EP4640886A1

EP4640886A1 - Hot-rolled steel sheet and manufacturing method therefor

Info

Publication number: EP4640886A1
Application number: EP23907639.1A
Authority: EP
Inventors: Je-Woong LEE; Sung-Il Kim; Il-Cheol YI; Hyun-taek NA
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2022-12-21
Filing date: 2023-12-15
Publication date: 2025-10-29
Also published as: JP2025541892A; MX2025007283A; EP4640886A4; CN120390819A; KR20240098898A; WO2024136354A1

Abstract

A hot-rolled steel sheet and a manufacturing method therefor are provided. The hot-rolled steel sheet of the present invention includes, by wt%, 0.020-0.080% of C, 0.01-0.50% of Si, 0.8-1.8% of Mn, 0.010-0.100% of Al, 0.001-0.020% of P, 0.001-0.010% of S, 0.001-0.010% of N, 0.010-0.120% of Ti, 0.010-0.050% of Nb, and the remainder of Fe and inevitable impurities, satisfies an X value of 3.50-6.00 as defined by equation 1, and has a micro-structure in which the sum of ferrite and bainite phases is at least 90 area% and the sum of the remaining martensite and MA phases is less than 10 area%.

Description

Technical Field

The present disclosure relates to a high-strength composite structure hot-rolled steel sheet having excellent thermal stability, which may be applicable to chassis parts of an automobile, and a method for manufacturing the same.

Background Art

A general high-strength hot-rolled steel sheet used for automobile chassis and frames has been designed to have high-strength and a reduced weight. Also, excellent formability may be required considering the shape of parts, and a specific level of bake hardening amount, which indicates the degree of hardening after painting, may be required to maximize durability of the parts. Also, there may be cases in which heat may be applied to a portion or the entirety of a steel sheet and parts for various purposes during the manufacturing process and use, and strength of the steel sheet and parts may change due to the heating process, such that durability may be reduced.
Generally, during the heating process, as the amount of solid-solute carbon in the structure increases, clustering may be formed at dislocations, grain boundaries, or the like, and carbides may be formed. Also, the microstructure of the steel, such as martensite, bainite, and retained austenite, may also change, such that strength of the steel may rapidly change, which may affect formability and durability. The changes in the microstructure and properties of steel during the heating process may vary depending on an initial steel composition and a microstructure and may be largely dependent on heat treatment conditions such as heating temperature and holding time, such that the existing technologies have focused simply on suppressing a decrease in strength at high temperatures of 600°C or higher.
Cited documents 1 and 2 propose a technology to ensure high-temperature strength by adding Cr, Mo, Nb, V, or the like, and heat-treating the steel sheet after hot rolling, but this is a technology suitable for the process of manufacturing a thick steel sheet for construction. Also, considering environmental factors such as fires inevitably heating steel for construction, alloy components such as Cr, Mo, Nb, and V may be added to the steel such that a certain level of strength may be ensured even when exposed to a high-temperature environment of 600°C or higher for a long time, but there may be a problem in that the manufacturing costs may be excessive because expensive alloy components may be used and a heat treatment process may be required to secure properties at the steel sheet stage. In particular, the steel sheet may have excessive thermal stability for use in the case in which the steel sheet is exposed to a heating environment of 600°C or lower for a short period of time.
Cited document 3 proposes a technology for securing strength of the welded heat-affected zone by adding Ti, Nb, Cr, Mo, or the like. During arc welding, the area adjacent to the welding material melted by welding heat may be heated to a high temperature of 600°C or higher, and in particular, there may be a case in which the area is heated to a temperature higher than an austenite region. Accordingly, addition of Cr and Mo may increase hardenability of the steel, and when cooling, low-temperature phases such as bainite and martensite phases may be formed, thereby securing strength. However, this concept based on maximizing hardenability may have limitations in applying to an automotive steel sheet requiring high formability even after necessary heat treatment after manufacturing the steel sheet.

[Prior art]

[Patent Reference]

(Cited document 1) Korean Registered Patent Publication No. 10-0358939
(Cited document 2) Korean Registered Patent Publication No. 10-1290382
(Cited document 3) Korean Registered Patent Publication No. 10-0962745

Detailed description of present disclosure

Technical problems to solve

An aspect in the present disclosure is to provide a high-strength composite structure hot-rolled steel sheet having excellent bake hardening amount and thermal stability and a method for manufacturing the same.
The technical problems to be achieved in the present disclosure are not limited to the technical problems mentioned above, and other technical problems not mentioned herein will be clearly understood by a person having ordinary skill in the technical field to which the present disclosure belongs from the description below.

Solution to Problem

An aspect in the present disclosure provides a hot-rolled steel sheet including

by weight%, C:0.020-0.080%, Si:0.01-0.50%, Mn:0.8-1.8%, Al:0.010-0.100%, P:0.001-0.020%, S:0.001-0.010%, N:0.001-0.010%, Ti: 0.010-0.120%, Nb:0.010-0.050%, and a balance of Fe and inevitable impurities, wherein an X value defined by relational expression 1 as below satisfies 3.50-6.00, and
having a microstructure including a sum of ferrite and bainite phases: 90 area% or more, and a sum of remainder of martensite and MA phases: less than 10 area%,
wherein the steel sheet has tensile strength of 590 MPa or more, hole expandability (HER₀) of 40% or more, and a bake hardening amount (BH₂): bake hardening amount before heat treatment) of 30 MPa or more, the steel sheet maintaining a bake hardening amount (BH_h) after heat treatment at 300-600°C of 30 MPa or more, and when ΔTS is defined as in relational expression 2 below, an absolute value of ΔTS × BH_h ^-1 is 0.70 or less:

$\begin{array}{l} X = A * B \\ A = 1.3 * [Mn] + 200 * [C] \\ B = (Nb / 93 + Ti * / 48) / (C / 12 + N / 14) \\ Ti * = Ti - 3.42 N - 1.5 S \end{array}$
where Mn, C, Nb, Ti, N and S represent weight% thereof

$ΔTS = {TS}_{h} - {TS}_{0}$
TS_h : tensile strength after heat treatment, TS₀ : tensile strength before heat treatment
BH_h : bake hardening amount after heat treatment
the hot-rolled steel sheet includes one or more of Mo, Cr, V, Ni, and B in a total amount of 1.500% or less.

Hot dip galvanized plating is formed on a surface of the hot-rolled steel sheet.
Another aspect in the present disclosure provides a method for manufacturing the hot-rolled steel sheet including

reheating a steel slab having the alloy composition described above and of which an X value defined by relational expression 1 as below satisfies 3.50-6.00 in a temperature range of 1100-1350°C;
manufacturing a hot-rolled steel sheet by hot-rolling the reheated steel slab in a range of 850-1150°C; and
cooling the hot-rolled steel sheet to a temperature in a range of 400-550°C at an average cooling rate of 10-70°C/sec, and coiling the steel sheet.

The method may further include pickling and oiling the coiled hot-rolled steel sheet.
The method may further include pickling the coiled hot-rolled steel sheet, heating the steel sheet to a temperature range of 450-750°C, and immersing the steel sheet in a plating bath containing 0.01-30% of Mg, 0.01-50% of Al, and the remainder of zinc in weight%, thereby forming a hot-dip galvanizing layer on the surface thereof.

Advantageous Effects of Invention

According to an aspect in the present disclosure, a hot-rolled steel sheet having excellent thermal stability, which may have tensile strength of 590 MPa or more, hole expandability (HER₀) of 40% or more, a bake hardening amount (BH₂) of 30 MPa or more, a bake hardening amount (BH_h) after heat treatment at 300-600°C of 30 MPa or more, and an absolute value of ΔTS × BH_h ^-1 of 0.70 or less, may be effectively provided.
Accordingly, the hot-rolled steel sheet may be effectively applied to members of automobile chassis parts and parts used in a lower arm, a reinforcement, a connecting material, and a frame.

Brief Description of Drawings

FIG. 1 is an image indicating changes (absolute value of ΔTS × BH_h ^-1) in tensile strength and bake hardening amount before and after heat treatment for X values of steel sheets of inventive examples and comparative examples.

Best Mode for Invention

Hereinafter, the present disclosure will be described.
The inventors of the present disclosure studied various steels with various components and different microstructures to expand applicability of hot-rolled chassis parts, and it was confirmed that changes in the room-temperature tensile strength measured after heat treatment in the temperature range of 100-600°C depended on the slope of the dynamic strength value measured during heating of the steel. From the results, relational expression 1 to determine the component content of C, Mn, Si, Ti, and Nb, which are major components of steel, was derived, and the optimal steel microstructure was formed based on relational expression 1, such that a high-strength composite structure hot-rolled steel sheet having tensile strength of 590MPa or more, a yield ratio of 0.7 or more, hole expandability (HER₀) value of 40% or more, and a bake hardening amount (BH₂) of 30MPa or more was manufactured, and the present disclosure was suggested. The hot-rolled steel sheet manufactured by this technology may maintain a bake hardening amount (BH_h) of 30 MPa or more after heat treatment at 100-600°C, and may have excellent thermal stability with a ΔTS × BH_h ^-1 absolute value of 0.70 or less. Accordingly, when used as an actual component, heat treatment may be performed at a relatively low temperature for a short period of time, such that application may be expanded, and may be easily used when manufacturing a plated hot-rolled steel sheet using hot dip galvanizing, or the like.
The hot-rolled steel sheet of the present disclosure may include, by weight%, C:0.020-0.080%, Si:0.01-0.50%, Mn:0.8-1.8%, Al:0.010-0.100%, P:0.001-0.020%, S:0.001-0.010%, N:0.001-0.010%, Ti: 0.010-0.120%, Nb:0.010-0.050%, and a balance of Fe and inevitable impurities, wherein an X value defined by relational expression 1 as below satisfies 3.50-6.00, and having a microstructure including a sum of ferrite and bainite phases: 90 area% or more, and a sum of remainder of martensite and MA phases: less than 10 area%, wherein the steel sheet may have tensile strength of 590 MPa or more, hole expandability (HER₀) of 40% or more, and a bake hardening amount (BH₂) : bake hardening amount before heat treatment) of 30 MPa or more, the steel sheet maintaining a bake hardening amount (BH_h) after heat treatment at 300-600°C of 30 MPa or more, and when ΔTS is defined as in relational expression 2 below, an absolute value of ΔTS × BH_h ^-1 may satisfy 0.70 or less.
In the description below, the composition of the steel sheet provided in the present disclosure may be described in detail. In this case, unless otherwise indicated, the content of each component may indicate weight%.

C: 0.020-0.080%

C may be economical and effective for strengthening steel, and when the added amount increases, the precipitation strengthening effect or the low-temperature phase fraction may increase, such that tensile strength may increase. However, when the content is less than 0.020%, sufficient precipitation strengthening effect and low-temperature phase formation may be difficult, such that it may be difficult to ensure target strength and bake hardening amount, and when the content exceeds 0.080%, excessive low-temperature phase and carbide may be formed, which may deteriorate formability and weldability. Also, after the excessive content of C is added, the low-temperature phase may deteriorate and additional excess carbide may be formed during heat treatment in the range of 100-600°C, which may significantly decrease strength and bake hardening amount after heat treatment, and may further deteriorate formability. Accordingly, in the present disclosure, it may be preferable to limit the C content to 0.020-0.080%. More preferably, the content may be limited to the range of 0.030-0.072%.

Si: 0.01-0.50%

Si may deoxidize molten steel and may have a solid-solution strengthening effect, and may be advantageous in improving formability by delaying the formation of coarse carbide. Si may also have an effect of suppressing the formation of carbide during heat treatment in the range of 100-600°C. However, when the content is less than 0.01%, the effect of delaying carbide formation may be insignificant, such that it may be difficult to improve formability and thermal stability may be reduced. When the content exceeds 0.50%, red scales may be formed on the surface of the steel sheet by Si during hot-rolling, which may significantly deteriorate the surface quality of the steel sheet, and may also decrease ductility and weldability. Accordingly, in the present disclosure, it may be preferable to limit the Si content to the range of 0.01-0.50%. More preferably, the content may be limited to the range of 0.10-0.43%.

Mn: 0.8-1.8%

Similarly to Si, Mn may be effective for solid-solution strengthening of steel and may increase hardenability of steel, such that the formation of low-temperature phase may be facilitated. However, when the content is less than 0.8%, the effect of addition may not be obtained, and when the content exceeds 1.8%, hardenability may increase significantly, such that the martensite phase fraction may increase, and the segregation zone may be greatly developed in the thickness center of slab during slab casting in the casting process, which may deteriorate formability. Also, carbide may be easily formed during heat treatment in the range of 100-600°C, such that there may be significant changes in strength and bake hardening amount. Accordingly, it may be preferable to limit the content of Mn to 0.8-1.8% in the present disclosure. More preferably, the content may be limited to the range of 0.8-1.5%.

P: 0.001-0.020%

Similarly to Si, P may have both solid-solution strengthening and ferrite transformation promotion effects. However, manufacturing steel with less than 0.001% of P may increase manufacturing costs, which may be economically disadvantageous, and may be insufficient to obtain strength. When the content exceeds 0.020%, brittleness may occur due to grain boundary segregation, and microcracks may be easily created during forming, which significantly deteriorate ductility and impact resistance properties. Accordingly, it may be preferable to limit the P content to the range of 0.001-0.02%.

S: 0.001-0.010%

S may be impurities in steel, and when the content thereof exceeds 0.010%, S may combine with Mn, or the like, and may form a non-metallic inclusion, and accordingly, fine cracks may be easily created during cutting processing of steel. When the content is less than 0.001%, it may take a great deal of time during steelmaking, which may reduce productivity. Accordingly, in the present disclosure, it may be preferable to limit the S content to 0.001-0.010%.

Sol.Al: 0.010-0.100%

Sol.Al may be mainly added for deoxidation, and when the content thereof is less than 0.010%, the addition effect thereof may be insufficient, and when the content exceeds 0.100%, Sol.Al may combine with nitrogen such that AlN may be formed, and corner cracks may be easily created in the slab during continuous casting and casting, and defects may occur due to the formation of inclusion. Accordingly, in the present disclosure, it may be preferable to limit the Sol.Al content to 0.010-0.100%.

N: 0.001-0.010%

N may be a representative solid-solution strengthening element along with C, and may form coarse precipitates along with Ti, Al, or the like. Generally, the solid-solution strengthening effect of N may be more excellent than that of carbon, but toughness may decrease significantly as the amount of N in the steel increases. Also, in order to manufacture steel with less than 0.001% of N, it may take a great deal of time during steelmaking operation, which may lower productivity. Accordingly, in the present disclosure, it may be preferable to limit the N content to 0.001-0.010%.

Ti: 0.010-0.120%

Ti may be a representative precipitation strengthening element along with Nb and V, and may form coarse TiN in steel due to strong affinity with N. TiN may have the effect of suppressing grain growth during the heating process for hot-rolling. Also, Ti remaining after reacting with nitrogen may become solid-solute in steel and may combine with carbon and TiC precipitate may be formed, such that Ti may be a useful component for improving strength of steel. Accordingly, when the Ti content is less than 0.010%, the effect may not be obtained, and when the Ti content exceeds 0.120%, formability may deteriorate due to the occurrence of coarse TiN and coarsening of TiC precipitate. Accordingly, in the present disclosure, it may be preferable to limit the Ti content to 0.010-0.120%. More preferably, the content may be limited to the range of 0.070-0.115%.

Nb: 0.010-0.050%

Nb may be a representative precipitation strengthening element along with Ti and V, and may be effective in improving strength and impact toughness of steel due to the grain refinement effect caused by precipitation during hot-rolling and the delay in recrystallization. However, when the Nb content is less than 0.010%, the effect may not be obtained, and when the Nb content exceeds 0.050%, formability may deteriorate due to the formation of elongated grains and the formation of coarse composite precipitates due to excessive recrystallization delay during hot-rolling. Accordingly, in the present disclosure, it may be preferable to limit the Nb content to 0.010-0.050%. More preferably, the content may be limited to the range of 0.015-0.040%.
Also, in the present disclosure, one or more of Mo, Cr, V, Ni, and B may be additionally added if necessary, and in this case, the total content thereof may be within 1.500%.
Mo and Cr may delay ferrite transformation and may be advantageous in ensuring a low-temperature transformation structure such as bainite, and both elements may also contribute to ensuring strength by forming carbide by combining with C. Ni may be an austenite stabilizing element and may have a greater hardenability effect than the two elements, such that Ni may greatly contribute to increasing strength by ensuring a low-temperature transformation structure. B may also be an effective hardenability element and may produce the same effect even in much smaller amount (tens of ppm units) than the above elements. V may be a precipitation element precipitating in a low-temperature temperature range as compared to Nb and Ti, and may have the advantage of increasing strength through a precipitation strengthening effect.

Relational expression 1

In the present disclosure, the C, Mn, Ti, Nb and N contents may be controlled such that the X value defined by relational expression 1 as below may satisfy 3.50-6.00.
In the present disclosure, the main elements determining the strength and microstructure of the steel sheet may be hardenability elements such as C and Mn and precipitation elements such as Ti and Nb. Also, in order to reduce the changes in strength after heat treatment (in order to increase thermal stability), the fraction of low-temperature transformation phases such as bainite and martensite among the constituent microstructures may need to be reduced. In particular, when the fraction of martensite, which is a hard structure, is high, there may be a significant decrease in strength depending on the heat treatment conditions. Also, the proper design of precipitation strengthening elements may also be important for reducing the strength difference before and after heat treatment. This may be because the re-solid-solution, generation, and size change of precipitates may occur during the reheating, hot-rolling, and coiling processes, and also the additional heat treatment process, and may affect the material strength. The present disclosure may derive and present X, an alloy composition design factor for ensuring proper thermal stability, from C, Mn, which are major hardenability elements, and Nb and Ti, which are precipitated elements, related to the microstructure and precipitation strengthening.
When the X value in the present disclosure is less than 3.50, it may be highly likely that sufficient hardenability elements or precipitation hardening elements may not be added, and in this case, it may be difficult to ensure the desired tensile strength, and the material deviation may be large due to the deviation in the hot-rolled manufacturing conditions. Also, when the X value exceeds 6.00, hardenable elements or precipitation hardening elements may be excessively added, such that strength may increase excessively, and elongation may decrease, or due to a decrease in strength or re-precipitation caused by softening of the secondary hard phase after additional heat treatment, material changes may become severe. $\begin{array}{l} X = A * B \\ A = 1.3 * [Mn] + 200 * [C] \\ B = (Nb / 93 + Ti * / 48) / (C / 12 + N / 14) \\ Ti * = Ti - 3.42 N - 1.5 S \end{array}$
where Mn, C, Nb, Ti, N and S represent weight% thereof
In the present disclosure, other components and the remainder may be Fe and inevitable impurities.
Also, the hot-rolled steel sheet of the present disclosure may have a microstructure including a sum of ferrite and bainite phases: 90 area% or more, and a sum of remainder of martensite and MA phases: less than 10 area%. When the phase fraction of the sum of ferrite and bainite phases is less than 90 area%, the remaining pearlite or the sum of martensite and MA phases may exceed 10%, such that hole expandability may deteriorate. Hole expandability may be greatly affected by the microstructural composition of the steel sheet, and in particular, when a soft phase and a hard phase are complexly formed in the steel sheet, hole expandability may be greatly deteriorated due to the hardness difference between each the phases. In particular, as the fraction of pearlite or martensite, which is a hard microstructure, increases, cracks may be easily created at the interphase interface during hole expansion, which deteriorates the hole expandability, and thus, it may be necessary to limit the phase fraction.
The hot-rolled steel sheet of the present disclosure having the microstructure described above may have tensile strength of 590 MPa or more, hole expandability (HER₀) of 40% or more, and a bake hardening amount (BH₂) of 30 MPa or more.
Also, the hot-rolled steel sheet of the present disclosure may have a bake hardening amount (BH_h) after heat treatment at 300-600°C of 30 MPa or more, and may have a high-temperature bake hardening properties satisfying the absolute value of ΔTS × BH_h ^-1 of 0.70 or less when defining ΔTS as in relational expression 2 below. That is, the hot-rolled steel sheet of the present disclosure may maintain a bake hardening amount (BH_h) of 30MPa or more even after heat treatment at 300-600°C, such that a plated steel sheet may be effectively manufactured using a subsequent hot-dip galvanizing process, or the like. $ΔTS = {TS}_{h} - {TS}_{0}$
TS_h : tensile strength after heat treatment, TS₀ : tensile strength before heat treatment
Thereafter, the method for manufacturing the hot-rolled steel sheet of the present disclosure according to a preferable embodiment of the present disclosure may be described in detail.
The method for manufacturing the hot-rolled steel sheet of the present disclosure may include reheating a steel slab having the above-described alloy composition and of which the X value defined by relational expression 1 satisfies 3.5-6.0, in the temperature range of 1100-1350°C; manufacturing a hot-rolled steel sheet by hot-rolling the reheated steel slab in a range of 850-1150°C; and cooling the hot-rolled steel sheet to a temperature in a range of 400-550°C at an average cooling rate of 10-70°C/sec, and coiling the steel sheet.

Reheating

First, in the present disclosure, the steel slab having the alloy composition described above may be reheated in the temperature range of 1100-1350°C. In this case, when the reheating temperature is lower than 1100°C, the re-solid-solution rate of precipitates including Ti, Nb, Mo and V may decrease, such that the formation of fine precipitates may be reduced in the process after hot-rolling. When the content exceeds 1350°C, strength may decrease due to coarsening of austenite grains, and thus, it may be preferable to limit the reheating temperature to 1100-1350°C.

Hot-rolling

Thereafter, in the present disclosure, the hot-rolled steel sheet may be manufactured by hot-rolling the reheated steel slab in the range of 850-1150°C. In this case, when hot-rolling is started at a temperature higher than 1150°C, the temperature of the hot-rolled steel sheet may increase, such that the grain size may become coarse and the surface quality of the hot-rolled steel sheet may deteriorate. Also, when the hot-rolling is terminated at a temperature lower than 850°C, the elongated grains may develop due to excessive recrystallization delay, such that anisotropy may become severe and formability may deteriorate.

Cooling and coiling

In the present disclosure, the hot-rolled steel sheet may be cooled to a temperature in the range of 400-550°C at an average cooling rate of 10-70°C/sec, and may be coiled.
When the hot-rolled steel sheet is cooled below 400°C and coiled, low-temperature phases such as martensite and MA phases may be unnecessarily formed in the steel, which decreases thermal stability of the structure, such that formability may deteriorate both before and after heat treatment, and the decrease in strength may increase after heat treatment. When the coiling is performed by cooling above 550°C, the appropriate fractions of the bainite, martensite, and MA phases may not be ensured, such that it may be impossible to ensure the BH value both before and after heat treatment.
Also, when the average cooling rate is less than 10°C/sec during cooling, the grains of the matrix structure may become coarse and the microstructure may become uneven. When the average cooling rate exceeds 70°C/sec, the low-temperature phase fraction may increase, causing problems similar to those mentioned above when coiling below 400°C.
Preferably, the cooling temperature may be limited to 400-500°C.
The present disclosure may additionally include pickling and oiling the coiled hot-rolled steel sheet, if necessary.
Also, if necessary, the method may further include pickling the coiled hot-rolled steel sheet, heating the steel sheet to a temperature range of 450-750°C, and immersing the steel sheet in a plating bath containing 0.01-30% of Mg, 0.01-50% of Al, and the remainder of zinc in weight%, thereby forming a hot-dip galvanizing layer on the surface thereof. Mode for Invention
Hereinafter, the present disclosure may be described more specifically through examples. However, it should be noted that the examples as below are only intended to describe the present disclosure in greater detail, and are not intended to limit the scope of the rights in the present disclosure.

(Example)

The steel slabs having the alloy compositions listed in Table 1 below was prepared, and these slabs were reheated at 1200°C. Thereafter, the reheated slabs were hot-rolled, cooled, and coiled under the conditions listed in Table 2 below and hot-rolled steel sheets were manufactured. In this case, the cooling rate of the coiled hot-rolled steel sheets was maintained at the level of 0.5-10°C/s.
The microstructure composition and fraction of the hot-rolled steel sheets manufactured as described above were measured, and the results are listed in Table 2 below.
Specifically, the fractions of ferrite phase (F), bainite phase (B), martensite phase (M), and pearlite phase (P) were measured from the results analyzed at ×3000 and ×5000 magnifications using SEM. To identify martensite and MA phases, etching with Nital and Lepera was performed and analysis was performed at ×1000 magnification using an optical microscope and an image analyzer.
Also, the tensile strength (TS₀), bake hardening amount (BH₂), and hole expandability (HER₀) of the manufactured hot-rolled steel sheet were measured, and the results are listed in Table 3.
The tensile strength and bake hardening amount were tested by collecting DIN standard test pieces in the rolling direction, and the tensile evaluation was performed at room temperature.
The bake hardening amount was measured as the difference between the strength value at 2% pre-strain at room temperature and the strength value after carrying out heat treatment at 170°C for 20 minutes after 2% pre-strain and cooling at room temperature. In particular, when measuring the bake hardening amount, the strength after heat treatment at 170°C for 20 minutes after pre-strain was measured as the lower yield strength, and the lower bake hardening amount was measured.
The hole expandability was the average value of the results evaluated three times at room temperature. Specifically, the test was stopped when cracks were created visually in each test, and the long axis length of the crack was measured to evaluate the hole expandability, which was measured regardless of the rolling direction of the specimen.

Also, additional heat treatment was performed on the steel sheets of comparative example 1-9 and inventive example 1-5 under the conditions listed in Table 4 below. That is, the hot-rolled steel sheets were heat treated at the heat treatment temperature of 500°C for 10 minutes, and air-cooled to room temperature. After performing this heat treatment, the mechanical properties of the steel sheets before and after the heat treatment were evaluated, and are also listed in Table 4 below. Specifically, the tensile strength and bake hardening amount of the steel sheet after heat treatment were measured, and the results were compared with the tensile strength and bake hardening amount of the steel sheet before heat treatment. In this case, the method of measuring the tensile strength and bake hardening amount after heat treatment is as described above.

[Table 1]

Classi ficati on	Alloy composition (weight%)											X value
Classi ficati on	C	Si	Mn	Al	P	S	N	Ti	Nb	Cr	Mo	X value
Compar ative steel 1	0.09	0.33	1.9	0.029	0.009	0.002	0.004	0.011	0.05	-	0.009	1.10
Compar ative steel 2	0.011	0.34	1.6	0.025	0.011	0.002	0.003	0.101	0.021	0.592	-	7.77
Compar ative steel 3	0.075	0.61	1.7	0.023	0.011	0.002	0.003	0.105	0.045	0.037	-	6.38
Compar ative steel 4	0.043	0.31	2.1	0.028	0.012	0.001	0.003	0.033	0.042	0.009	-	2.37
Comparative steel 5	0.021	0.11	0.6	0.031	0.011	0.002	0.003	0.032	0.045	0.008	0.009	2.22
Compar ative steel 6	0.078	0.11	1.7	0.028	0.012	0.002	0.003	0.151	0.015	0.395	0.135	8.04
Compar ative steel 7	0.071	0.11	1.7	0.028	0.011	0.002	0.004	0.092	0.079	0.395	0.135	6.40
Compar ative steel 8	0.072	0.13	1.4	0.032	0.011	0.001	0.004	0.065	0.015	0.581	-	3.09
Compar ative steel 9	0.072	0.13	1.4	0.032	0.011	0.001	0.004	0.065	0.015	0.581	-	3.09
Invent ive steel 1	0.074	0.11	0.8	0.028	0.009	0.001	0.004	0.093	0.021	-	-	4.53
Invent ive steel 2	0.032	0.32	0.8	0.025	0.012	0.002	0.004	0.072	0.036	-	0.011	3.88
Invent ive steel 3	0.041	0.43	0.8	0.024	0.011	0.002	0.003	0.095	0.036	0.011	0.093	5.32
Invent ive steel 4	0.049	0.35	1.3	0.025	0.009	0.001	0.004	0.112	0.015	0.395	0.135	5.73
Invent ive steel 5	0.071	0.35	1.5	0.028	0.009	0.002	0.003	0.113	0.015	0.598	-	5.90

* In Table 1, the residuals are Fe and inevitable impurities.

[Table 2]

Classif ication	FDT(°C)	Cooling		Steel sheet microstructure phase fraction (area%)					Notes
Classif ication	FDT(°C)	CT(°C)	Cooling rate (°C/s)	F	B	M	MA	P	Notes
Compara tive steel 1	890	440	40	44	44	9	3	0	Compara tive example 1
Compara tive steel 2	890	440	40	98	2	0	0	0	Compara tive example 2
Compara tive steel 3	900	480	35	66	5	20	9	0	Compara tive example 3
Compara tive steel 4	900	480	40	75	6	17	2	0	Compara tive example 4
Compara tive steel 5	900	480	45	96	3	1	0	0	Compara tive example 5
Compara tive steel 6	900	480	50	89	6	3	2	0	Compara tive example 6
Compara tive steel 7	900	480	55	87	7	4	2	0	Compara tive example 7
Compara tive steel 8	890	650	50	91	3	0	0	6	Compara tive example 8
Compara tive steel 9	890	100	65	69	5	24	2	0	Compara tive example 9
Inventi ve steel 1	890	440	50	70	25	3	2	0	Inventi ve example 1
Inventi ve steel 2	900	480	60	86	11	2	1	0	Inventi ve example 2
Inventi ve steel 3	880	480	60	62	33	3	2	0	Inventi ve example 3
Inventi ve steel 4	900	480	55	66	30	2	2	0	Inventi ve example 4
Inventi ve steel 5	890	445	55	65	30	3	2	0	Inventi ve example 5
Inventi ve steel 1	890	440	95	57	28	13	2	0	Compara tive example 10
Inventi ve steel 3	880	480	3	75	13	0	0	12	Compara tive example 11
Inventi ve steel 4	900	480	2	77	10	0	0	13	Compara tive example 12
Compara tive steel 8	890	460	60	75	18	4	3	0	Compara tive example 13
Compara tive steel 9	890	460	40	80	17	3	0	0	Compara tive example 14

* In Table 2, F represents ferrite, B represents bainite, M represents martensite, MA represents martensite and austenite constituent, and P represents pearlite.

[Table 3]

Classi ficati on	Heat treatment conditions		Properties before heat treatment				Properties after heat treatment		Absolu te value of △TS × BH_h ^-1
Classi ficati on	Tempera ture (°C)	Time (min)	TS_o	BH2	HER_o	YRo	TS_h	BH_h	Absolu te value of △TS × BH_h ^-1
Compar ative exampl e 1	500	10	932	98	29	0.79	901	25	1. 24
Compar ative exampl e 2	500	10	510	1	89	0.88	509	1	1.00
Compar ative exampl e 3	500	10	947	70	35	0.71	812	33	4.09
Compar ative example 4	500	10	791	61	33	0.72	765	30	0.87
Compar ative exampl e 5	500	10	325	1	102	0.88	324	1	1.00
Compar ative exampl e 6	500	10	817	5	49	0.89	814	3	1.00
Compar ative exampl e 7	500	10	865	3	37	0.91	863	1	2.00
Compar ative exampl e 8	500	10	713	2	52	0.88	711	1	2.00
Compar ative exampl e 9	500	10	801	81	39	0.65	743	21	2.76
Invent ive exampl e 1	500	10	734	65	63	0.79	733	39	0.03
Invent ive exampl e 2	500	10	633	57	81	0.81	631	40	0.06
Invent ive exampl e 3	500	10	765	65	55	0.77	765	41	0.00
Invent ive exampl e 4	500	10	813	67	60	0.79	810	41	0.07
Invent ive exampl e 5	500	10	821	76	53	0.76	818	42	0.07
Compar ative exampl e 10	500	10	781	19	24	0.77	697	15	5.60
Compar ative exampl e 11	500	10	742	17	22	0.81	735	4	1.75
Comparative exampl e 12	500	10	752	14	19	0.82	740	3	4.00
Compar ative exampl e 13	500	10	772	39	38	0.78	765	8	0.88
Compar ative exampl e 14	500	10	764	37	40	0.79	760	3	1.33

* In Table 3, TS₀ and BH₂ represent the tensile strength and bake hardening amount before heat treatment, respectively. TS_h and BH_h represent the tensile strength and bake hardening amount after heat treatment, respectively.

ΔTS = {TS}_{h} - {TS}_{0}

As listed in Table 1-3, inventive examples 1-5 satisfying the steel sheet component range, X value, and manufacturing conditions suggested in the present disclosure were able to ensure the target material. Specifically, the hot-rolled steel sheets of inventive examples 1-5 exhibited tensile strength of 590 MPa or more, hole expandability (HER₀) of 40% or more, and a bake hardening amount (BH₂) of 30 MPa or more. Also, after additional heat treatment, the bake hardening amount (BH_h) was maintained at 30 MPa or more, and the absolute value of ΔTS × BH_h ^-1 satisfied 0.70 or less, such that the hot-rolled steel sheet may be effectively applied to various plating processes.
Differently from the above examples, comparative steels 1-9 may not satisfy the component range and/or manufacturing process conditions suggested in the present disclosure.
Specifically, comparative examples 1, 3, and 4 did not satisfy the relational expression due to the excess of C and Mn content, and the martensite phase and MA phase were unnecessarily formed, such that the hole expandability value of the steel sheet was deteriorated or a decrease in tensile strength after heat treatment was large.
In comparative examples 2 and 5, C and Mn contents were insufficient, and sufficient low-temperature phase fraction was not ensured due to the decrease in hardenability, such that strength of the steel sheet did not satisfy 590 MPa, and the bake hardening amount was deteriorated both before and after heat treatment.
In comparative steel 6 and comparative example 7, Ti and Nb contents were excessive, and the low-temperature phase fraction was not ensured due to excessive carbide formation, and the bake hardening amount deteriorated both before and after heat treatment, and the hole expandability also remained at the lower limit or did not reach the target due to the increase in coarse precipitates.
Also, in comparative examples 8-9, the X value by relational expression 1 was beyond the range of the present disclosure and the coiling temperature was beyond the suggested range of the present disclosure. That is, when the coiling temperature exceeded as in comparative example 8, it may be difficult to form a low-temperature phase within the structure, such that it may be difficult to ensure a bake hardening amount before and after heat treatment. When the temperature was not satisfied as in comparative example 9, the low-temperature phase fraction within the structure increased unnecessarily, such that the yield ratio deteriorated, and also the strength and bake hardening values before and after heat treatment were greatly changed.
In comparative examples 10-12, the alloy composition was within the range of the present disclosure, but the manufacturing conditions did not satisfy the range of the present disclosure. In comparative example 10, the cooling rate to the coiling temperature was too high, such that the fraction of martensite in the microstructure exceeded 10%, such that deterioration of hole expandability and strength deviation after heat treatment were severe. In comparative examples 11 and 12, the cooling rate was too low to the coiling temperature, such that the pearlite fraction exceeded 10% in the microstructure, such that hole expandability and the bake hardening properties were deteriorated, and there was a decrease in the strength/baking hardening value after heat treatment.
In comparative examples 13-14, as mentioned above, the X value by relational expression 1 was beyond the range of the present disclosure, but the manufacturing conditions satisfied the range of the present disclosure, and the absolute value of ΔTS × BH_h ^-1 did not satisfy 0.7 or lower depending on changes in the tensile strength/baking hardening value before and after heat treatment.
FIG. 1 is an image indicating changes (absolute value of ΔTS × BH_h ^-1) in tensile strength and bake hardening amount before and after heat treatment for X values of steel sheets of inventive examples and comparative examples. As indicated in FIG. 1, the hot-rolled steel sheet of the inventive example of the present disclosure had excellent thermal stability with an absolute value of ΔTS × BH_h ^-1 of 0.70 or less.
Although the present disclosure has been described in detail through embodiments above, other forms of embodiments may also be possible. Accordingly, the technical spirit and scope of the claims described below are not limited to the embodiments.

Claims

A hot-rolled steel sheet, comprising:
by weight%, C:0.020-0.080%, Si:0.01-0.50%, Mn:0.8-1.8%, Al:0.010-0.100%, P:0.001-0.020%, S:0.001-0.010%, N:0.001-0.010%, Ti: 0.010-0.120%, Nb:0.010-0.050%, and a balance of Fe and inevitable impurities, wherein an X value defined by relational expression 1 as below satisfies 3.50-6.00, and

having a microstructure comprising a sum of ferrite and bainite phases: 90 area% or more, and a sum of remainder of martensite and MA phases: less than 10 area%,

wherein the steel sheet has tensile strength of 590 MPa or more, hole expandability (HER₀) of 40% or more, and a bake hardening amount (BH₂: bake hardening amount before heat treatment) of 30 MPa or more,

the steel sheet maintaining a bake hardening amount (BH_h) after heat treatment at 300-600°C of 30 MPa or more, and

when ΔTS is defined as in relational expression 2 below, an absolute value of ΔTS × BH_h ^-1 is 0.70 or less: $\begin{array}{l} X = A * B \\ A = 1.3 * [Mn] + 200 * [C] \\ B = (Nb / 93 + Ti * / 48) / (C / 12 + N / 14) \\ Ti * = Ti - 3.42 N - 1.5 S \end{array}$

where Mn, C, Nb, Ti, N and S represent weight% thereof $ΔTS = {TS}_{h} - {TS}_{0}$

TS_h : tensile strength after heat treatment, TS₀ : tensile strength before heat treatment

BH_h : bake hardening amount after heat treatment
The hot-rolled steel sheet of claim 1, wherein the hot-rolled steel sheet comprises one or more of Mo, Cr, V, Ni, and B in a total amount of 1.500% or less.
The hot-rolled steel sheet of claim 1, wherein hot dip galvanizing plating is formed on a surface of the hot-rolled steel sheet.
A method for manufacturing the hot-rolled steel sheet,
the method comprising:
reheating a steel slab comprising, by weight%, C:0.020-0.080%, Si:0.01-0.50%, Mn:0.8-1.8%, Al:0.010-0.100%, P:0.001-0.020%, S:0.001-0.010%, N:0.001-0.010%, Ti: 0.010-0.120%, Nb:0.010-0.050%, and a balance of Fe and inevitable impurities, wherein an X value defined by relational expression 1 as below satisfies 3.50-6.00, in a temperature range of 1100-1350°C;

manufacturing a hot-rolled steel sheet by hot-rolling the reheated steel slab in a range of 850-1150°C; and

cooling the hot-rolled steel sheet to a temperature in a range of 400-550°C at an average cooling rate of 10-70°C/sec, and coiling the steel sheet. $\begin{array}{l} X = A * B \\ A = 1.3 * [Mn] + 200 * [C] \\ B = (Nb / 93 + Ti * / 48) / (C / 12 + N / 14) \\ Ti * = Ti - 3.42 N - 1.5 S \end{array}$

where Mn, C, Nb, Ti, N and S represent weight% thereof
The method of claim 4, wherein the hot-rolled steel sheet comprises one or more of Mo, Cr, V, Ni, and B in a total amount of 1.500% or less.
The method of claim 4, further comprising:
pickling and oiling the coiled hot-rolled steel sheet.
The method of claim 4, further comprising:
pickling the coiled hot-rolled steel sheet, heating the steel sheet in a temperature range of 450-750°C, and hot-dip galvanizing the steel sheet.