012	990	894	872	890
Comparati ve Example 9	Comparati ve Steel 5	1270	1056	996	956	903	927
Comparati ve Example 10	Comparati ve Steel 6	1210	997	921	864	795	850
Comparati ve Example 11	Inventive Steel 15	1220	1023	964	891	836	872
Comparati ve Example 12	Inventive Steel 16	1260	1042	976	938	882	899

[Table 4]

Division	Steel type No.	Primary cooling rate (°C/s)	Primary cooling stop temperatu re (°C)	Secondary cooling rate (°C/s)	Coiling temperatu re (CT) (°C)	SCT (°C)	Ms (°C)
Inventive Example 1	Inventive Steel 1	75	286	28	146	839	273
Inventive Example 2	Inventive Steel 2	75	314	29	129	825	300
Inventive Example 3	Inventive Steel 3	65	320	18	141	877	293
Inventive Example 4	Inventive Steel 4	70	320	30	115	871	292
Inventive Example 5	Inventive Steel 5	65	318	26	100	834	299
Inventive Example 6	Inventive Steel 6	75	305	16	144	857	289
Inventive Example 7	Inventive Steel 7	65	266	15	101	937	255
Inventive Example 8	Inventive Steel 8	85	292	19	114	914	274
Inventive Example 9	Inventive Steel 9	80	288	15	126	876	265
Inventive Example 10	Inventive Steel 10	85	306	15	108	906	281
Comparati ve Example 1	Comparati ve Steel 1	85	319	28	125	746	294
Comparati ve Example 2	Comparati ve Steel 2	85	225	23	114	950	202
Comparati ve Example 3	Comparati ve Steel 3	80	301	24	100	880	290
Comparati ve Example 4	Comparati ve Steel 4	85	362	22	150	801	342
Comparati ve Example 5	Inventive Steel 11	80	298	17	113	897	279
Comparati ve Example 6	Inventive Steel 12	70	276	23	140	894	253
Comparati ve Example 7	Inventive Steel 13	65	316	20	135	859	288
Comparati ve Example 8	Inventive Steel 14	85	304	23	133	879	280
Comparati ve Example 9	Comparati ve Steel 5	85	254	22	110	930	228
Comparati ve Example 10	Comparati ve Steel 6	75	273	23	120	850	252
Comparati ve Example 11	Inventive Steel 15	85	293	32	20	886	274
Comparati ve Example 12	Inventive Steel 16	70	274	27	320	908	262
SCT(°C) = 741+134C-137Si+75.4Mn-21.4Cr+24.8Mo-1391Nb-13Ti+19330B
Ms(°C) = 430-380C-13.4Si-47.3Mn-16Cr-24.2Mo

[Table 5]

Divisio n	Microstructure							Mechanical properties
	Central portion			Surface layer portion				Hard ness of cent ral port ion (HrC )	Hard ness of surf ace laye r port ion( HrC)	Bend ing work abil ity (R/t )	Tens ile stre ngth (MPa )	Yiel d stre ngth (MPa )
	M+ATM (area %)	At least one of P and B (area %)	AR	B (area %)	At least one of F, M, and P (area % )	GND (¹⁴m^-2)	Thic knes s (µm)	Hard ness of cent ral port ion (HrC )	Hard ness of surf ace laye r port ion( HrC)	Bend ing work abil ity (R/t )	Tens ile stre ngth (MPa )	Yiel d stre ngth (MPa )
Inventi ve Example 1	100	0	10	98	2	2.3	58	45.9	41.9	1.7	1435	1182
Inventi ve Example 2	100	0	15	100	0	2.1	53	44.8	38.8	1.7	1394	1141
Inventi ve Example 3	98	2	9	100	0	2.4	48	46.5	40.5	2.3	1512	1264
Inventi ve Example 4	100	0	12	100	0	2.2	44	44	39	1	1363	1108
Inventi ve Example 5	98	2	15	98	2	2.8	52	47.7	43.7	2.3	1547	1286
Inventi ve Example 6	95	5	12	100	0	2.7	57	47.6	42.6	2	1588	1346
Inventi ve Example 7	95	5	8	98	2	2.8	44	47.7	43.7	2.3	1590	1348
Inventi ve Example 8	98	2	8	98	2	2.6	58	46.2	42.2	2	1484	1235
Inventi ve Example 9	95	5	10	95	5	2.8	43	47.1	44.1	2	1611	1371
Inventi ve Example 10	98	2	12	100	0	2.4	49	46	40	2	1490	1241
Compara tive Example 1	85	15	17	100	0	1.5	47	42.7	37	2.3	1311	994
Compara tive Example 2	96	4	8	5	95	3.2	0	47.9	47	3.6	1601	1359
Compara tive Example 3	100	0	8	85	15	3.1	18	46.7	42	4.3	1580	1310
Compara tive Example 4	75	25	8	95	5	1.4	15	39	37.5	1.6	1198	984
Compara tive Example 5	95	5	7	90	10	2.4	20	45.1	42	3.6	1464	1227
Compara tive Example 6	95	5	3	100	0	2.6	49	46	40	3.3	1518	1270
Compara tive Example 7	100	0	12	90	10	3	15	49.7	46.5	3.3	1592	1351
Compara tive Example 8	98	2	8	90	10	2.4	8	45.4	42.6	3.6	1454	1203
Compara tive Example 9	100	0	7	95	5	3.6	34	51.4	48.9	5.6	1762	1536
Compara tive Example 10	100	0	10	95	5	3.3	49	50.7	47.5	4.3	1753	1525
Compara tive Example 11	100	0	9	95	5	2.3	47	45.9	43.8	3.6	1435	1184
Compara tive Example 12	80	20	9	100	0	1.4	44	41.2	37.2	3	1167	980
M: Martensite, ATM: Auto-tempered Martensite, P: Pearlite, B: Bainite, F: Ferrite
AR: Aspect ratio of old austenite, GND: Average potential density

As can be seen from Tables 1 to 5 above, Inventive Examples 1 to 10, which satisfy the alloy composition and manufacturing conditions proposed by the present disclosure, satisfy the microstructure conditions of the present disclosure and have excellent mechanical properties.
Comparative Example 1 showed that an X value of Relational Expression 1 was excessively high, so that a driving force for forming precipitates was excessive, and accordingly, carbon and nitrogen did not sufficiently play a role in improving strength within martensite and bainite, resulting in insufficient hardness. In addition, it can be seen that GND falls short of the scope of the present disclosure.
Comparative Example 2 showed that a T value of Relational Expression 1 was excessively high, which failed to properly form bainite or martensite, so that an entire steel sheet was comprised of almost martensite without a surface layer portion, resulting in high surface hardness and poor bending workability. In addition, it was confirmed that GND exceeded the scope of the present disclosure.
Comparative Example 3 showed that an X value of Relational Expression 1 was excessively low, so there was almost no driving force for forming precipitates, and accordingly, the role of carbon and nitrogen in improving strength within martensite and bainite was excessive, resulting in poor bending workability. In addition, it was confirmed that GND exceeded the scope of the present disclosure.
Comparative Example 4 showed that a T value of Relational Expression 1 was excessively low, which failed to properly form bainite or martensite, resulting in excessive low hardness. In addition, it can be seen that the GND falls short of the range of the present disclosure.
Comparative Example 5 showed that the alloy composition of the present disclosure is satisfied but a cooling start temperature is high, a surface layer portion was cooled rapidly and a driving force to form martensite than bainite was strong so that a thickness of the surface layer portion was insufficient. As a result, it was shown that the bending workability was poor.
Comparative Example 6 showed that the alloy composition of the present disclosure is satisfied, but a finishing rolling temperature was high, so an aspect ratio of old austenite of martensite of the central portion is as small as 3, and as a result, it can be seen that the bending workability is inferior to the strength.
Comparative Examples 7 and 8 showed that the alloy composition of the present disclosure is satisfied, but a difference in temperatures between RDT and RST or FDT and FST was not sufficiently obtained, so a driving force to form bainite in the surface layer portion is low, and a surface layer having a sufficient thickness could not be obtained. As a result, it can be seen that the bending workability was inferior.
Comparative Examples 9 and 10 showed that the manufacturing conditions of the present disclosure are all satisfied due to excessive high content of carbon, but it can seen that both the surface layer portion and the central portion have excessively high hardness, and as a result, the bending workability is inferior. In addition, it was confirmed that GND exceeds the range of the present disclosure.
Comparative Example 11 showed that the alloy composition of the present disclosure is satisfied, but a coiling temperature is excessively low, so the auto-tempering effect is not obtained, and thus it can be seen that bending workability is inferior to the strength.
Comparative Example 12 showed that the alloy composition of the present disclosure is satisfied, but has excessively high coiling temperature so has a weak driving force to form martensite and bainite to have low hardness and strength. In addition, it can be seen that GND falls short of the range of the present disclosure.
FIG. 1 is a graph illustrating a relationship between bending workability (R/t) according to tensile strength of Inventive Examples 1 to 10 and Comparative Examples 1 to 12. As shown in FIG. 1, in the case of Comparative Examples 1 to 12, it can be seen that the bending workability (R/t) tends to deteriorate overall as the tensile strength increases. In the case of Inventive Examples 1 to 10, the bending workability is relatively good compared to the tensile strength, so that a value of [bending workability (R/t) ÷ (tensile strength - 1000)] × 1000 is 5 or less, and therefore, it can be confirmed that the bending workability is excellent.
FIG. 2 is a photograph of Inventive Example 1 observed using an electron microscope. FIG. 3 is a photograph of a surface layer portion of Inventive Example 1 observed using an electron microscope. FIG. 4 is a photograph of a central portion of Inventive Example 1 observed using an electron microscope. As can be seen from FIGS. 2 to 4, in the case of Inventive Example 1, it could be confirmed that the microstructures of the surface layer portion and the central portion that the present disclosure intends to be obtained are formed.

Claims

A hot-rolled steel sheet, comprising by weight%: 0.17 to 0.26% of C, 0.01 to 0.5% of Si, 0.3 to 2.0% of Mn, 0.005 to 0.5% of Cr, 0.005 to 0.55% of Mo, 0.005 to 0.05% of Nb, 0.005 to 0.08% of Ti, 0.005 to 0.2% of V, 0.01 to 0.5% of Al, 0.003 to 0.05% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.0005 to 0.005% of B, with a remainder of Fe and other inevitable impurities, and satisfying the following Relational Expressions 1 and 2,
wherein a microstructure has a central portion including by area%, 90% or more of the sum of martensite and auto-tempered martensite and 10% or less of at least one of pearlite and bainite, and a surface layer portion including by area%, 90% or more of bainite and 10% or less of at least one of ferrite, martensite, and pearlite,

wherein the surface layer portion has an average thickness of 30 to 200 µm,

wherein the surface layer portion has an average potential density of 1.7×10¹⁴ to 3.0×10¹⁴m^-2,

wherein the central portion has an aspect ratio of old austenite of 5 or more, 0.1 ≤ X = (Nb/9+Ti*/5+V/5)/(C/12+N/14) ≤ 0.6 1.5 ≤ T = Mn+2.8Mo+1.5Cr+500B ≤ 5.0 $Ti * = Ti - 3.42 N - 1.5 S$

in the Relational Expression 1, Ti* is represented by the [Equation 1], and a content of each alloy composition in the Relational Expressions 1 and 2 and Equation 1 refers to % by weight.
The hot-rolled steel sheet of claim 1, wherein the central portion has Rockwell hardness of 44 to 50HrC.
The hot-rolled steel sheet of claim 1, wherein the surface layer portion has Rockwell hardness of 38 to 46HrC.
The hot-rolled steel sheet of claim 1, wherein the hot-rolled steel sheet has a value of [bending workability (R/t) ÷ (tensile strength - 1000)] × 1000 of 5 or less.
A method for manufacturing a hot-rolled steel sheet, comprising operations of:
heating a slab including by weight%, 0.17 to 0.26% of C, 0.01 to 0.5% of Si, 0.3 to 2.0% of Mn, 0.005 to 0.5% of Cr, 0.005 to 0.55% of Mo, 0.005 to 0.05% of Nb, 0.005 to 0.08% of Ti, 0.005 to 0.2% of V, 0.01 to 0.5% of Al, 0.003 to 0.05% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.0005 to 0.005% of B, with a remainder of Fe and other inevitable impurities, and satisfying the following Relational Expressions 1 and 2, to a temperature within a range of 1150 to 1350°C;

completing rough rolling at a rough rolling temperature (RDT) of 880 to SCT+170°C based on 1/2t of the heated slab, where t is a thickness of a steel material, to obtain a bar;

completing finishing rolling at a finishing rolling temperature (FDT) of 780 to SCT+50°C based on 1/2t of the bar, where t is a thickness of a steel material, to obtain a hot-rolled steel sheet;

initiating cooling of the hot-rolled steel sheet at a cooling start temperature (WCT) of 700°C to SCT+10°C and primarily cooling the same 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

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 same,

wherein a surface temperature (RST) of the bar at the end of the rough rolling is controlled to be 750 to RDT-40°C, and a surface temperature (FST) of the hot-rolled steel sheet at the end of the finishing rolling is controlled to be 700 to FDT-40°C, 0.1 ≤ X = (Nb/9+Ti*/5+V/5)/(C/12+N/14) ≤ 0.6 1.5 ≤ T = Mn+2.8Mo+1.5Cr+500B ≤ 5.0 $Ti * = Ti - 3.42 N - 1.5 S$ SCT(°C) = 741+134C-137Si+75.4Mn- 21.4Cr+24.8Mo-1391Nb-13Ti+19330B $Ms (° C) = 430 - 380 C - 13.4 Si - 47.3 Mn - 16 Cr - 24.2 Mo$

in the Relational Expression 1, Ti* is represented by the [Equation 1], and a content of each alloy composition in the Relational Expressions 1 and 2 and Equations 1 to 3 refers to % by weight.
The method for manufacturing a hot-rolled steel sheet of claim 5, wherein the surface temperature of the bar and the surface temperature of the hot-rolled steel sheet are controlled by a water spray device.
The method for manufacturing a hot-rolled steel sheet of claim 5, further comprising:
after the coiling, pickling and oiling the coiled hot-rolled steel sheet.