EP4600399A1

EP4600399A1 - Steel sheet and method for manufacturing same

Info

Publication number: EP4600399A1
Application number: EP23907585.6A
Authority: EP
Inventors: Yeon-Sang Ahn; Hang-Sik Cho; Joo-Hyun Ryu
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2022-12-21
Filing date: 2023-12-14
Publication date: 2025-08-13
Also published as: WO2024136299A1; CN120303434A; MX2025005651A; EP4600399A4; JP2025537573A; KR20240098907A

Abstract

The present invention relates to a high-strength steel sheet used for a vehicle structural member, etc. and, more specifically, to a steel sheet having excellent formability and crashworthiness, and a method for manufacturing same.

Description

Technical Field

The present disclosure relates to a high-strength steel sheet used for automobile structural members, or the like, and more particularly, a steel sheet having excellent formability and fracture resistance and a method for manufacturing the same.

Background Art

As environmental and safety regulations in the automobile industry have been increasingly stringent, carbon dioxide emission regulations and fuel efficiency regulations have also become more stringent. The Insurance Institute for Highway Safety in the United States has gradually strengthened crash safety regulations for occupant protection, and has required rigorous crash performance such as 25% small overlap since 2013.
The only solution to these environmental and safety issues may be reduction of automobile weight. To achieve weight reduction, high-strength steel may be necessary, and excellent formability may also be required while ensuring high-strength. Also, to assure the crash performance of the body, fracture resistance may need to be improved.
Generally, high-strength automobile materials may be classified into precipitation-strengthened steel, bake-hardened steel, solid solution-strengthened steel, and transformation-strengthened steel.
The transformation-strengthened steels may include dual phase steel (DP steel), complex phase steel (CP steel), and transformation induced plasticity (TRIP steel). These transformation-strengthened steels are also referred to as advanced high-strength steel (AHSS).
Among these, DP steel may ensure high-strength as hard martensite is finely and homogeneously dispersed within soft ferrite, and CP steel may include two or three phases of ferrite, martensite, and bainite, and precipitation hardening elements such as Ti and Nb may be added to DP to improve strength. TRIP steel may include finely and homogeneously dispersed residual austenite, and when processed at room temperature, the residual austenite phase may transform into martensite, thereby ensuring high-strength and high ductility.
Recently, a steel sheet for automobiles has been required to have higher strength for improved fuel efficiency and durability, and for collision safety and passenger protection, an ultra-high-strength steel sheet with a tensile strength of 980 MPa or more may be increasingly used for a vehicle body structure or as a reinforcing material.
In particular, high-strength steel having excellent yield strength may be employed for structural members such as a member, a seat rail, and a pillar to improve collision resistance of the vehicle body. As the yield strength (YS) of the structural member increases as compared to tensile strength (TS), that is, the yield ratio (YR = YS / TS), the member may be advantageous for collision resistance.
However, since strength of the steel sheet increases, ductility may decrease, such that formability and processability may deteriorate, and the development of materials which may supplement the problems may be necessary. In other words, in order to secure both crash stability and part formability, it may be essential to develop a material having high yield strength and excellent ductility.
In order to improve a yield ratio (YR) of steel, yield strength may need to be improved relative to tensile strength. A representative method for achieving this may be a method of using water cooling during continuous annealing. In this method, cracks may be created in the annealing process, the steel may be immersed in water to form martensite, and through a tempering process, a steel sheet having a microstructure of tempered martensite may be manufactured.
As a prior art related to the above technology, cited document 1 discloses a martensitic steel having a volume ratio of martensite of 80 to 97% by continuously annealing steel including carbon (C) of 0.18% or more, cooling to room temperature, and overaging at a temperature of 120 to 300°C for 1 to 15 minutes. When manufacturing ultra-high-strength steel by the water-cooling and tempering method as above, the yield ratio may be relatively high, but the shape quality of the coil may deteriorate due to temperature deviation in the width direction and length direction, such that cracks may be created during forming and workability may be reduced.
Cited document 2 discloses a method for manufacturing a high-strength steel sheet having a complex phase mainly including martensite, in which fine precipitated copper particles having a particle size of 1 to 100 nm are dispersed in the structure to improve workability. However, to precipitate fine copper particles, Cu may be excessively added, 2-5 weight%, such that red-hot embrittlement may occur due to Cu, and manufacturing costs may increase excessively.
Cited document 3 is a precipitation-strengthened steel sheet having ferrite as a base structure and including 2 to 10 area% of pearlite, and attempts to improve strength through precipitation strengthening and grain refinement by adding carbon and nitride forming elements such as Nb, Ti, and V. This steel sheet may have good hole expandability, but may have limitations in increasing tensile strength, may have high yield strength and low ductility, such that cracks may be created during press forming.
Cited document 4 discloses a method for manufacturing a cold rolled steel sheet simultaneously ensuring high-strength and high ductility by using tempered martensite and also having an excellent plate shape after continuous annealing. However, the carbon content may be high, 0.2% or more, such that weldability may deteriorate and due to the large amount of Si content, denter defects may occur in the furnace.

(Cited document 1) Japanese Laid-Open Patent Publication No. 1992-289120
(Cited document 2) Japanese Laid-Open Patent Publication No. 2005-264176
(Cited document 3) Korean Laid-Open Patent Publication No. 2015-0073844
(Cited document 4) Japanese Laid-Open Patent Publication No. 2010-090432

Detailed description of present disclosure

Technical problems to solve

An aspect of the present disclosure is to provide a steel sheet suitable for an automobile structural member, and more particularly, a steel sheet having high strength and also excellent formability and fracture resistance, and a method for manufacturing same.
The purpose of the present invention is not limited to the above-described features. The subject matter of the present disclosure may be understood from the entire descriptions in this specification, and a person having ordinary knowledge in the technical field to which the present disclosure belongs may have no difficulty in understanding the additional subject matter of the present disclosure.

Solution to Problem

An aspect of the present disclosure provides a steel sheet comprising, by weight%, carbon (C): 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn): 1.9-3.0%, molybdenum (Mo): 0.3% or less, chromium (Cr): 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities,

wherein C, Si, Al, Mn, Cr and Mo satisfy relational expression 1 as below,
wherein a microstructure comprises, by area%, ferrite: 10-35%, residual austenite: 3-15%, fresh martensite: 20% or less (excluding 0%), and a remainder of one of tempered martensite and bainite, and
wherein, in fresh martensite, a fraction of fresh martensite in which a distance between fresh martensite phases is three or more times a grain size of fresh martensite is 30% or more:

$(10 C + Si + Al) / (Mn + 1.3 Cr + 2.7 Mo) \geq 0.7$
where each element indicates a weight content.

A fraction of fresh martensite satisfies relational expression 2 as below:

$({FM}_{TM + B} / {FM}_{T}) \times 100 \geq 80 %$
where FM_T indicates a total fraction of fresh martensite, and FM_TM+B indicates a fraction of fresh martensite in contact with tempered martensite or bainite in the fraction of fresh martensite.
Yield strength (YS), uniform elongation (U-El) and tensile strength (TS) of the steel sheet satisfy relational expression 3:

$YS \times U - El / TS \geq 6$
Yield strength (YS), post uniform elongation (P-El) and tensile strength (TS) of the steel sheet satisfy relational expression 4 as below:

$YS \times P - El / TS \geq 3$

the steel sheet comprises boron (B): 0.01% or less.
the steel sheet further includes one or more of titanium (Ti): 0.05% or less and niobium (Nb): 0.05% or less.
the steel sheet further comprises a zinc-based plating layer.

An aspect of the present disclosure provides a method of manufacturing a steel sheet comprising heating a steel slab comprising, by weight%, carbon (C): 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn): 1.9-3.0%, molybdenum (Mo): 0.3% or less, chromium (Cr): 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities, wherein C, Si, Al, Mn, Cr and Mo satisfy relational expression 1 as below, in a temperature range of 1100-1300°C;

manufacturing a hot-rolled steel sheet by finishing hot-rolling the reheated steel slab at a temperature range of Ar3 or higher;
winding the hot-rolled steel sheet in a temperature range of 400-700°C;
cooling the hot-rolled steel sheet to room temperature at a cooling rate of 0.1°C/s after the winding;
manufacturing a cold-rolled steel sheet by cold-rolling the hot-rolled steel sheet at a total reduction ratio of 30-80% after the cooling;
continuously annealing the cold-rolled steel sheet at a temperature of Ac1+30°C-Ac3+30°C;
primary-cooling the cold-rolled steel sheet having gone through the continuous annealing to 450-700°C at a cooling rate of 10°C/s or lower;
secondary-cooling the steel sheet to 250-500°C at a cooling rate of 5°C/s or more after the primary-cooling; and
reheating the secondary-cooled cold-rolled steel sheet at a temperature of 490°C or lower and holding the steel sheet for 20 seconds or more,
wherein the cold-rolling is performed with a cumulative reduction ratio of 25% or more of first 1-3 stands.

$(10 C + Si + Al) / (Mn + 1.3 Cr + 2.7 Mo) \geq 0.7$
where each element indicates a weight content.
the continuous annealing is performed at a temperature of 800-880°C.
the cooling rate during the secondary-cooling is higher than the cooling rate during the primary-cooling.
the secondary-cooling is performed in a hydrogen quenching facility using hydrogen (H₂) gas.

The method further comprises holding the steel sheet for 30 seconds after the secondary-cooling.
The method further comprises hot-dip galvanizing the steel sheet in a plating bath at 430-490°C after the reheating and holding.
The method further comprises performing alloying heat treatment after the hot-dip galvanizing.
The method further comprises cooling the steel sheet to a temperature of Ms-100°C or lower at a cooling rate of 5°C/s or higher and performing a skin pass of less than 2% after the alloying heat treatment.

Advantageous Effects of Invention

According to an aspect of the present disclosure, by providing a steel sheet having high tensile strength of 980 MPa or more and excellent formability, processing defects occurring during processing of parts may be prevented, thereby manufacturing various parts having complex shapes. Also, by improving fracture resistance, the steel sheet may contribute to improvement of safety of parts and vehicle.
The various beneficial advantages and effects of present disclosure are not limited to the above-described contents, and will be more easily understood in the process of describing the specific embodiment of present disclosure.

Brief Description of Drawings

FIG. 1 is may be a graph indicating changes in relational expression 3 according to relational expression 1 in the result of embodiment of present disclosure.
FIG. 2 may be a graph indicating changes in relational expression 4 according to relational expression 1 in the result of embodiment of present disclosure.

Best Mode for Invention

The terms used in this specification are intended to describe the present disclosure and are not intended to limit the present disclosure. An expression used in the singular encompasses the expression of the plural, unless it has a clearly different meaning in the context.
The term "comprising or including" used in this specification is to specify the components and may not to exclude the presence or addition of other components.
Unless otherwise indicated, the terms, including technical and scientific terms, used in this specification may mean the same as generally understood by a person of ordinary skill in the art to which the present disclosure belongs. Terms defined in the dictionary are to be interpreted to have a meaning consistent with the relevant technical literature and the present disclosure.
The inventors of the present disclosure have conducted in-depth research to provide a high-strength steel sheet with excellent formability and fracture resistance.
As a result, by optimizing the alloy composition and manufacturing conditions of the steel, it was confirmed that a preferable structure could be obtained to ensure target properties, and thus a steel sheet suitable for structural members for automobiles requiring processing into complex shapes could be provided, and the present disclosure was completed.
Hereinafter, the present disclosure may be described in detail.
First, an alloy composition of the steel sheet according to an aspect of the present disclosure may be described in detail.
Unless otherwise described in the present disclosure, the content of each element may be based on weight, and the ratio of the structure may be based on area.
The steel sheet may comprise carbon (C): 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn): 1.9-3.0%, molybdenum (Mo): 0.3% or less, chromium (Cr): 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities.

Carbon (C): 0.1-0.2%

C may be an important element added to strengthen the transformation structure of steel. C may promote high strength of steel and may facilitate the formation of martensite in complex phase steel. As the content of the C increases, the amount of martensite may increase.
When the C content exceeds 0.2%, strength due to the formation of martensite may increase, but the strength difference with ferrite having a low carbon concentration may increase. Due to this strength difference, fracture may easily occur at an interfacial surface between phases during plastic deformation, ductility and work hardening rate may decrease. Also, weldability may be deteriorated, such that welding defects may occur during processing of parts, and liquid metal embrittlement (LME) cracks may occur during welding, which may impair performance of the parts. When the C content is less than 0.1%, it may be difficult to ensure a target level of strength, and it may be difficult to ensure a predetermined fraction of residual austenite phase required for ductility. It may be preferable for the C content to be 0.10-0.20%. More preferably, the content may be 0.12-0.18%.

Silicon (Si): 0.5-1.3%

Si may be a ferrite stabilizing element and may facilitate the transformation of ferrite and may promote enrichment of carbon (C) into untransformed austenite, thereby contributing to the formation of martensite. Also, Si may have excellent solid-solution strengthening ability, which may be effective in increasing strength of ferrite and reducing the difference in hardness between phases. Also, by effectively suppressing the precipitation of carbides in bainite during holding in the bainite region, Si may promote the enrichment of C into untransformed austenite, thereby delaying the transformation of martensite during low-temperature rapid cooling, and forming residual austenite necessary for ductility, and thus, Si may be useful for improving ductility of the steel sheet. In other words, Si may be useful in that Si may ensure strength without decreasing ductility of the steel sheet.
When the content of Si exceeds 1.3%, surface scale defects may occur, which adversely affect the plating surface quality, may impede chemical processability, and may deteriorate weldability, such that welding defects may occur during processing of parts. In particular, LME cracks may occur during welding, which decreases performance of parts. When the content is less than 0.5%, it may be difficult to ensure a predetermined fraction of residual austenite phase required for ductility, and solid-solution hardenability may deteriorate, such that strength of ferrite may be lowered, and there may be limitations in reduction of the hardness difference between phases, and accordingly, formability may decrease. It may be more preferable that the content of Si be 0.50-1.30%. More preferably, the content may be 0.7-1.2%.

Aluminum (Al) 0.5% or less (excluding 0%)

Aluminum, more preferably acid-soluble aluminum (Sol.Al), may be added for grain refinement and deoxidation of steel, and may be a ferrite stabilizing element similarly to Si. Al may be useful for distributing carbon in ferrite to austenite to improve hardenability of martensite. Also, Al may effectively suppress precipitation of carbides in bainite during holding in the bainite region during annealing, thereby promoting C enrichment into untransformed austenite, and delaying martensite transformation during low-temperature rapid cooling, and Al may generate a residual austenite phase, thereby improving ductility of the steel sheet.
When the content of Al exceeds 0.5%, inclusions may be excessively formed during the steelmaking continuous casting operation, it may be highly likely that surface defects may occur on the surface of the steel sheet, and manufacturing costs may also increase. Also, weldability may be deteriorated, which may cause welding defects during processing of parts. Accordingly, the Al may be included in 0.5% or less, and 0% may be excluded. It may be preferable that the Al content be 0.50% or less.

Manganese (Mn): 1.9-3.0%

Mn may refine particles without decreasing ductility and may completely precipitate sulfur (S) in the steel as MnS, thereby preventing hot brittleness caused by FeS formation, and Mn may also be effective for strengthening steel. Also, Mn may facilitate the formation of martensite by lowering a critical cooling rate at which the martensite phase is obtained in complex phase steel.
When the Mn content is less than 1.9%, it may be difficult to ensure the target strength in the present disclosure. When the content exceeds 3.0%, it may be highly likely that there may be a problem in weldability and hot-rollability, and martensite may be formed excessively, the material may become unstable, Mn-bands may be formed within the structure, such that there may be a high risk that defects such as processing cracks and plate breakage may occur. Also, during the annealing process, Mn oxide may be dissolved on the surface, which may significantly deteriorate the surface quality.
Accordingly, in the present disclosure, the Mn may be included at 1.9-3.0%, and 1.90-3.00% may be more preferable. More preferably, the content may be 2.0-2.8%.

Molybdenum (Mo): 0.3% or less

Mo may be selectively included to delay the transformation of austenite into pearlite and to improve the refinement and strength of ferrite. Mo may improve hardenability of the steel and may control the yield ratio by forming martensite finely at the grain boundary. However, as Mo is an expensive element, the manufacturing cost may increase as the content thereof increases, which may be economically disadvantageous.
In order to sufficiently obtain the above-mentioned effect, Mo may be added up to 0.3%. When the Mo content exceeds 0.3%, the alloy cost may increase rapidly, which reduces economic feasibility, and the grain refinement effect and solid-solution strengthening effect may occur excessively, such that ductility of the steel may decrease. It may be preferable that the Mo content be 0.30% or less.

Chromium (Cr): 1.0% or less (excluding 0%)

Cr may be added to improve hardenability of the steel and to ensure high strength, and may be important in formation of martensite. Also, Cr may be preferable for the production of complex phase steel having high ductility by minimizing the decrease in elongation as compared to the increase in strength. In particular, Cr-based carbides such as Cr₂₃C₆ may be formed during the hot-rolling process, and some of these carbides may be dissolved during the annealing process and some may remain undissolved, such that the solid-solution C content in martensite may be controlled below an appropriate level after cooling, such that yield point elongation (YP-El) may be prevented, and thus, Cr may be preferable for the production of complex phase steel having a low yield ratio.
When the Cr content exceeds 1.0%, the above-mentioned effect may be saturated, and the hot-rolled strength may increase excessively and the cold-rollability may deteriorate. Also, the Cr-based carbides may be formed excessively and coarsen, the martensite size may become coarse after annealing, causing a decrease in elongation. It may be more preferable that the Cr content be 1.00% or less.

Phosphorus (P): 0.1% or less

P may be a substitutional element having a large solid-solution strengthening effect, and may improving the anisotropy in the plane, may not significantly damage formability and may be most preferable element for ensuring strength. However, when the content of the P is excessive, the possibility of brittle fracture may greatly increase, slab plate breakage may occur during hot-rolling and the plating surface properties may be degraded. Accordingly, P may be included at 0.1% or less, and 0% may be excluded considering the level inevitably added during the steel manufacturing process.

Sulfur (S): 0.1% or less

S may be impurities inevitably added to steel, and as S may deteriorate ductility and weldability, and thus, it may be preferable to manage the content thereof as low as possible. In particular, since S may cause red-hot brittleness, it may be preferable to control the content to 0.01% or less. However, considering the level inevitably added during the steel manufacturing process, 0% may be excluded.

Boron (B): 0.01% or less

B may delay the transformation of austenite into pearlite during the cooling process during annealing, and may be selectively included to ensure hardenability to suppress formation of ferrite and promote the formation of martensite. When the content of B exceeds 0.01%, B may be excessively concentrated on the steel surface, such that plating adhesion may deteriorate. Accordingly, B may be included in 0.01% or less. It may be more preferable that the content of B be 0.010% or less.
In addition to the components mentioned above, titanium (Ti): 0.05% or less and niobium (Nb): 0.05% or less may be further comprised.
Ti and Nb may be effective for increasing strength of steel and refining grains by forming nano precipitates. When these elements are added, the elements may combine with carbon and may form fine nano precipitates, and these nano precipitates may strengthen the base structure and may thus reduce the difference in hardness between phases.
When the content of Ti and Nb exceeds 0.05%, the manufacturing cost may increase, and precipitates may be excessively formed, such that ductility may greatly decrease. Accordingly, when adding one or more of the Ti and Nb, each of Ti and Nb may be included in 0.05% or less. It may be more preferable that the content of Ti and Nb be 0.050% or less.
A remainder of the present disclosure is iron (Fe). However, in a general manufacturing process, inevitable impurities may be inevitably added from raw materials or an ambient environment, and thus, impurities may not be excluded. A person skilled in the art of a general manufacturing process may be aware of the impurities, and thus, the descriptions of the impurities may not be provided in the present disclosure.
As for the steel sheet, among the alloy compositions described above, it may be effective for C, Si, Al, Mn, Cr and Mo to satisfy relational expression 1 as below. $(10 C + Si + Al) / (Mn + 1.3 Cr + 2.7 Mo) \geq 0.7$
(Herein, each element indicates a weight content)
When relational expression 1 is satisfied, residual austenite which may contribute to ductility may be sufficiently ensured. Also, by optimizing the manufacturing process, a distance between fresh martensite phases having high hardness, which may receive large local stress concentration during press forming or impact, may be widened, and distribution may be finely and evenly distributed around the tempered martensite and bainite structures, such that, by relieving a local stress concentration occurring in fresh martensite, the creation and coalescence of voids occurring around fresh martensite may be delayed, and excellent formability without cracks may be ensured during the press forming process. Also, by delaying fracture even in the event of a vehicle collision, fracture resistance may also be improved. When relational expression 1 is not satisfied, it may be difficult to ensure a sufficient residual austenite fraction and fine fresh martensite may not be uniformly distributed, it may not be easy to ensure formability and fracture resistance. It may be preferable that relational expression 1 be 0.70 or less.
Thereafter, the microstructure of the steel sheet of the present disclosure may be described in detail. The microstructure may comprise, by area fraction, ferrite: 10-35%, residual austenite: 3-15%, fresh martensite: 20% or less (excluding 0%), and the remainder may include one or more of tempered martensite and bainite.
The ferrite may be important for ensuring formability and strength. As the fraction of ferrite increases, carbon may be accumulated in untransformed austenite, such that the martensite transformation temperature may be lowered below room temperature, and may thus contribute to ensuring residual austenite at room temperature. Also, since ferrite is easily transformed into a soft phase, ferrite may contribute to ensuring ductility. To this end, ferrite may be effective at 10% or more. When the fraction of the ferrite exceeds 35%, strength may decrease and the desired strength may not be ensured, and when it is less than 10%, the ductility contribution effect of ferrite may not be expected, and the accumulation of carbon in the austenite may be impaired, such that it may be difficult to ensure the desired fraction of residual austenite.
Also, in the present disclosure, when bainite is generated or the generated martensite may be tempered, carbon included in bainite or martensite may move to the surrounding untransformed austenite and may accumulate, such that the martensite transformation temperature may be lowered to room temperature or lower, and accordingly, residual austenite may be ensured at room temperature. When the residual austenite is 3% or more, transformation-induced plasticity may occur during forming, and it may be preferable to ensure ductility of the steel sheet. However, when the residual austenite is excessive, the plated steel sheet may tend to be vulnerable to liquid metal brittleness during spot welding of automobile parts, such that the content may be effective not to exceed 15%.
Fresh martensite may be effective in ensuring strength, and when the fraction exceeds 20%, fresh martensite may not be finely and evenly distributed around the tempered martensite or bainite structure. When fresh martensite is concentrated, a large local stress concentration may be received during forming or impact in accordance with high hardness properties, which may cause brittleness or fracture. In other words, formability may be reduced.
It may be effective that the remainder of the microstructure be at least one of tempered martensite and bainite phases. When at least one of the phases is included at 40% or more, carbon may be concentrated in untransformed austenite, and residual austenite contributing to ductility may be generated in a fraction of 3-15%, and the fraction of fresh martensite may also be controlled to 20% or less. Accordingly, fresh martensite may be dispersed finely and evenly around tempered martensite or bainite structures. However, when the fraction is less than 40%, the desired residual austenite fraction may not be ensured, and the fraction of fresh martensite ultimately generated may also exceed 20%, such that it may be difficult to evenly and finely disperse fine fresh martensite around tempered martensite or bainite.
Specifically, fresh martensite may be the microstructure having the highest strength among the microstructure phases. When fresh martensite is finely and evenly dispersed, and deformation is applied during the press forming process of a part, the local stress concentration applied to fresh martensite may be distributed to neighboring microstructures and may be alleviated. Accordingly, since the generation and coalescence of pores are delayed, excellent formability may be obtained without processing cracks during the forming of a part. When fresh martensite is coarse and unevenly distributed, stress may be locally concentrated in the fresh martensite such that the creation and coalescence of pores may be facilitated. Accordingly, processing cracks may occur during component molding, which may reduce formability. Also, when fresh martensite is finely and evenly distributed, local stress concentration is alleviated, such that the creation and coalescence of pores are delayed, thereby improving fracture resistance in the event of a vehicle crash. On the other than, when fresh martensite is coarse and unevenly distributed, stress may be locally concentrated in the fresh martensite such that the creation and coalescence of pores may be facilitated, which may deteriorate the fracture resistance in the event of a vehicle crash.
Also, in the steel sheet, the fresh martensite may have a fraction of fresh martensite in which the distance between fresh martensite phases may be three or more times of the grain size of fresh martensite. When the fraction of fresh martensite in which the distance between fresh martensite phases (L) is three or more times of the grain size (d) of the fresh martensite (L > 3d) is 30% or more, the local stress concentration in the fresh martensite previously described may be prevented and the effect of delaying the creation and coalescence of pores may be maximized, thereby improving formability and impact fracture resistance. When the fraction is less than 30%, the stress may be locally concentrated in the fresh martensite may be increased, such that the creation and coalescence of pores may be facilitated, thereby deteriorating formability and impact fracture resistance.
It may be effective that the fraction of the fresh martensite satisfy relational expression 2. This may indicate that implementing a microstructure in which fine fresh martensite is uniformly dispersed around tempered martensite or bainite may be effective in terms of formability and fracture resistance. When the fraction of fresh martensite in contact with tempered martensite or bainite mentioned in relational expression 2 exceeds 80% of the total fraction of fresh martensite, the evenly distributed effect of fine fresh martensite may maximize the local stress concentration relief effect and the creation and coalescence delay effect in fresh martensite, thereby improving formability and impact fracture resistance. However, when the fraction is less than 80%, the evenly distributed effect of fine fresh martensite may disappear, which increases local stress concentration in fresh martensite, such that the creation and coalescence of pores may be facilitated, thereby deteriorating the formability and impact fracture resistance. $({FM}_{TM + B} / {FM}_{T}) \times 100 \geq 80 %$
(Herein, FM_T indicates a total fraction of fresh martensite, and FM_TM+B indicates a fraction of fresh martensite in contact with tempered martensite or bainite in the fraction of fresh martensite)
The steel sheet of the present disclosure may have high strength of tensile strength (TS) of 980 MPa or more, yield strength (YS) of 700 MPa or more, and total elongation (T-El) of 13% or more. Also, as for the steel sheet of the present disclosure, relationship between tensile strength(TS), yield strength(YS) and uniform elongation(U-El) may satisfy relational expression 3 below, and tensile strength(TS), yield strength(YS) and post uniform elongation(P-El) may satisfy relational expression 4 below, thereby providing a steel sheet having excellent formability and fracture resistance. In relational expressions 3 and 4 below, the units of YS and TS may be MPa, and the units of U-El and P-El may be %. $YS \times U - El / TS \geq 6$ $YS \times P - El / TS \geq 3$
In the tensile test of a material, the strain until breakage in the load-displacement diagram may be defined as total elongation (T-El), the strain until the maximum load point may be defined as uniform elongation (U-El), and the strain from the maximum load point to breakage may be defined as post uniform elongation (P-El). When deformation is applied to a material, the deformation of the entire material may occur uniformly up to the uniform elongation point, and beyond this point, necking may occur in a predetermined part of the material. Accordingly, when the uniform elongation is excellent, excellent formability may be obtained without necking even at high strains during molding of parts. As mentioned above, relational expression 3 may indicate that, when fine fresh martensite may be evenly distributed by controlling the microstructure, excellent formability may be obtained despite high yield strength. Also, when the post uniform elongation is excellent, more deformation may have to accompany until the breakage of the final material even when necking occurs. It may be preferable that relational expression 3 be 6.0 or more. Relational expression 4 may indicate that, as mentioned above, when fine fresh martensite is evenly distributed by controlling the microstructure, yield strength may be high and the post uniform elongation may also be excellent, such that the generation and coalescence of pores in the structure during a vehicle collision may be delayed, and impact fracture resistance may thus be excellent. It may be preferable that relational expression 4 be 3.0 or more. In summary, when relational expressions 3 and 4 are satisfied, it may indicate that the steel material may be excellent, which may have excellent part formability, and excellent fracture resistance during a vehicle collision although yield strength is high.
The steel sheet of the present disclosure may be implemented as a cold-rolled steel sheet, and also a hot-dip galvanized steel sheet including a zinc-based plating layer on at least one surface of the cold-rolled steel sheet, or an alloyed hot-dip galvanized steel sheet alloyed with the hot-dip galvanized steel sheet.
Although not specifically limited, the zinc-based plating layer may be a zinc-plated layer including mainly zinc, or a zinc alloy plating layer including aluminum and/or magnesium in addition to zinc.
In the description below, another embodiment of the present disclosure, a method for manufacturing a steel sheet, may be described in detail.
Briefly, in the present disclosure, a target steel sheet may be manufactured through a process of [steel slab reheating - hot-rolling - winding - cold-rolling - continuous annealing - cooling - reheating and holding], and processes such as [hot-dip galvanized - alloying heat treatment] may be performed.
The conditions for each step may be described in detail below.

Steel slab heating

First, a steel slab satisfying the aforementioned alloy composition and relational expression 1 may be prepared, and the steel slab may be heated. This process may be performed to smoothly perform the subsequent hot-rolling process and to sufficiently obtain properties of the target steel sheet.
The heating process may be not specifically limited, and normal heating conditions may be sufficient. As a preferable example, the heating may be performed in a temperature range of 1100-1300°C. When the heating temperature is less than 1100°C, friction between the steel sheet and the rolling mill may increase, the load applied to the roller during hot rolling may increase rapidly. When the temperature exceeds 1300°C, the energy cost required for temperature increase may increase, and the amount of surface scale may also increase, which may lead to material loss.

Hot-rolling

According to the above, the heated steel slab may be finished hot-rolled at the Ar3 transformation point or higher and a hot-rolled steel sheet may be manufactured, and in this case, the hot-rolling conditions may not be particularly limited and may be performed at a normal hot-rolling temperature. As a preferable example, the finishing hot-rolling may be performed at a temperature range of 800-1000°C.

Winding

The hot-rolled steel sheet manufactured as above may be winded, and in this case, the winding may be performed at a temperature range of 400-700°C.
When the winding temperature is less than 400°C, the strength of the hot-rolled steel sheet may be excessively high, which may cause a rolling load during subsequent cold-rolling. Also, the cost and time required to cool the hot-rolled steel sheet to the winding temperature may be excessive, which may increase the process cost. When the temperature exceeds 700°C, excessive scale may occur on the surface of the hot-rolled steel sheet, which may cause surface defects and may deteriorate the plating properties.

Cooling

It may be preferable to cool the winded hot-rolled steel sheet to room temperature at a cooling rate of 0.1°C/s or lower (excluding 0°C/s). Here, cooling may indicate an average cooling rate.
By cooling the winded hot-rolled steel sheet at a constant rate as above, a hot-rolled steel sheet in which carbides becoming nucleation sites of austenite are finely dispersed may be obtained. That is, during the hot rolling process, fine carbides may be evenly dispersed in the steel, and the carbides may be dissolved during the subsequent annealing, such that an austenite phase may be finely dispersed and formed in the steel, and accordingly, after the annealing is completed, a uniformly dispersed fine martensite phase may be obtained.

Cold-rolling

According to the above, the wound hot-rolled steel sheet may be cold-rolled to manufacture a cold-rolled steel sheet, and in this case, the cold-rolling may be performed with a cold reduction ratio (total reduction ratio) of 30-80%.
In particular, in the present disclosure, by performing the initial stand, preferably the 1-3 stands, with a cumulative reduction ratio of 25% or more during the cold-rolling, the stored energy in the steel may be increased, thereby acting as a driving force to promote the recrystallization of ferrite in the subsequent annealing process. Accordingly, the effect of lowering the fraction of unrecrystallized ferrite in the steel may be provided.
When unrecrystallized ferrite is present in the steel, deformation and stress may be locally concentrated such that ductility of the steel may be deteriorated, whereas recrystallized ferrite relieves deformation and stress concentration and may thus contribute to improving ductility.
When the cumulative reduction ratio of the initial 1-3 stands during cold-rolling is less than 25%, or the cold reduction ratio (total reduction ratio) until the final stand is less than 30%, it may be difficult to ensure the target thickness, and may be also difficult to correct the shape of the steel sheet. Also, the fraction of unrecrystallized ferrite may increase, such that ductility may decrease. When the cold reduction ratio until the final stand during cold-rolling exceeds 80%, the strength may increase, which may cause roll load during cold-rolling, and it may be highly likely that cracks may be created at the edge portion of the steel sheet.
In the present disclosure, the cold-rolling may be performed using a rolling mill consisting of 5 or 6 stands, but an embodiment thereof is not limited thereto.

Continuous annealing

It may be preferable to continuously anneal the cold-rolled steel sheet manufactured as above. The continuous annealing treatment may be performed, for example, in a continuous alloying galvanizing furnace.
The continuous annealing may be a process for forming ferrite and austenite phases simultaneously with recrystallization and decomposing carbon.
The continuous annealing treatment may be performed in a temperature range of Ac1+30°C-Ac3+30°C, and more preferably in a temperature range of 800-880°C.
When the temperature during continuous annealing is lower than Ac1+30°C, recrystallization may not be sufficient, and it may be be difficult to form sufficient austenite, and thus, the target-level fractions of tempered martensite, fresh martensite, and bainite after annealing may not be ensured. When the temperature exceeds Ac3+30°C, the austenite grain size may become coarse, and the fine residual austenite phase may not be evenly formed around the hard phase. Also, the productivity may decrease, and the formation of surface concentrated substances may be intensified due to elements decreasing wettability of hot-dip galvanized such as Si, Mn, and B due to high-temperature annealing, the plating surface quality may not be ensured.

Sequential cooling

According to the above, it may be preferable to sequentially cool the continuously annealed cold-rolled steel sheet.
Specifically, the cooling may be preferably performed at an average cooling rate of 10°C/s or lower (excluding 0°C/s) to 450-670°C (the cooling in this case is referred to as primary-cooling), and at an average cooling rate of 5°C/s or higher to 250-500°C (the cooling in this case is referred to as secondary-cooling).

Primary-cooling

The present disclosure may be required to ensure the fractions of martensite and bainite in the subsequent secondary-cooling process in order to form 40% or more of one or more of fractions of tempered martensite and bainite as the final structure. To this end, the primary-cooling may be performed at an average cooling rate of 10°C/s or lower (excluding 0°C/s) to 450-670°C.
Specifically, when the subsequent secondary-cooling is terminated below Ms (martensite transformation initiation temperature), a relatively large amount of martensite phase may be formed, and to this end, it may be preferable to control the termination temperature of the primary-cooling to be as low as possible. Also, when the subsequent secondary-cooling is terminated in the bainite temperature region, a bainite phase may be formed relatively preferably, and to this end, it may be preferable to control the termination temperature of the primary-cooling to be higher.
Cooling during primary-cooling may be performed at an average cooling rate of 10°C/s or lower (excluding 0°C/s) up to 450-700°C, and preferably, when subsequent secondary-cooling is terminated below Ms, it may be preferable to perform the primary-cooling within a temperature range of 450-600°C, and when subsequent secondary-cooling is terminated in the bainite temperature region, it may be preferable to perform the primary-cooling within a temperature range of 550-700°C.
When the termination temperature during primary-cooling is less than 450°C, the facility cooling the atmosphere gas in the annealing furnace may be heavily loaded, and the cooling rate may increase such that the ferrite phase formed during cooling may not be sufficiently ensured. When the termination temperature exceeds 700°C, an excessively high cooling speed may be required during subsequent cooling (secondary-cooling), which may be disadvantageous.
Also, when the average cooling rate during the primary-cooling exceeds 10°C/s, carbon diffusion may not occur sufficiently. Considering productivity, the primary-cooling may be performed at an average cooling rate of 1°C/s or more.

Secondary-cooling

After completing the primary-cooling under the conditions described above, it may be preferable to perform secondary-cooling, and in this case, the target microstructure may be formed by controlling the cooling termination temperature and the cooling rate.
When the secondary-cooling is performed below Ms, quenching martensite may be formed, and as the temperature decreases, the fraction of quenching martensite may increase, which may induce improvement in the strength of the steel sheet. Also, during the subsequent heat treatment (the reheating process of the present disclosure), martensite may be tempered and may become tempered martensite, the supersaturated carbon in the martensite may be distributed to the surrounding untransformed austenite, such that stability of the residual austenite may increase and ductility may improve.
When cooling is performed at a temperature exceeding Ms during the secondary-cooling, the fraction of bainite may be increased. In this case, as the precipitation of carbides is delayed by the effects of Si and Al during the bainite transformation process, carbon may be distributed from the bainite to the surrounding untransformed austenite, stability of the residual austenite may be increased and ductility may be improved.
When the termination temperature of the secondary-cooling is less than 250°C, the fraction of quenching martensite may increase excessively, such that the fraction of the residual austenite phase may decrease and the shape of the steel sheet may be deteriorated. When the temperature exceeds 500°C, bainite may not be sufficiently formed, such that the fraction of the residual austenite phase may decrease, and the fraction of the fresh martensite phase may increase significantly in the subsequent process, which may cause an excessive increase in strength.
Also, when the average cooling rate during the secondary-cooling is less than 5°C/s, the pearlite phase may be formed, such that the bainite phase may not be formed to the target level. An upper limit of the average cooling rate may not be particularly limited, and may be appropriately selected by a general engineer in consideration of the specifications of the cooling facility. For example, the cooling may be performed at 100°C/s or lower.
Also, the secondary-cooling may use a hydrogen cooling facility using hydrogen gas (H₂ gas). In this way, by performing cooling using a hydrogen cooling facility, the effect of suppressing surface oxidation occurring during the secondary-cooling may be obtained. In this case, the type of gas used in the hydrogen cooling facility is not limited, but for example, the gas may be controlled with 60-70% of hydrogen (H₂) and the remainder of nitrogen (N₂).
When performing the cooling sequentially as described above, the cooling rate during the secondary-cooling may be higher than the cooling rate during the primary-cooling.

Holding

After completing the secondary-cooling as described above, a process of holding 30 seconds or more in the cooled temperature range may be performed.
Through the holding process, the effect of tempering martensite or further increasing the amount of bainite transformation may be obtained. When the holding time is less than 30 seconds, it may be difficult to expect the above-described effect.

Reheating and holding

The microstructure intended in the present disclosure may be formed through the process of reheating and holding the cold-rolled steel sheet having gone through the sequential cooling as described above. Specifically, it may be preferable to perform a process of reheating the secondary-cooled cold-rolled steel sheet to a temperature of 490°C or lower and holding the steel sheet for 20 seconds or more.
By reheating at the above-described temperature and holding, the quenching martensite generated in the previous cooling process may be transformed into tempered martensite, and also bainite transformation may be accompanied.
During the tempering process, carbon supersaturated in martensite may be redistributed to the surrounding untransformed austenite. Also, in the case in which the secondary-cooling is terminated when exceeding Ms, the fraction of bainite may increase significantly during the reheating and holding process. During this process, the carbon released from bainite may be redistributed to untransformed austenite, such that the stability of residual austenite may improve and ductility may increase.
However, when the temperature during the reheating is excessively high, carbides in tempered martensite and bainite may become coarse, such that strength may decrease, and the carbon redistribution effect to untransformed austenite due to the formation of coarse carbides may decrease, the residual austenite fraction may decrease and it may be difficult to expect an improvement in ductility.
Accordingly, the reheating temperature may be performed at 490°C or lower, and more preferably at 470°C or lower.
As described above, it may be preferable to sufficiently implement the above-described effect by holding the cold-rolled steel sheet reheated to 490°C or lower at that temperature for 20 seconds or more. However, when the holding time is excessive and exceeds 5 minutes, the tempering effect of martensite may become excessive, which decrease the strength, such that it may be preferable not to exceed 5 minutes.
In the present disclosure, a plated steel sheet may be provided by plating a cold-rolled steel sheet manufactured as the above.

Hot-dip galvanized

It may be preferable to manufacture a hot-dip galvanized steel sheet by immersing the steel sheet in a hot-dip galvanizing bath after undergoing the reheating and holding process as the above.
In this case, the hot-dip galvanized may be performed under normal conditions, but may be performed at a temperature range of 430-490°C, for example. Also, the composition of the hot-dip galvanizing bath is not particularly limited during the hot-dip galvanizing, and may be a pure zinc plating bath or a zinc-based alloy plating bath including Si, Al, Mg, or the like.

Alloying heat treatment

Also, if necessary, the alloyed hot-dip galvanized steel sheet may be obtained by performing alloying heat treatment on the hot-dip galvanized steel sheet. In the present disclosure, the alloying heat treatment process conditions may not be particularly limited, and any normal conditions may be used. For example, the alloying heat treatment process may be performed at a temperature range of 480-600°C.
Also, the final cooling and the skin pass process may be performed after the hot-dip galvanizing or the alloying heat treatment.

Final cooling

The steel sheet having gone through the hot-dip galvanizing or the alloying heat treatment as above may be finally cooled and fresh martensite may be further formed. In this case, the final cooling may be preferably performed at a cooling rate of 5°C/s or higher to a temperature of Ms or lower.
When the cooling rate is less than 5°C/s during the cooling process, the fresh martensite phase may not be ensured to the intended level during the cooling process. The upper limit of the cooling rate is not particularly limited, but the cooing may be performed at 50°C/s or lower to form a predetermined fraction of the pressed martensite phase.

Skin pass rolling

Furthermore, if necessary, by performing a skin pass rolling on the finally cooled hot-dip galvanized steel sheet or alloyed hot-dip galvanized steel sheet, a large amount of dislocations may be formed in the steel, thereby further improving bake hardenability. In this case, it may be preferable that the reduction ratio be less than 2% (excluding 0%). When the reduction ratio is 2% or more, the ratio may be preferable in terms of dislocation formation, but side effects such as plate breakage may occur due to facility capacity limitations.

Mode for Invention

Hereinafter, preferable embodiments of the present disclosure will be described with reference to the accompanied drawings. However, the embodiment of the present disclosure may be modified with various other forms, and the scope of the present disclosure is not limited to the embodiment described below. The embodiments below are intended to help understand the present disclosure, and the scope of rights in the present disclosure should not be limited to the following embodiment, and may be determined not only by the claims described below but also by equivalents thereof.

(Embodiment 1)

Steel slabs having the alloy compositions in Table 1 below were heated at a temperature of 1200°C, and each of the heated slabs was finishing hot-rolled at 900°C, which is higher than the Ar3 transformation point temperature, and hot-rolled steel sheets were manufactured. The steel sheets were pickled, each hot-rolled steel sheet was wound at the temperature disclosed in Table 2 and cooled to room temperature at a cooling rate of 0.1°C/s or lower. Thereafter, the cooled hot-rolled steel sheet was cold-rolled and cold-rolled steel sheets was manufactured.
During the cold-rolling, the cumulative reduction ratio of stands 1-3 was determined to be 25%, and the overall reduction ratio was determined to be 60%.
Thereafter, each cold-rolled steel sheet was continuously annealed under the conditions listed in Table 2, and primary-cooling, secondary-cooling, holding, and reheating were performed. After the secondary-cooling and holding processes were completed, reheating to a temperature of 490°C or lower was performed, and holding at the above temperature for 30 seconds was performed. The holding process after the secondary-cooling was performed for 30 seconds.

Thereafter, zinc plating was performed in a hot-dip galvanizing bath at 460°C, and final cooling to room temperature at a cooling rate of 5°C/s was performed, and skin pass was performed at less than 2%, thereby manufacturing a hot-dip galvanized steel sheet. In this case, alloying heat treatment was performed after the zinc plating for some steels.

[Table 1]

Classification	Composition (wt%)											Relational expression 1
Classification	C	Si	Mn	P	S	Sol-Al	Nb	Ti	B	Cr	Mo	Relational expression 1
Steel type 1	0.15	1.0	2.25	0.008	0.002	0.030	0.00	0.020	0.0011	0.70	0.000	0.80
Steel type 2	0.18	0.6	2.10	0.006	0.002	0.230	0.00	0.000	0.0015	0.81	0.010	0.83
Steel type 3	0.14	1.2	2.50	0.007	0.002	0.310	0.00	0.015	0.0013	0.51	0.020	0.90
Steel type 4	0.13	1.1	2.80	0.008	0.003	0.430	0.00	0.000	0.0010	0.40	0.000	0.85
Steel type 5	0.11	1.0	2.10	0.005	0.003	0.030	0.02	0.000	0.0012	0.65	0.000	0.72
Steel type 6	0.17	0.8	2.50	0.004	0.002	0.230	0.00	0.000	0.0000	0.80	0.020	0.76
Steel type 7	0.16	0.3	2.55	0.013	0.015	0.341	0.01	0.020	0.0011	0.30	0.152	0.67
Steel type 8	0.11	0.6	2.58	0.013	0.002	0.032	0.02	0.016	0.0000	0.20	0.000	0.61
Steel type 9	0.14	0.6	2.50	0.010	0.002	0.035	0.02	0.040	0.0021	0.25	0.200	0.60
Steel type 10	0.13	0.6	2.58	0.012	0.002	0.051	0.00	0.021	0.0000	0.59	0.103	0.54
Steel type 11	0.12	0.5	2.40	0.010	0.001	0.025	0.03	0.015	0.0015	0.20	0.400	0.46

Thereafter, the mechanical properties and microstructure properties of each sample were evaluated, and the results are listed in Table 3 below. In this case, the tensile test for each test sample was performed in the L direction using the DIN standard and the tensile properties were evaluated. In this case, during the tensile test of the material, in the load-displacement diagram, the strain to the breakage was defined as total elongation (T-El), the strain to the maximum load point was defined as uniform elongation (U-El), and the strain from the maximum load point to the breakage was defined as post uniform elongation (P-El), and values thereof were listed. As for the microstructure fraction, at the 1/4t point of the plate thickness of the annealed steel sheet, the base structure was nital-corroded, and the fractions of tempered martensite (TM), bainite (B), ferrite (F), fresh martensite (FM), and residual austenite (A) were measured using FE-SEM, image analyzer, EBSD, and XRD.

[Table 2]

Steel type	Hot-rolling	Annealing				Classification
Steel type	Winding temperature (°C)	Annealing temperature (°C)	Slowcooling temperature (°C)	Rapid cooling temperature (°C)	Reheating temperature (°C)	Classification
Steel type 1	530	870	500	300	460	Inventive steel 1
Steel type 2	520	850	600	340	460	Inventive steel 2
Steel type 3	560	840	550	450	460	Inventive steel 3
Steel type 4	500	850	600	480	460	Inventive steel 4
Steel type 5	480	860	500	330	460	Inventive steel 5
Steel type 6	650	850	570	300	460	Inventive steel 6
Steel type 7	760	840	650	300	480	Comparative steel 1
Steel type 8	540	850	650	600	450	Comparative steel 2
Steel type 9	530	770	650	350	460	Comparative steel 3
Steel type 10	620	850	680	230	430	Comparative steel 4
Steel type 11	350	830	720	400	460	Comparative steel 5
Steel type 1	530	770	650	600	480	Comparative steel 6
Steel type 2	520	920	650	150	460	Comparative steel 7
Steel type 5	480	910	750	650	470	Comparative steel 8

[Table 3]

Classifica tion	Phase fraction						Mechanical properties							Classific ation
Classifica tion	F (%)	TM (%)	B (%)	FM (%)	A (%)	Relational expression 2	YS (MPa)	TS (MPa)	T-El (%)	U-El (%)	P-El (%)	Relational expression 3	Relational expression 4	Classific ation
Steel type 1	24	50	11	7	8	91	788	1026	16. 5	10.0	5	7.7	5.0	Inventive steel 1
Steel type 2	28	43	9	11	9	88	730	1031	16.9	10.3	6.6	7.3	4.7	Inventive steel 2
Steel type 3	30	0	46	11	13	87	702	1018	19.2	11.5	7.7	7.9	5.3	Inventive steel 3
Steel type 4	32	0	45	14	9	83	709	1033	18.1	11.0	7.0	7.6	4.8	Inventive steel 4
Steel type 5	25	47	13	10	5	85	761	1075	14.1	9.1	5.0	6.4	3.5	Inventive steel 5
Steel type 6	12	65	10	5	8	93	835	1100	16.1	10.1	5.5	7.7	4.2	Inventive steel 6
Steel type 7	11	50	14	23	2	70	800	1156	12.0	8.2	3.8	5.7	2.6	Comparati ve steel 1
Steel type 8	45	0	30	22	3	73	525	1018	15.7	10.5	5.2	5.4	2.7	Comparati ve steel 2
Steel type 9	44	29	2	22	3	72	591	1039	14.0	10.1	3.9	5.7	2.2	Comparati ve steel 3
Steel type 10	14	60	3	21	2	74	830	1173	10. 8	7.0	3.8	5.0	2.7	Comparati ve steel 4
Steel type 11	40	0	30	28	2	67	631	1084	11.2	7.3	3.9	4.2	2.3	Comparati ve steel 5
Steel type 1	65	0	8	27	0	65	560	950	13.1	9.1	4	5.4	2.4	Comparati ve steel 6
Steel type 2	3	90	0	6	1	98	990	1145	9.2	6.1	3.1	5.3	2.7	Comparati ve steel 7
Steel type 5	10	0	18	61	1	20	918	1230	6.3	3.1	3.2	2.3	2.4	Comparati ve steel 8

In Tables 1 to 3, relational expressions 1 to 4 are as below. $(10 C + Si + Al) / (Mn + 1.3 Cr + 2.7 Mo) \geq 0.7$
(Herein, each element indicates a weight content) $({FM}_{TM + B} / {FM}_{T}) \times 100 \geq 80 %$
(Herein, FM_T indicates a total fraction of fresh martensite, and FM_TM+B indicates a fraction of fresh martensite in contact with tempered martensite or bainite in the fraction of fresh martensite)

$YS \times U - El / TS \geq 6$

$YS \times P - El / TS \geq 3$
FIG. 1 is a graph indicating the relationship between relational expression 1 and relational expression 3 for inventive steels 1 to 6 and comparative steels 1 to 5 in the embodiment, and FIG. 2 may be a graph indicating the relationship between relational expression 1 and relational expression 4 for inventive steel 1 to 6 and comparative steel 1 to 5. When the steel composition relational expression 1 of relational expression 1 is 0.7 or higher, relational expressions 3 and 4 were satisfied. As described above, in order to improve formability and impact fracture resistance in the present disclosure, relational expression 3 and relational expression 4 may need to be satisfied. To this end, it may be effective to ensure the microstructure presented in the present disclosure, and to this end, it may be important to enrich C in untransformed austenite. In order to enrich C in untransformed austenite, a predetermined amount of C content may need to be included in the tempering composition, and additionally, the formation of carbides may need to be delayed in the bainite transformation region by adding Si and Al. Also, proper ferrite may need to be ensured to sufficiently enrich C in austenite, but when hardenability elements such as Mn, Cr, and Mo are added excessively, the formation of ferrite may be suppressed, such that C may not be sufficiently enriched in austenite.
When relational expression 1 is 0.7 or more, the fractions of the target ferrite and residual austenite may be ensured, and accordingly, fresh martensite may also be finely dispersed, thereby increasing uniform elongation and non-uniform ductility, and as a result, relational expressions 3 and 4 may also be satisfied.
When the relational expression 1 is less than 0.7, the target residual austenite and ferrite fraction may not be ensured, and as a result, the fresh martensite may not be finely dispersed, such that the uniform elongation and post uniform elongation may decrease, and relational expression 3 and relational expression 4 may not be satisfied.
From the Tables 1 to 3, it may be confirmed that inventive steels 1 to 6, which satisfied the alloy compositions and manufacturing conditions suggested in the present disclosure, had tensile strength of 980 MPa or more and satisfied relational expressions 3 and 4. Accordingly, excellent formability and fracture resistance were ensured.
In contrast, in comparative steels 1 to 8, which deviated from the composition range or relational expression 1 suggested in the present disclosure or did not satisfy the manufacturing conditions, the intended microstructure was not formed, and the intended physical properties in the present disclosure were not ensured.

Claims

A steel sheet, comprising:
by weight%, carbon (C): 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn) : 1.9-3.0%, molybdenum (Mo) : 0.3% or less, chromium (Cr) : 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities,

wherein C, Si, Al, Mn, Cr and Mo satisfy relational expression 1 as below,

wherein a microstructure comprises, by area%, ferrite: 10-35%, residual austenite: 3-15%, fresh martensite: 20% or less (excluding 0%), and a remainder of one of tempered martensite and bainite, and

wherein, in fresh martensite, a fraction of fresh martensite in which a distance between fresh martensite phases is three or more times a grain size of fresh martensite is 30% or more: $(10 C + Si + Al) / (Mn + 1.3 Cr + 2.7 Mo) \geq 0.7$

where each element indicates a weight content.
The steel sheet of claim 1, wherein a fraction of fresh martensite satisfies relational expression 2 as below: $({FM}_{TM + B} / {FM}_{T}) \times 100 \geq 80 %$ where FM_T indicates a total fraction of fresh martensite, and FM_TM+B indicates a fraction of fresh martensite in contact with tempered martensite or bainite in the fraction of fresh martensite.
The steel sheet of claim 1, wherein yield strength (YS), uniform elongation (U-El) and tensile strength (TS) of the steel sheet satisfies relational expression 3: $YS \times U - El / TS \geq 6$
The steel sheet of claim 1, wherein yield strength (YS), post uniform elongation (P-El) and tensile strength (TS) of the steel sheet satisfies relational expression 4 as below: $YS \times P - El / TS \geq 3$
The steel sheet of claim 1, wherein the steel sheet comprises boron (B): 0.01% or less.
The steel sheet of claim 1, wherein the steel sheet further comprises one or more of titanium (Ti): 0.05% or less and niobium (Nb): 0.05% or less.
The steel sheet of claim 1, wherein the steel sheet further comprises a zinc-based plating layer.
A method of manufacturing a steel sheet, the method comprising:
heating a steel slab including, by weight%, carbon (C) : 0.1-0.2%, silicon (Si): 0.5-1.3%, aluminum (Al): 0.5% or less (excluding 0%), manganese (Mn): 1.9-3.0%, molybdenum (Mo) : 0.3% or less, chromium (Cr) : 1% or less (excluding 0%), phosphorus (P): 0.1% or less, sulfur (S): 0.1% or less, and a balance of Fe and inevitable impurities, wherein C, Si, Al, Mn, Cr and Mo satisfy relational expression 1 as below, in a temperature range of 1100-1300°C;

manufacturing a hot-rolled steel sheet by finishing hot-rolling the reheated steel slab at a temperature range of Ar3 or higher;

winding the hot-rolled steel sheet in a temperature range of 400-700°C;

cooling the hot-rolled steel sheet to room temperature at a cooling rate of 0.1°C/s after the winding;

manufacturing a cold-rolled steel sheet by cold-rolling the hot-rolled steel sheet at a total reduction ratio of 30-80% after the cooling;

continuously annealing the cold-rolled steel sheet at a temperature of Ac1+30°C-Ac3+30°C;

primary-cooling the cold-rolled steel sheet having gone through the continuous annealing to 450-700°C at a cooling rate of 10°C/s or lower;

secondary-cooling the steel sheet to 250-500°C at a cooling rate of 5°C/s or more after the primary-cooling; and

reheating the secondary-cooled cold-rolled steel sheet at a temperature of 490°C or lower and holding the steel sheet for 20 seconds or more,

wherein the cold-rolling is performed with a cumulative reduction ratio of 25% or more of first 1-3 stands. $(10 C + Si + Al) / (Mn + 1.3 Cr + 2.7 Mo) \geq 0.7$

where each element indicates a weight content.
The method of claim 8, wherein the continuous annealing is performed at a temperature of 800-880°C.
The method of claim 8, wherein the cooling rate during the secondary-cooling is higher than the cooling rate during the primary-cooling.
The method of claim 8, wherein the secondary-cooling is performed in a hydrogen quenching facility using hydrogen (H₂) gas.
The method of claim 8, further comprising:
holding the steel sheet for 30 seconds after the secondary-cooling.
The method of claim 8, further comprising:
hot-dip galvanizing the steel sheet in a plating bath at 430-490°C after the reheating and holding.
The method of claim 13, further comprising:
performing alloying heat treatment after the hot-dip galvanizing.
The method of claim 14, further comprising:
cooling the steel sheet to a temperature of Ms-100°C or lower at a cooling rate of 5°C/s or higher and performing a skin pass of less than 2% after the alloying heat treatment.