CN113825852B - Steel sheet having high strength and high formability and method for manufacturing same - Google Patents

Steel sheet having high strength and high formability and method for manufacturing same Download PDF

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CN113825852B
CN113825852B CN202080035734.7A CN202080035734A CN113825852B CN 113825852 B CN113825852 B CN 113825852B CN 202080035734 A CN202080035734 A CN 202080035734A CN 113825852 B CN113825852 B CN 113825852B
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steel sheet
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temperature
rolled steel
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CN113825852A (en
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严浩溶
具南勋
金愍城
吴圭真
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Hyundai Steel Co
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

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  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
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  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Sheet Steel (AREA)
  • Heat Treatment Of Steel (AREA)

Abstract

A steel sheet having high strength and formability according to an aspect of the present invention includes, in wt%, 0.05% to 0.15% of carbon (C), more than 0% to 0.4% or less of silicon (Si), 4.0% to 9.0% of manganese (Mn), more than 0% to 0.3% or less of aluminum (Al), 0.02% or less of phosphorus (P), 0.005% or less of sulfur (S), 0.006% or less of nitrogen (N), and the balance of iron (Fe) and other unavoidable impurities. The steel sheet has a microstructure consisting of ferrite and retained austenite. The grain size of the microstructure is 3 μm or less. The steel sheet has a Yield Strength (YS) of 800MPa or more, a Tensile Strength (TS) of 980MPa or more, an Elongation (EL) of 25% or more, and a Hole Expansion Ratio (HER) of 20% or more.

Description

Steel sheet having high strength and high formability and method for manufacturing same
Technical Field
The present invention relates to a steel sheet and a method for manufacturing the same, and more particularly, to a steel sheet having high strength and high formability and a method for manufacturing the same.
Background
In recent years, the strength of automobile steel sheets has been rapidly increased from the viewpoint of safety and weight reduction of automobiles. In order to ensure the safety of passengers, steel plates for automotive structural members need to have sufficient impact toughness by increasing the strength or thickness thereof. Further, these steel sheets need to have sufficient formability to be applied to automobile parts, and in order to improve fuel efficiency of automobiles, it is necessary to reduce the weight of automobile bodies. Therefore, research has been conducted to continuously significantly enhance the automobile steel sheet and improve the formability thereof.
Currently, as high-strength steel sheets for automobiles having the above-described characteristics, dual-phase steels having strength and elongation secured by two phases (ferrite phase and martensite phase) and transformation induced plasticity steels having strength and elongation secured by phase transformation of retained austenite in the final structure during plastic deformation have been proposed.
The related art includes korean patent application No. 10-2016-0077463 (titled "ultra-high strength, high ductility steel sheet having excellent yield strength and method for manufacturing the same").
Disclosure of Invention
Technical problem
The present invention is directed to provide a steel sheet having high formability and high strength and a method for manufacturing the same.
Technical scheme
In one aspect of the present invention, there is provided a steel sheet having high strength and high formability, which includes, in wt%, carbon (C) in an amount of 0.05% to 0.15%, silicon (Si) in an amount of greater than 0% and less than or equal to 0.4%, manganese (Mn) in an amount of 4.0% to 9.0%, aluminum (Al) in an amount of greater than 0% and less than or equal to 0.3%, phosphorus (P) in an amount of 0.02% or less, sulfur (S) in an amount of 0.005% or less, nitrogen (N) in an amount of 0.006% or less, and iron (Fe) and other unavoidable impurities as a balance. The steel sheet has a microstructure consisting of ferrite and retained austenite. The microstructure has a grain size of 3 μm or less. The steel sheet has a Yield Strength (YS) of 800MPa or more, a Tensile Strength (TS) of 980MPa or more, an Elongation (EL) of 25% or more, and a Hole Expansion Ratio (HER) of 20% or more.
In an exemplary embodiment, the steel sheet may further include one or more elements of niobium (Nb), titanium (Ti), vanadium (V), and molybdenum (Mo), and a content of each of the elements may be greater than 0% and less than or equal to 0.02 wt%.
In an exemplary embodiment, the steel sheet may further include boron (B) in an amount of greater than 0% and less than or equal to 0.001 wt%.
In exemplary embodiments, the volume fraction of retained austenite in the microstructure may be 10 to 30 volume%.
In one aspect of the present invention, there is provided a method of manufacturing a steel sheet having high strength and high formability, the method including the steps of: (a) Manufacturing a hot-rolled steel sheet from a steel slab including, in wt.%, carbon (C) in an amount of 0.05% to 0.15%, silicon (Si) in an amount of greater than 0% and less than or equal to 0.4%, manganese (Mn) in an amount of 4.0% to 9.0%, aluminum (Al) in an amount of greater than 0% and less than or equal to 0.3%, phosphorus (P) in an amount of 0.02% or less, sulfur (S) in an amount of 0.005% or less, nitrogen (N) in an amount of 0.006% or less, and iron (Fe) and other unavoidable impurities as a remainder; (b) manufacturing a cold-rolled steel sheet by cold-rolling the hot-rolled steel sheet; (c) Subjecting the cold-rolled steel sheet to a first heat treatment at a temperature of AC3 to (AC 3+ 15) ° C; and (d) performing a second heat treatment at the critical zone temperature on the cold-rolled steel sheet subjected to the first heat treatment. The cold-rolled steel sheet after the step (d) has a microstructure consisting of ferrite and retained austenite.
In an exemplary embodiment, the steel slab may further include one or more elements of niobium (Nb), titanium (Ti), vanadium (V), and molybdenum (Mo), and a content of each of the elements may be greater than 0% and less than or equal to 0.02 wt%.
In an exemplary embodiment, the steel slab may further include boron (B) in an amount greater than 0% and less than or equal to 0.001 wt%.
In exemplary embodiments, the step (c) may include the step of cooling the heat-treated cold-rolled steel sheet to a temperature of 350 ℃ to 450 ℃ at a cooling rate of 4 ℃/s to 10 ℃/s.
In an exemplary embodiment, the step (d) may include a step of cooling the heat-treated cold rolled steel sheet to a temperature of 350 ℃ to 450 ℃ at 4 ℃/s to 10 ℃/s.
In an exemplary embodiment, step (a) may include the steps of: (a 1) reheating a steel slab to a temperature of 1150 ℃ to 1250 ℃; (a2) Hot rolling the reheated slab to a finish rolling exit temperature of 925 ℃ to 975 ℃; and (a 3) cooling the hot rolled steel sheet to a temperature of 700 ℃ to 800 ℃ at a cooling rate of 10 ℃/s to 30 ℃/s, followed by winding.
In an exemplary embodiment, between steps (a) and (b), the method may further include the step of performing a softening heat treatment on the hot rolled steel sheet at a temperature of 550 ℃ to 650 ℃.
In exemplary embodiments, the cold-rolled steel sheet after step (d) may have a Yield Strength (YS) of 800MPa or more, a Tensile Strength (TS) of 980MPa or more, an Elongation (EL) of 25% or more, and a Hole Expansion Ratio (HER) of 20% or more.
In exemplary embodiments, the cold rolled steel sheet after step (d) may have a grain size of 3 μm or less.
Advantageous effects
According to the present invention, a steel sheet having a microstructure composed of ultra-fine grained ferrite and retained austenite can be manufactured through component system control and process condition control. The steel sheet may have high strength due to ferrite having fine grains, and high strength and elongation due to the presence of retained austenite in an amount of 10 to 30 vol% in the microstructure. In addition, since the shape of the microstructure is controlled, the steel sheet may have a high Hole Expansion Ratio (HER). Therefore, a steel sheet having high formability and high strength can be effectively obtained.
Drawings
Fig. 1 is a process flow diagram schematically showing a method of manufacturing a steel sheet having high strength and high formability according to an exemplary embodiment of the present invention.
Figure 2 shows the results of high temperature tensile testing of a sample of a comparative component system of the present invention.
FIG. 3 shows the results of high temperature tensile testing of samples of the embodied component systems of the present invention.
Fig. 4 is a photograph showing the microstructure of a high-strength steel sheet according to an exemplary embodiment of the present invention.
Detailed Description
The present invention will be described in detail below with reference to the accompanying drawings so that those skilled in the art to which the present invention pertains can easily carry out the present invention. The present invention may be embodied in many different forms and is not limited to the embodiments described herein. In the present specification, the same reference numerals denote the same or similar components. In addition, when known functions and configurations may unnecessarily obscure the subject matter of the present invention, a detailed description thereof will be omitted.
According to an exemplary embodiment of the present invention, a steel sheet having high strength and high formability may have a final microstructure consisting of fine-grained ferrite and residual austenite present in an amount of 10 to 30 vol%. Accordingly, the steel sheet may have high strength, high elongation, and high Hole Expansion Ratio (HER).
First, in order to impart high elongation to the steel sheet, the steel sheet sufficiently contains retained austenite at a level of 10 to 30 vol%. The retained austenite can increase the elongation of the steel sheet in substantially the same manner as the conventional transformation induced plasticity steel. In order to secure the fraction of retained austenite, austenite stabilizing elements may be appropriately added to the steel sheet as described later. Further, as described later, the first annealing heat treatment and the second annealing heat treatment may be continuously performed, and the second annealing heat treatment may be performed at the critical zone temperature.
Then, in order to make the steel sheet have a high hole expansion ratio, a phase boundary between the hard phase and the soft phase, which may serve as a crack formation site in the steel sheet, is reduced. For this reason, the steel sheet may not include a hard phase, such as martensite and bainite, in its final microstructure. In addition, in order to make the steel sheet have a high hole expansion ratio, interfaces between precipitates and crystal grains are reduced. For this reason, the contents of the precipitate forming element (e.g., titanium, niobium, and vanadium) and the precipitate growth inhibiting element (e.g., molybdenum) can be controlled. In addition, in order to provide a steel sheet with a high hole expansion ratio, the fraction of High Angle Grain Boundaries (HAGB) in the final structure may be increased. As an example, the high angle grain boundary may mean a grain boundary in which an angle between adjacent grains is 15 ° or more. In addition, the shape of the microstructure can be controlled to give a high hole expansion ratio to the steel sheet. As described later, in order to increase the fraction of high angle grain boundaries and control the shape of the microstructure, the annealing heat treatment may be performed in two steps (including a first heat treatment and a second heat treatment).
Then, in order to impart high strength to the steel sheet, the crystal grains of the final microstructure are refined. By the above annealing heat treatment performed in two steps, the grain sizes of ferrite and retained austenite can be controlled to 3 μm or less. Further, the first annealing heat treatment may be performed at a temperature of AC3 to (AC 3+ 15) ° c.
Hereinafter, a steel sheet having high formability and high strength according to an exemplary embodiment of the present invention having the above-described characteristics will be described in more detail.
Steel sheet having high strength and high formability
The high-strength steel sheet according to an exemplary embodiment of the present invention includes, in wt%, carbon (C) in an amount of 0.05% to 0.15%, silicon (Si) in an amount of greater than 0% and less than or equal to 0.4%, manganese (Mn) in an amount of 4.0% to 9.0%, aluminum (Al) in an amount of greater than 0% and less than or equal to 0.3%, phosphorus (P) in an amount of 0.02% or less, sulfur (S) in an amount of 0.005% or less, nitrogen (N) in an amount of 0.006% or less, and the balance iron (Fe) and other unavoidable impurities. In addition, the high strength steel sheet further includes one or more elements of niobium (Nb), titanium (Ti), vanadium (V), and molybdenum (Mo), and the content of each of the elements may be greater than 0% and less than or equal to 0.02 wt%. Further, the high-strength steel sheet may further include more than 0% and 0.001% by weight or less of boron (B).
The effects and contents of various components included in the high strength cold rolled steel sheet according to an exemplary embodiment of the present invention will be described in detail below (the contents of the various components are given in wt% based on the total weight of the steel sheet and are hereinafter represented by%).
Carbon (C): 0.05 to 0.15 percent
Carbon (C) is the most important alloying element in steel making and the main objective in the present invention is to provide substantial strengthening and stabilizing of austenite. The high carbon (C) concentration in austenite improves austenite stability, so that it is easy to ensure proper austenite to improve material properties. However, an excessively high carbon (C) content may cause a decrease in weldability due to an increase in carbon equivalent, and a large amount of precipitated cementite structure (e.g., pearlite) may be formed during cooling. For this reason, it is preferable to add carbon (C) in an amount of 0.05 to 0.15% by weight based on the total weight of the steel sheet. If the carbon content is less than 0.05%, it may be difficult to secure the strength of the steel sheet, and when the carbon content is more than 0.15%, the toughness and ductility of the steel sheet may be deteriorated.
Silicon (Si): more than 0 and less than or equal to 0.4 percent
Silicon (Si) is an element that suppresses carbide formation in ferrite and increases the diffusion rate of austenite by increasing the activity of carbon (C). Silicon (Si) is also a well-known ferrite stabilizing element that increases ductility by increasing the fraction of ferrite during cooling. In addition, silicon has a very high ability to inhibit carbide formation, and thus is an essential element to ensure the TRIP effect by increasing the carbon concentration in the residual austenite during bainite formation. However, if silicon (Si) is added in an amount greater than 0.4%, silicon dioxide (SiO) may be formed on the surface of the steel sheet during machining 2 ) The rolling load in the hot rolling process is increased, and a large amount of red scale is generated. Therefore, it is preferable to add silicon (Si) in an amount of 0.4% or less by total weight of the steel sheet.
Manganese (Mn): 4.0 to 9.0 percent
Manganese (Mn) is an austenite stabilizing element. With the addition of manganese (Mn), ms (martensite formation start temperature) gradually decreases, thereby exhibiting the effect of increasing the fraction of retained austenite after heat treatment.
Manganese is contained in an amount of 4.0% to 9.0% by weight of the total steel sheet. If manganese is added in an amount of less than 4.0%, the above effects cannot be sufficiently ensured. On the other hand, if manganese is added in an amount of more than 9.0%, weldability may be reduced due to an increase in carbon equivalent, and manganese oxide (MnO) may be formed on the surface of the steel sheet during the process, resulting in a reduction in platability due to a reduction in wettability of the corresponding portion.
Aluminum (Al): more than 0 and less than or equal to 0.3 percent
Like silicon (Si), aluminum (Al) is considered as an element that stabilizes ferrite and suppresses the formation of carbides. In addition, aluminum has an effect of increasing the equilibrium temperature, and thus when aluminum (Al) is added, there is an advantage in that a suitable heat treatment temperature range is widened. However, if aluminum is excessively added in an amount of more than 0.3%, problems may occur in continuous casting due to AIN precipitation. Therefore, aluminum may be added in an amount greater than 0% and less than or equal to 0.3% of the total weight of the steel sheet.
At least one of niobium (Nb), titanium (Ti), vanadium (V), and molybdenum (Mo): each of which is more than 0% and less than or equal to 0.2%
Niobium (Nb), titanium (Ti), vanadium (V) and molybdenum (Mo) may optionally be included in the steel. First, niobium (Nb), titanium (Ti), and vanadium (V) are elements precipitated in the form of carbides in steel, and are added to secure strength by carbide precipitation. Titanium (Ti) can suppress crack formation during continuous casting by suppressing the formation of AIN. However, if niobium (Nb), titanium (Ti), and vanadium (V) are added in amounts greater than 0.2%, respectively, it may form coarse precipitates, which have disadvantages in that the amount of carbon in the steel is reduced, material properties are deteriorated, and manufacturing costs are increased due to the addition of niobium (Nb), titanium (Ti), and vanadium (V). In addition, if titanium is excessively added, clogging of the nozzle may be caused during continuous casting. Therefore, when at least one of niobium (Nb), titanium (Ti), and vanadium (V) is added, each of niobium (Nb), titanium (Ti), and vanadium (V) may be added in an amount of greater than 0% and less than or equal to 0.2% of the total weight of the steel sheet, respectively.
In addition, molybdenum (Mo) may be used to control the size of carbides by inhibiting the growth of carbides. However, if molybdenum is added in an amount of more than 0.2%, there are disadvantages in that the above-mentioned effects are saturated and the manufacturing cost is increased.
Boron (B)
Boron (B) may be optionally added to the steel sheet, and may act as a grain boundary strengthening element. Boron may be added in an amount of more than 0% by total weight of the steel sheet and less than or equal to 0.001% by total weight of the steel sheet. If boron is added in an amount greater than 0.001%, it is possible to reduce the high temperature ductility of the steel sheet by forming nitrides (e.g., BN).
Other elements
Phosphorus (P), sulfur (S), and nitrogen (N) may be inevitably added to steel during steel making. I.e. ideally these elements are preferably not included, but may also be included in a certain amount, since it is difficult in terms of process technology to remove these elements completely.
The role of phosphorus (P) in steel may be similar to that of silicon. However, if phosphorus is added in an amount greater than 0.02% by total weight of the steel sheet, weldability of the steel sheet may be reduced and brittleness thereof may be increased, resulting in deterioration of material properties. Therefore, the amount of phosphorus added may be controlled to 0.02% or less of the total weight of the steel sheet.
Sulfur (S) may inhibit toughness and weldability of steel, and thus its content may be controlled to 0.005% or less of the total weight of the steel sheet.
If nitrogen (N) is present in excess in the steel, a large amount of nitrides may be precipitated, resulting in deterioration of ductility of the steel sheet. Therefore, the content of nitrogen (N) can be controlled to 0.006% or less of the total weight of the steel sheet.
The high-strength steel sheet of the present invention having the above alloy composition has a microstructure composed of ferrite and retained austenite. In this case, the volume fraction of the retained austenite in the microstructure may be 10 to 30 volume%. The grains of the high strength steel sheet may be fine grains having a size of 3 μm or less. The fraction of high angle grain boundaries in the grains may be 70% or more.
The material properties of the high strength steel sheet may include a Yield Strength (YS) of 800MPa or more, a Tensile Strength (TS) of 980MPa or more, an Elongation (EL) of 25% or more, and a Hole Expansion Ratio (HER) of 20% or more.
Therefore, the high strength steel sheet according to the embodiment of the present invention may be applicable to fields requiring high strength and high formability.
The above-described high strength steel sheet according to an embodiment of the present invention may be manufactured by the method of the following exemplary embodiment. The present invention is directed to provide a steel sheet having excellent elongation, hole expansion ratio and strength by using alloy components having suitably controlled composition ratios and performing a two-step annealing heat treatment after performing a hot rolling process and a cold rolling process, and a method of manufacturing the same.
Method for manufacturing steel sheet having high strength and high formability
Fig. 1 is a process flow diagram schematically showing a method of manufacturing a steel sheet having high strength and high formability according to an exemplary embodiment of the present invention.
Referring to fig. 1, the method of manufacturing a steel sheet includes the steps of: (S110) reheating the billet; (S120) manufacturing a hot-rolled steel sheet by hot-rolling the slab; (S130) cold rolling the hot rolled steel sheet; and (S140) annealing the cold-rolled steel sheet.
First, the step of reheating the slab (S110) is a step of preparing a slab comprising, in wt%, carbon (C) in an amount of 0.05% to 0.15%, silicon (Si) in an amount of more than 0% and less than or equal to 0.4%, manganese (Mn) in an amount of 4.0% to 9.0%, aluminum (Al) in an amount of more than 0% and less than or equal to 0.3%, phosphorus (P) in an amount of 0.02% or less, sulfur (S) in an amount of 0.005% or less, nitrogen (N) in an amount of 0.006% or less, and the balance iron (Fe) and other unavoidable impurities, and reheating the slab to re-dissolve components segregated during casting and homogenize as-cast components. Meanwhile, the steel slab may further include one or more elements of niobium (Nb), titanium (Ti), vanadium (V), and molybdenum (Mo), and a content of each of the elements may be greater than 0% and less than or equal to 0.02 wt%. The steel slab may further contain more than 0% and 0.001 wt% or less of boron (B).
The slab reheating temperature is preferably about 1150 to 1250 c so that a normal hot exit temperature can be ensured. If the reheating temperature is less than 1150 deg.c, there may be a problem in that a hot rolling load is rapidly increased, and if the reheating temperature is more than 1250 deg.c, it may be difficult to secure strength of a finally manufactured steel sheet due to coarsening of initial austenite grains.
Then, after reheating the slab, a hot rolling step (S120) is performed, which is a step of forming a hot rolled steel sheet by performing hot rolling by a conventional method and finish rolling at a temperature of 925 ℃ to 975 ℃. Considering that the billet of the present invention has a high content of alloying elements (e.g., manganese), it can be finish rolled at a high temperature of 925 to 975 ℃. After the finish rolling, the hot rolled steel sheet is cooled to a temperature of 700 ℃ to 800 ℃ at a cooling rate of 10 ℃/s to 30 ℃/s, and then wound. The cooling method may be performed using a water-free cooling method. After cooling, the hot rolled steel sheet may have a full martensite structure.
According to some exemplary embodiments, before cold rolling a hot rolled steel sheet having a full martensite structure, a softening heat treatment may be performed to reduce a rolling load during the cold rolling. The softening heat treatment may be performed at a temperature of 550 ℃ to 650 ℃. If the temperature of the softening heat treatment is less than 550 c, martensite produced after hot rolling may not be recrystallized and may be tempered only, so that supersaturated carbon in the form of cementite may be formed in the structure and spheroidization may occur. In this case, since martensite may exhibit brittleness, breakage of the steel sheet may occur during the cold rolling. On the other hand, if the temperature of the softening heat treatment is higher than 650 ℃, austenite may be excessively formed, and the austenite may form martensite during cooling, so that the effect of the softening heat treatment may not work. By the softening heat treatment performed in the above temperature range, the martensite structure after hot rolling can be converted into a composite structure of ferrite and retained austenite.
Then, the cold rolling step (S130) is a step of cold rolling the hot-rolled steel sheet after pickling. The cold rolling may be performed under a condition that the hot rolled steel sheet is cold rolled at a reduction ratio of 40 to 60%. By cold rolling, the composite structure of ferrite and retained austenite after softening heat treatment can be transformed into a composite structure of ferrite and martensite.
Then, the annealing heat treatment step (S140) may include a step of performing a first heat treatment on the cold-rolled steel sheet at a temperature of AC3 to (AC 3+ 15) ° c, and a step of performing a second heat treatment on the cold-rolled steel sheet subjected to the first heat treatment at a critical zone temperature. The temperature of AC3 to (AC 3+ 15) ° c in the first heat treatment step may be, for example, a temperature of 735 ℃ to 750 ℃. The critical zone temperature in the second heat treatment step may be, for example, a temperature of 640 ℃ to 660 ℃.
In exemplary embodiments, the first heat treatment may transform a composite structure of ferrite and martensite in the steel sheet after cold rolling into a martensite structure. In the first heat treatment, the heat treatment is performed by: heating the cold-rolled steel sheet to a target temperature of 735 ℃ to 750 ℃ at a heating rate of 1 ℃/s to 3 ℃/s, and maintaining the cold-rolled steel sheet at the target temperature for 40 seconds to 120 seconds.
If the heat treatment temperature is less than 735 ℃, austenite grains having a sufficient size may not be ensured at a target temperature, and a composite structure of martensite and ferrite may be formed after the heat treatment, and thus the strength and ductility of the final structure after the annealing heat treatment may be reduced. On the other hand, if the heat treatment temperature is higher than 750 ℃, the size of austenite grains at the target temperature may be excessively increased, which is disadvantageous to ensure the stabilization of austenite in the final structure after the annealing heat treatment, and thus the strength of the steel sheet may be poor.
Further, if the heating rate is less than 1 ℃/s, the retention time at the target temperature of 735 ℃ to 750 ℃ may be longer than the upper limit of the range of 40 seconds to 120 seconds, and thus the austenite grain size at the target temperature may be excessively increased. On the other hand, if the heating rate is greater than 3 ℃/s, the retention time at the target temperature of 735 ℃ to 750 ℃ may be shorter than the lower limit of the range of 40 seconds to 120 seconds, and therefore austenite grains having a sufficient size cannot be ensured at the target temperature.
Then, the heat-treated cold-rolled steel sheet is cooled to a temperature of 350 ℃ to 450 ℃ at a cooling rate of 4 ℃/s to 10 ℃/s. In an exemplary embodiment, the cold rolled steel sheet cooled to the above temperature may be aged for 120 seconds to 330 seconds.
The cold-rolled steel sheet having undergone the first heat treatment may be continuously subjected to the second heat treatment. In an exemplary embodiment, in the second heat treatment, the heat treatment is performed by: heating the cold-rolled steel sheet to a target temperature of 640 ℃ to 660 ℃ at a heating rate of 1 ℃/s to 3 ℃/s, and maintaining the cold-rolled steel sheet at the target temperature for 40 seconds to 120 seconds. Since the second heat treatment is performed at the intercritical temperature corresponding to the target temperature range, the martensite structure after the first heat treatment can be transformed into a structure consisting of ferrite and retained austenite. In this case, the volume fraction of the retained austenite may be 10 to 30 volume%.
If the second heat treatment temperature is less than 640 deg.c, the austenite structure formed at the target temperature may be too small and the austenite stability may increase, and thus, the austenite in the microstructure after cooling may not undergo phase transformation during plastic deformation, and thus the strength and ductility of the steel sheet may decrease. On the other hand, if the second heat treatment temperature is higher than 660 ℃, the austenite structure formed at the target temperature may be excessive and the austenite stability may be reduced, and thus, martensite may be formed in the microstructure after cooling, resulting in a reduction in ductility and hole expansion ratio of the steel sheet.
If the heating rate is less than 1 deg.C/s, unwanted cementite may be formed or spheroidization may occur before the cold-rolled sheet reaches the above critical zone temperature range, resulting in deterioration of material properties of the steel sheet. If the heating rate is greater than 3 deg.C/s, the steel sheet may not be maintained within the target temperature range for 40 seconds to 120 seconds, thereby failing to secure a sufficient fraction of retained austenite in the final structure.
Then, the heat-treated cold-rolled steel sheet is cooled to a temperature of 350 ℃ to 450 ℃ at a cooling rate of 4 ℃/s to 10 ℃/s. In an exemplary embodiment, the cold rolled steel sheet cooled to the above temperature may be aged for 120 seconds to 330 seconds.
By the above method, a steel sheet having high strength and high formability according to an exemplary embodiment of the present invention may be manufactured.
The steel sheet of the present invention manufactured by the above method may have a Yield Strength (YS) of 800MPa or more, a Tensile Strength (TS) of 980MPa or more, an Elongation (EL) of 25% or more, and a Hole Expansion Ratio (HER) of 20% or more.
As described above, in the manufacturing method according to the exemplary embodiment of the present invention, the predetermined amount of the austenite stabilizing element as described above may be added to the steel billet. Further, since the first annealing heat treatment and the second annealing heat treatment are continuously performed, the final microstructure of the steel sheet may be a composite structure consisting of fine-grained ferrite and 10 to 30 vol% of retained austenite. Since the steel sheet has a sufficient fraction of retained austenite, the steel sheet may have a high elongation of 25% or more due to the transformation induced plasticity property of the retained austenite.
Further, the phase boundary between the hard phase and the soft phase can be reduced by controlling so that the hard phase (e.g., martensite and bainite) is not included in the final microstructure as described above. In addition, the interface between the precipitates and the crystal grains can be reduced by controlling the contents of the precipitate forming elements (e.g., titanium, niobium, and vanadium) and the precipitate growth inhibiting elements (e.g., molybdenum) in the steel slab component system. Further, by performing the annealing heat treatment in two steps (including the first heat treatment step and the second heat treatment step) within a predetermined temperature range, the fraction of High Angle Grain Boundaries (HAGB) in the final structure can be increased. In the first heat treatment, since the high dislocation density exists in the martensite formed through the cold rolling process, recrystallization may actively occur before the martensite is reversely transformed into austenite. In the second heat treatment, the martensite formed by the first heat treatment is heat-treated, and thus recrystallization is relatively suppressed before the martensite is reversely transformed into austenite, whereby the fraction of high angle grain boundaries in the final microstructure can be increased to 70% or more of the grains. Therefore, the steel sheet may have a high hole expansion ratio of 20% or more.
Then, the grains of the final microstructure may be refined to give high strength to the steel sheet. In particular, the grain size of the initial austenite can be optimized by performing the first heat treatment at a temperature of AC3 to (AC 3+ 15) ° c. Further, by performing the second heat treatment in the intercritical temperature range, the grain sizes of ferrite and residual austenite in the final microstructure can be controlled to 3 μm or less.
Modes for the invention
The constitution and effect of the present invention will be described in more detail with reference to preferred embodiments of the present invention. However, the following examples are provided to aid understanding of the present invention, and the scope of the present invention is not limited to the following examples.
Example 1
Steel billets having the comparative component system and the embodied component system shown in table 1 below were produced by a continuous casting process. Test specimens were prepared from individual billets and subjected to a high temperature tensile test. In the case of the comparative component system, the contents of silicon and aluminum are greater than the upper limit of the range of contents of silicon and aluminum according to the exemplary embodiment of the present invention.
[ Table 1]
Figure BDA0003351403000000121
Figure 2 shows the results of the high temperature tensile test of the comparative component system samples of the present invention and figure 3 shows the results of the high temperature tensile test of the embodied component system samples of the present invention. Specifically, the results of the high-temperature tensile test were obtained by heating each of the comparative component system specimen and the conducted component system specimen to a temperature of 700 ℃, 750 ℃,800 ℃, 850 ℃, 900 ℃, 950 ℃, 1000 ℃ and 1100 ℃, and then subjecting each specimen to the tensile test at the above-described temperature. With respect to the high temperature tensile test, FIG. 3 shows a curve 201 obtained by heating the test specimen to a temperature higher than 1100 deg.C, and then cooling the test specimens to the tensile test temperature at a cooling rate of-1 deg.C/s, respectively, and a curve 202 obtained by cooling the test specimen to the tensile test temperature at a cooling rate of-20 deg.C/s, respectively. In general, when the area reduction rate at a predetermined temperature is 50% or more, it is determined that ductility at the predetermined temperature is ensured.
Referring to fig. 2, in the case of the comparative component system sample, the area reduction rate at 1100 ℃ was 55%, the area reduction rate in the temperature range of 700 ℃ to 800 ℃ was 50%, and the area reduction rate in the temperature range of 800 ℃ to 1050 ℃ was less than 50% (50% being the target value). On the other hand, referring to fig. 3, the area reduction rate in the temperature range of 800 ℃ to 1100 ℃ exceeds 50% (50% is the target value).
Referring to fig. 2 and 3, in the case of the comparative component system sample, unlike the embodied component system sample of the embodiment of the present invention, high temperature ductility cannot be secured at a high temperature range equal to or higher than 800 ℃ in which continuous casting is performed according to the embodiment of the present invention, and thus cracks may be generated during continuous casting, thereby failing to secure a good blank.
Table 2 below shows the rolling force of each pass calculated by simulating hot rolling according to an exemplary embodiment of the present invention for each of the comparative component system sample and the practical component system sample.
[ Table 2]
Figure BDA0003351403000000131
Referring to table 2 above, it can be seen that in order to ensure that each rolling pass produces the same reduction, a greater rolling force must be applied to the comparative compositional system sample than to the rolling force applied to the embodied compositional system sample. It was determined that a relatively high load was applied to the rolling mill during hot rolling of the comparative component system samples.
Example 2
The samples prepared from the embodied composition systems shown in table 1 above were subjected to the first annealing heat treatment process and the second annealing heat treatment process, respectively, according to table 3 below. In the case of comparative examples 1 and 3, the second annealing temperature was less than 640 ℃ (640 ℃ is the lower limit of the second annealing temperature according to the embodiment of the present invention). In the case of comparative examples 2 and 4, the second annealing temperature was higher than 660 ℃ (660 ℃ is the upper limit of the second annealing temperature according to an embodiment of the present invention). In the case of comparative examples 5 to 7, the first annealing temperature was higher than 750 ℃ (750 ℃ is the upper limit of the first annealing temperature according to an embodiment of the present invention). Further, in the case of comparative example 7, the second annealing temperature was higher than 660 ℃ (660 ℃ is an upper limit of the second annealing temperature according to an embodiment of the present invention). In the case of comparative examples 8 to 11, the first annealing heat treatment was not performed, and only the second annealing heat treatment was performed. Further, in the case of comparative example 11, the second annealing temperature was higher than 660 ℃ (660 ℃ is an upper limit of the second annealing temperature according to an embodiment of the present invention).
[ Table 3]
Figure BDA0003351403000000132
Figure BDA0003351403000000141
Table 4 shows the evaluation results of the material properties of the samples of comparative examples 1 to 11 and examples 1 to 6 subjected to the annealing heat treatment according to table 3.
[ Table 4]
Figure BDA0003351403000000142
Figure BDA0003351403000000151
The target values of the material properties of the high strength steel sheet according to the exemplary embodiment of the present invention are yield strength of 800MPa or more, tensile strength of 980MPa or more, elongation of 25% or more, retained austenite volume fraction of 10% to 30%, large angle grain boundary (HAGB) fraction of 70% or more, and hole expansion ratio of 20% or more. The samples of examples 1 to 6 satisfied all of the above-mentioned target values. In the case of comparative example 1, the fraction of elongation and High Angle Grain Boundaries (HAGB) was lower than the target value. In the case of comparative example 2, the elongation was lower than the target value. In the case of comparative example 3, the tensile strength, elongation, and fraction of High Angle Grain Boundaries (HAGB) were lower than the target values. In the case of comparative example 4, the elongation, tensile strength × elongation, average grain size, and fraction of High Angle Grain Boundaries (HAGB) were lower than the target values. In the case of comparative example 5, the tensile strength, elongation, average grain size, and fraction of High Angle Grain Boundaries (HAGB) were lower than the target values. In the case of comparative examples 6 and 7, the yield strength, tensile strength, elongation, average grain size, and fraction of High Angle Grain Boundaries (HAGB) were lower than the target values. In the case of comparative example 8, the tensile strength, elongation, and fraction of High Angle Grain Boundaries (HAGB) were lower than the target values. In the case of comparative examples 9 to 11, the elongation, the fraction of High Angle Grain Boundaries (HAGB), and the pore expansion ratio were lower than the target values.
Fig. 4 is a photograph showing the microstructure of a high-strength steel sheet according to an exemplary embodiment of the present invention. Specifically, fig. 4 is a microstructure photograph of the sample of example 1. Referring to table 4 and fig. 4, 17% of retained austenite and remaining ferrite were observed in the sample of example 1 by volume fraction.
Example 3
The samples prepared from the embodied composition systems shown in table 1 above were subjected to a first annealing heat treatment process and a second annealing heat treatment process according to table 5 below.
[ Table 5]
Figure BDA0003351403000000161
Figure BDA0003351403000000171
Referring to table 5 above, in the case of comparative example 12, the heating rate during the first annealing heat treatment was greater than 3 ℃/s (3 ℃/s is an upper limit of the heating rate during the first annealing heat treatment according to the exemplary embodiment of the present invention), and the first annealing holding time did not satisfy 40 seconds or more. In the case of comparative example 13, the heating rate during the first annealing heat treatment was less than 1 ℃/s (1 ℃/s is a lower limit of the heating rate during the first annealing heat treatment according to the exemplary embodiment of the present invention), and the first annealing hold time exceeded 120 seconds (120 seconds is an upper limit). In the case of comparative example 14, the heating rate during the first annealing heat treatment was less than 1 ℃/s (1 ℃/s is a lower limit of the heating rate during the first annealing heat treatment according to the exemplary embodiment of the present invention), and the first annealing hold time exceeded 120 seconds (120 seconds is an upper limit). Further, the cooling rate is less than 4 ℃/s (4 ℃/s is a lower limit). Examples 7 to 10 satisfy the first and second annealing heat treatment conditions according to the exemplary embodiment of the present invention.
Table 6 below shows the evaluation results of the material properties of the samples of comparative examples 12 to 14 and examples 7 to 10, which were subjected to the annealing heat treatment according to table 5 above.
[ Table 6]
Figure BDA0003351403000000172
Figure BDA0003351403000000181
Referring to table 6 above, in the case of comparative example 12, the target values of tensile strength and elongation were not reached. In the case of comparative example 13, the target values of tensile strength, elongation and average grain size were not reached. In the case of comparative example 14, the target values of yield strength, tensile strength, elongation and average grain size were not reached. Examples 7 to 10 met all target values for material properties according to embodiments of the present invention.
Example 4
The samples prepared from the embodied composition systems shown in table 1 above were subjected to the first annealing heat treatment process and the second annealing heat treatment process, respectively.
[ Table 7]
Figure BDA0003351403000000182
Referring to table 7 above, in the case of comparative example 15, the heating rate during the second annealing heat treatment was greater than 3 ℃/s (3 ℃/s is an upper limit of the heating rate during the second annealing according to the exemplary embodiment of the present invention), and the second annealing holding time did not satisfy 40 seconds or more. In the case of comparative example 16, the heating rate during the second annealing was less than 1 ℃/s (1 ℃/s is the lower limit of the heating rate during the second annealing according to the exemplary embodiment of the present invention), and the second annealing holding time exceeded 120 seconds (120 seconds is the upper limit). Examples 11 to 14 satisfied the first annealing heat treatment condition and the second annealing heat treatment condition according to the exemplary embodiment of the present invention.
Table 8 below shows the evaluation results of the material properties of the samples of comparative examples 15 to 16 and examples 11 to 14, which were subjected to the annealing heat treatment according to table 7.
[ Table 8]
Figure BDA0003351403000000191
Referring to table 8, in the case of comparative example 15, the target values of tensile strength and elongation were not reached. In the case of comparative example 16, the target value of elongation was not reached. Examples 11 to 14 satisfy all target values of material properties according to exemplary embodiments of the present invention.
Although the foregoing description has been made with reference to exemplary embodiments of the present invention, various changes or modifications may be made by those skilled in the art. Such changes and modifications are to be considered within the scope of the present invention unless they depart therefrom. Accordingly, the scope of the invention should be determined from the following claims.
Simple modifications or changes of the present invention can be easily implemented by those skilled in the art and are considered to be included in the scope of the present invention.

Claims (10)

1. A steel sheet having high strength and high formability, which comprises, in wt.%, carbon (C) in an amount of 0.05 to 0.15%, silicon (Si) in an amount of greater than 0% and less than or equal to 0.4%, manganese (Mn) in an amount of 4.0 to 9.0%, aluminum (Al) in an amount of greater than 0% and less than or equal to 0.3%, phosphorus (P) in an amount of 0.02% or less, sulfur (S) in an amount of 0.005% or less, nitrogen (N) in an amount of 0.006% or less, and the balance iron (Fe) and other unavoidable impurities, wherein the steel sheet comprises a microstructure consisting of ferrite and residual austenite,
wherein the volume fraction of retained austenite in the microstructure is 10 to 30 vol%,
wherein the microstructure has a grain size of 3 μm or less, a fraction of high angle grain boundaries in the grains is 70% or more, and the steel sheet has a Yield Strength (YS) of 800MPa or more, a Tensile Strength (TS) of 980MPa or more, an Elongation (EL) of 25% or more, and a Hole Expansion Ratio (HER) of 20% or more.
2. The steel sheet according to claim 1, comprising one or more elements of niobium (Nb), titanium (Ti), vanadium (V) and molybdenum (Mo), each of which is present in an amount greater than 0% and less than or equal to 0.02% by weight.
3. The steel sheet according to claim 1, further comprising boron (B) in an amount of greater than 0% and less than or equal to 0.001 wt%.
4. A method of manufacturing a steel sheet having high strength and high formability, the method comprising the steps of:
(a) Manufacturing a hot-rolled steel sheet from a steel slab comprising, in weight%: carbon (C) in an amount of 0.05% to 0.15%, silicon (Si) in an amount of greater than 0% and less than or equal to 0.4%, manganese (Mn) in an amount of 4.0% to 9.0%, aluminum (Al) in an amount of greater than 0% and less than or equal to 0.3%, phosphorus (P) in an amount of 0.02% or less, sulfur (S) in an amount of 0.005% or less, nitrogen (N) in an amount of 0.006% or less, and iron (Fe) and other unavoidable impurities as the balance;
(b) Manufacturing a cold-rolled steel sheet by cold-rolling the hot-rolled steel sheet;
(c) Performing a first heat treatment on a cold-rolled steel sheet by: heating the cold-rolled steel sheet to a temperature of 735 ℃ to 750 ℃ at a heating rate of 1 ℃/s to 3 ℃/s, and holding the cold-rolled steel sheet at the temperature for 40 seconds to 120 seconds; and
(d) Performing a second heat treatment on the cold-rolled steel sheet subjected to the first heat treatment by: heating the cold-rolled steel sheet to a temperature of 640 ℃ to 660 ℃ at a heating rate of 1 ℃/s to 3 ℃/s, and maintaining the cold-rolled steel sheet at the temperature for 40 seconds to 120 seconds,
wherein the cold rolled steel sheet after the step (d) has a microstructure consisting of ferrite and austenite,
wherein the volume fraction of retained austenite in the microstructure is 10 to 30 vol%,
wherein the cold rolled steel sheet after the step (d) has a grain size of 3 μm or less, and a fraction of high angle grain boundaries in the grains is 70% or more,
wherein the cold rolled steel sheet after step (d) has a Yield Strength (YS) of 800MPa or more, a Tensile Strength (TS) of 980MPa or more, an Elongation (EL) of 25% or more, and a Hole Expansion Ratio (HER) of 20% or more.
5. The method of claim 4 wherein said steel slab comprises one or more elements of niobium (Nb), titanium (Ti), vanadium (V) and molybdenum (Mo), each said element being present in an amount greater than 0% and less than or equal to 0.02% by weight.
6. The method of claim 4 wherein the steel slab further comprises boron (B) in an amount greater than 0% and less than or equal to 0.001% by weight.
7. The method as set forth in claim 4, wherein the step (c) comprises the step of cooling the cold-rolled steel sheet subjected to the heat treatment to a temperature of 350 ℃ to 450 ℃ at a cooling rate of 4 ℃/s to 10 ℃/s.
8. The method as set forth in claim 7, wherein the step (d) comprises the step of cooling the cold rolled steel sheet subjected to the heat treatment to a temperature of 350 ℃ to 450 ℃ at a cooling rate of 4 ℃/s to 10 ℃/s.
9. The method of claim 4, wherein step (a) comprises the steps of:
(a1) Reheating the billet to a temperature of 1150 ℃ to 1250 ℃;
(a2) Hot rolling the reheated slab to a finish rolling exit temperature of 925 ℃ to 975 ℃; and
(a3) The hot rolled steel sheet is cooled to a temperature of 700 ℃ to 800 ℃ at a cooling rate of 10 ℃/s to 30 ℃/s, followed by winding.
10. The method according to claim 4, further comprising a step of subjecting the hot rolled steel sheet to a softening heat treatment at a temperature of 550 ℃ to 650 ℃ between steps (a) and (b).
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