Patent Grant
12252759
The present invention relates to a steel sheet that may be used for automobile parts and the like, and to a steel sheet having high strength characteristics and excellent workability and a method for manufacturing the same.
In recent years, the automobile industry has been paying attention to ways to reduce material weight and secure occupant safety in order to protect the global environment. In order to meet these requirements for safety and weight reduction, the application of a high strength steel sheet is rapidly increasing. In general, it has been known that as the strength of the steel sheet increases, the workability of the steel sheet is lowered. Therefore, in the steel sheet for automobile parts, a steel sheet having excellent workability represented by ductility, bending formability, and hole expansion ratio while having high strength characteristics is required.
As a technique for improving workability of a steel sheet, a method of utilizing tempered martensite is disclosed in Patent Documents 1 and 2. Since the tempered martensite made by tempering hard martensite is softened martensite, there is a difference in strength between tempered martensite and existing untempered martensite (fresh martensite). Therefore, when fresh martensite is suppressed and tempered martensite is formed, the workability may increase.
However, by the techniques disclosed in Patent Documents 1 and 2, a balance (TSXE1) of tensile strength and elongation does not satisfy 22,000 MPa % or more, meaning that it is difficult to secure a steel sheet having superb strength and ductility.
Meanwhile, transformation induced plasticity (TRIP) steel using transformation-induced plasticity of retained austenite has been developed in order to obtain both high strength and excellent workability for automobile member steel sheets. Patent Document 3 discloses TRIP steel having excellent strength and workability.
Patent Document 3 discloses improving high ductility and workability by including polygonal ferrite, retained austenite, and martensite. However, it can be seen that Patent Document 3 uses bainite as a main phase, and thus, the high strength is not secured and a balance (TSXE1) of tensile strength and elongation also does not satisfy 22,000 MPa % or more.
That is, demand for a steel sheet having excellent workability, such as ductility, bending formability, and hole expansion ratio while having high strength, is not satisfied.
(Patent Document 1) Korean Patent Laid-Open Publication No. 10-2006-0118602
(Patent Document 2) Japanese Patent Laid-Open Publication No. 2009-019258
(Patent Document 3) Korean Patent Laid-Open Publication No. 10-2014-0012167
The present invention provides a high strength steel sheet having superb ductility, bending formability, and hole expansion ratio by optimizing a composition and microstructure of the steel sheet and a method for manufacturing the same.
An object of the present invention is not limited to the abovementioned contents. Additional problems of the present invention are described in the overall content of the specification, and those of ordinary skill in the art to which the present invention pertains will have no difficulty in understanding the additional problems of the present invention from the contents described in the specification of the present invention.
In an aspect of the present invention, a high strength steel sheet having excellent workability may include: by wt %, C: 0.25 to 0.75%, Si: 4.0% or less, Mn: 0.9 to 5.0%, Al: 5.0% or less, P: 0.15% or less, S: 0.03% or less, N: 0.03% or less, a balance of Fe, and unavoidable impurities; and, as microstructures, ferrite which is a soft structure, and tempered martensite, bainite, and retained austenite which are hard structures, in which the high strength steel sheet may satisfy the following [Relational Expression 1], [Relational Expression 2], and [Relational Expression 3].
0.4≤[H]F/[H]TM+B+γ≤0.9 [Relational Expression 1]
In the above Relational Expression 1, [H]F and [H]TM+B+γ may be nanohardness values measured using a nanoindenter, [H]F may be an average nanohardness value Hv of the ferrite which is the soft structure, and [H]TM+B+γ may be the average nanohardness value Hv of the tempered martensite, the bainite, and the retained austenite which are the hard structures.
V(1.2 μm,γ)/V(γ)≥0.1 [Relational Expression 2]
in Relational Expression 2, V(1.2 μm, γ) may be a fraction (vol %) of the retained austenite having an average grain size of 1.2 μm or more, and V(γ) may be the fraction (vol %) of the retained austenite of the steel sheet.
V(lath,γ)/V(γ)≥0.5 [Relational Expression 3]
In the above Relational Expression 3, V(lath, γ) may be the fraction (vol %) of the retained austenite in a lath form, and V(γ) may be the fraction (vol %) of the retained austenite of the steel sheet.
The high strength steel sheet may further include: any one or more of the following (1) to (9).
(1) one or more of Ti: 0 to 0.5%, Nb: 0 to 0.5%, and V: 0 to 0.5%
(2) one or more of Cr: 0 to 3.0% and Mo: 0 to 3.0%
(3) one or more of Cu: 0 to 4.5% and Ni: 0 to 4.5%
(4) B: 0 to 0.005%
(5) one or more of Ca: 0 to 0.05%, REM: 0 to 0.05% excluding Y, and Mg: 0 to 0.05%
(6) one or more of W: 0 to 0.5% and Zr: 0 to 0.5%
(7) one or more of Sb: 0 to 0.5% and Sn: 0 to 0.5%
(8) one or more of Y: 0 to 0.2% and Hf: 0 to 0.2%
(9) Co: 0 to 1.5%
A total content (Si+Al) of Si and Al may be 1.0 to 6.0 wt %.
The steel sheet may include, by volume fraction, 30 to 70% of tempered martensite, 10 to 45% of bainite, 10 to 40% of retained austenite, 3 to 20% of ferrite, and an unavoidable structure.
A balance BT·E of tensile strength and elongation expressed by the following [Relational Expression 4] may be 22,000 (MPa %) or more, a balance BT·H of tensile strength and hole expansion ratio expressed by the following [Relational Expression 5] may be 7*106 (MPa2%1/2) or more, and bendability BR expressed by the following [Relational Expression 6] may be 0.5 to 3.0.
BT·E=[Tensile Strength (TS, MPa)]*[Elongation (El, %)] [Relational Expression 4]
BT·H=[Tensile Strength (TS, MPa)]2*[Hole Extension Ratio (HER, %)]1/2 [Relational Expression 5]
BR=R/t [Relational Expression 6]
In the above Relational Expression 6, R may mean a minimum bending radius (mm) at which cracks do not occur after a 90° bending test, and t may mean a thickness (mm) of the steel sheet.
In another aspect of the present invention, a method for manufacturing a high strength steel sheet having excellent workability may include: providing a cold-rolled steel sheet including, by wt %, C: 0.25 to 0.75%, Si: 4.0% or less, Mn: 0.9 to 5.0%, Al: 5.0% or less, P: 0.15% or less, S: 0.03% or less, N: 0.03% or less, a balance of Fe, and unavoidable impurities; heating (primarily heating) the cold-rolled steel sheet to a temperature within a range of Ac1 or higher and less than Ac3, and maintaining (primarily maintaining) the primarily heated steel sheet for 50 seconds or more; cooling (primary cooling) the primarily heated steel sheet to a temperature within a range (primarily cooling stop temperature) of 600 to 850° C. at an average cooling rate of 1° C./s or more; cooling (secondarily cooling) the primarily cooled steel sheet to a temperature within a range of 300 to 500° C. at an average cooling rate of 2° C./s or more, and maintaining (secondarily maintaining) the secondarily cooled steel sheet in the temperature within a range for 5 seconds or more; cooling (tertiarily cooling) the secondarily cooled steel sheet to a temperature within a range (secondary cooling stop temperature) of 100 to 300° C. at an average cooling rate of 2° C./s or more; heating (secondarily heating) the tertiarily cooled steel sheet to a temperature within a range of 350 to 550° C., and maintaining (tertiarily maintaining) the secondarily heated steel sheet in the temperature within a range for 10 seconds or more; heating (quaternarily cooling) the secondarily heated steel sheet to a temperature within a range of 250 to 450° C., and maintaining (quaternarily maintaining) the quaternarily cooled steel sheet in the temperature within a range for 10 seconds or more; and cooling (fifth cooling) quaternarily cooled steel sheet to room temperature.
The steel slab may further include any one or more of the following (1) to (9).
(1) one or more of Ti: 0 to 0.5%, Nb: 0 to 0.5%, and V: 0 to 0.5%
(2) one or more of Cr: 0 to 3.0% and Mo: 0 to 3.0%
(3) one or more of Cu: 0 to 4.5% and Ni: 0 to 4.5%
(4) B: 0 to 0.005%
(5) one or more of Ca: 0 to 0.05%, REM: 0 to 0.05% excluding Y, and Mg: 0 to 0.05%
(6) one or more of W: 0 to 0.5% and Zr: 0 to 0.5%
(7) one or more of Sb: 0 to 0.5% and Sn: 0 to 0.5%
(8) one or more of Y: 0 to 0.2% and Hf: 0 to 0.2%
(9) Co: 0 to 1.5%.
A total content (Si+Al) of Si and Al included in the steel slab may be 1.0 to 6.0 wt %.
The providing of the cold-rolled steel sheet may include: heating a steel slab to 1000 to 1350° C.; performing finishing hot rolling in a temperature within a range of 800 to 1000° C.; coiling the hot-rolled steel sheet at a temperature within a range of 300 to 600° C.; performing hot-rolled annealing heat treatment on the coiled steel sheet in a temperature within a range of 650 to 850° C. for 600 to 1700 seconds; and cold rolling the hot-rolled annealing heat-treated steel sheet at a reduction ratio of 30 to 90%.
According to an aspect of the present invention, it is possible to provide a steel sheet particularly suitable for automobile parts because the steel sheet has superb strength as well as excellent workability such as ductility, bending formability, and hole expansion ratio.
The present invention relates to a high strength steel sheet having excellent workability and a method for manufacturing the same, and exemplary embodiments in the present invention will hereinafter be described. Exemplary embodiments in the present invention may be modified into several forms, and it is not to be interpreted that the scope of the present invention is limited to exemplary embodiments described below. The present exemplary embodiments are provided in order to further describe the present invention in detail to those skilled in the art to which the present invention pertains.
The inventors of the present invention recognized that, in a transformation induced plasticity (TRIP) steel including bainite, tempered martensite, retained austenite, and ferrite, when controlling a ratio of specific components included in the retained austenite and the ferrite to a certain range while promoting stabilization of the retained austenite, it is possible to simultaneously secure workability and strength of a steel sheet by reducing an interphase hardness difference between the retained austenite and the ferrite. Based on this, the present inventors have reached the present invention by devising a method capable of improving ductility and workability of the high strength steel sheet.
Hereinafter, a high strength steel sheet having excellent workability according to an aspect of the present invention will be described in more detail.
The high strength steel sheet having excellent workability according to an aspect of the present invention includes, by wt %, C: 0.25 to 0.75%, Si: 4.0% or less, Mn: 0.9 to 5.0%, Al: 5.0% or less, P: 0.15% or less, S: 0.03% or less, N: 0.03% or less, a balance of Fe, and unavoidable impurities, and includes, as microstructures, ferrite which is a soft structure, and tempered martensite, bainite, and retained austenite which are hard structures, and may satisfy the following [Relational Expression 1], [Relational Expression 2], and [Relational Expression 3].
0.4≤[H]F/[H]TM+B+γ≤0.9 [Relational Expression 1]
In the above Relational Expression 1, [H]F and [H]TM+B+γ are nanohardness values measured using a nanoindenter, [H]F is an average nanohardness value Hv of the ferrite which is the soft structure, and [H]TM+B+γ is an average nanohardness value Hv of tempered martensite, bainite, and retained austenite which are the hard structures.
V(1.2 μm, γ)/V(γ)≤0.1 [Relational Expression 2]
In the above Relational Expression 2, V(1.2 μm, γ) is a fraction (vol %) of the retained austenite having an average grain size of 1.2 μm or more, and V(γ) is the fraction (vol %) of the retained austenite of the steel sheet.
V(lath, γ)/V(γ)≥0.5 [Relational Expression 3]
In the above Relational Expression 3, V(lath, γ) is the fraction (vol %) of the retained austenite in a lath form, and V(γ) is the fraction (vol %) of the retained austenite of the steel sheet.
Hereinafter, compositions of steel according to the present invention will be described in more detail. Hereinafter, unless otherwise indicated, % indicating a content of each element is based on weight.
The high strength steel sheet having excellent workability according to an aspect of the present invention includes: by wt %, C: 0.25 to 0.75%, Si: 4.0% or less, Mn: 0.9 to 5.0%, Al: 5.0% or less, P: 0.15% or less, S: 0.03% or less, N: 0.03% or less, a balance of Fe, and unavoidable impurities. In addition, the high strength steel sheet may further include one or more of Ti: 0.5% or less (including 0%), Nb: 0.5% or less (including 0%), V: 0.5% or less (including 0%), Cr: 3.0% or less (including 0%), Mo: 3.0% or less (including 0%), Cu: 4.5% or less (including 0%), Ni: 4.5% or less (including 0%), B: 0.005% or less (including 0%), Ca: 0.05% or less (including 0%), REM: 0.05% or less (including 0%) excluding Y, Mg: 0.05% or less (including 0%), W: 0.5% or less (including 0%), Zr: 0.5% or less (including 0%), Sb: 0.5% or less (including 0%), Sn: 0.5% or less (including 0%), Y: 0.2% or less (including 0%), Hf: 0.2% or less (including 0%), Co: 1.5% or less (including 0%). In addition, a total content (Si+Al) of Si and Al may be 1.0 to 6.0%.
Carbon (C): 0.25 to 0.75%
Carbon (C) is an unavoidable element for securing strength of a steel sheet, and is also an element for stabilizing the retained austenite that contributes to the improvement in ductility of the steel sheet. Accordingly, the present invention may include 0.25% or more of carbon (C) to achieve such an effect. A preferable content of carbon (C) may exceed 0.25%, may be 0.27% or more, and may be 0.30% or more. The more preferable content of carbon (C) may be 0.31% or more. On the other hand, when the content of carbon (C) exceeds a certain level, cold rolling may become difficult due to an excessive increase in strength. Therefore, an upper limit of the content of carbon (C) of the present disclosure may be limited to 0.75%. The content of carbon (C) may be 0.70% or less, and the more preferable content of carbon (C) may be 0.67% or less.
Silicon (Si): 4.0% or less (excluding 0%)
Silicon (Si) is an element that contributes to improvement in strength by solid solution strengthening, and is also an element that improves workability by strengthening ferrite and homogenizing a structure. In addition, silicon (Si) is an element contributing to a generation of the retained austenite by suppressing precipitation of cementite. Therefore, in the present invention, silicon (Si) may be necessarily added to achieve such an effect. The preferable content of silicon (Si) may be 0.02% or more, and the more preferable content of silicon (Si) may be 0.05% or more. However, when the content of silicon (Si) exceeds a certain level, a problem of plating defects, such as non-plating, may be induced during plating, and weldability of a steel sheet may be lowered, so the present invention may limit the upper limit of the silicon (Si) content to 4.0%. The preferable upper limit of the content of silicon (Si) may be 3.8%, and the more preferable upper limit of the content of silicon (Si) may be 3.5%.
Aluminum (Al): 5.0% or less (excluding 0%)
Aluminum (Al) is an element performing deoxidation by combining with oxygen in steel. In addition, aluminum (Al) is also an element for stabilizing the retained austenite by suppressing precipitation of cementite like silicon (Si). Therefore, in the present invention, aluminum (Al) may be necessarily added to achieve such an effect. A preferable content of aluminum (Al) may be 0.05% or more, and a more preferable content of aluminum (Al) may be 0.1% or more. On the other hand, when aluminum (Al) is excessively added, inclusions in a steel sheet increase, and the workability of the steel sheet may be lowered, so the present invention may limit the upper limit of the content of aluminum (Al) to 5.0%. The preferable upper limit of the content of aluminum (Al) may be 4.75%, and the more preferable upper limit of the content of aluminum (Al) may be 4.5%.
Meanwhile, the total content (Si+Al) of silicon (Si) and aluminum (Al) is preferably 1.0 to 6.0%. Since silicon (Si) and aluminum (Al) are components that affect microstructure formation in the present invention, and thus, affect ductility, bending formability, and hole expansion ratio, the total content of silicon (Si) and aluminum (Al) is preferably 1.0 to 6.0%. The more preferable total content (Si+Al) of silicon (Si) and aluminum (Al) may be 1.5% or more, and may be 4.0% or less.
Manganese (Mn): 0.9 to 5.0%
Manganese (Mn) is a useful element for increasing both strength and ductility. Therefore, in the present disclosure, a lower limit of a content of manganese (Mn) may be limited to 0.9% in order to achieve such an effect. A preferable lower limit of the content of manganese (Mn) may be 1.0%, and a more preferable lower limit of the content of manganese (Mn) may be 1.1%. On the other hand, when manganese (Mn) is excessively added, the bainite transformation time increases and a concentration of carbon (C) in the austenite becomes insufficient, so there is a problem in that the desired austenite fraction may not be secured. Therefore, an upper limit of the content of manganese (Mn) of the present disclosure may be limited to 5.0%. A preferable upper limit of the content of manganese (Mn) may be 4.7%, and a more preferable upper limit of the content of manganese (Mn) may be 4.5%.
Phosphorus (P): 0.15% or less (including 0%)
Phosphorus (P) is an element that is contained as an impurity and deteriorates impact toughness. Therefore, it is preferable to manage the content of phosphorus (P) to 0.15% or less.
Sulfur (S): 0.03% or less (including 0%)
Sulfur (S) is an element that is included as an impurity to form MnS in a steel sheet and deteriorate ductility. Therefore, the content of sulfur (S) is preferably 0.03% or less.
Nitrogen (N): 0.03% or less (including 0%)
Nitrogen (N) is an element that is included as an impurity and forms nitride during continuous casting to cause cracks of slab. Therefore, the content of nitrogen (N) is preferably 0.03% or less.
Meanwhile, the steel sheet of the present invention has an alloy composition that may be additionally included in addition to the above-described alloy components, which will be described in detail below.
One or more of titanium (Ti): 0 to 0.5%, niobium (Nb): 0 to 0.5%, and vanadium (V): 0 to 0.5%
Titanium (Ti), niobium (Nb), and vanadium (V) are elements that make precipitates and refine crystal grains, and are elements that also contribute to the improvement in strength and impact toughness of a steel sheet, and therefore, in the present invention, one or more of titanium (Ti), niobium (Nb), and vanadium (V) may be added to achieve such an effect. However, when the content of titanium (Ti), niobium (Nb), and vanadium (V) exceed a certain level, respectively, excessive precipitates are formed to lower impact toughness and increase manufacturing cost, so the present invention may limit the content of titanium (Ti), niobium (Nb), and vanadium (V) to 0.5% or less, respectively.
One or more of chromium (Cr): 0 to 3.0% and molybdenum (Mo): 0 to 3.0%
Since chromium (Cr) and molybdenum (Mo) are elements that not only suppress austenite decomposition during alloying treatment, but also stabilize austenite like manganese (Mn), the present invention may add one or more of chromium (Cr) and molybdenum (Mo) to achieve such an effect. However, when the content of chromium (Cr) and molybdenum (Mo) exceeds a certain level, the bainite transformation time increases and the concentration of carbon (C) in austenite becomes insufficient, so the desired retained austenite fraction may not be secured. Therefore, the present invention may limit the content of chromium (Cr) and molybdenum (Mo) to 3.0% or less, respectively.
One or more of Cu: 0 to 4.5% and Ni: 0 to 4.5%
Copper (Cu) and nickel (Ni) are elements that stabilize austenite and suppress corrosion. In addition, copper (Cu) and nickel (Ni) are also elements that are concentrated on a surface of a steel sheet to prevent hydrogen from intruding into the steel sheet, to thereby suppress hydrogen delayed destruction. Accordingly, in the present invention, one or more of copper (Cu) and nickel (Ni) may be added to achieve such an effect. However, when the content of copper (Cu) and nickel (Ni) exceeds a certain level, not only excessive characteristic effects, but also an increase in manufacturing cost is induced, so the present invention may limit the content of copper (Cu) and nickel (Ni) to 4.5% or less, respectively.
Boron (B): 0 to 0.005%
Boron (B) is an element that improves hardenability to increase strength, and is also an element that suppresses nucleation of grain boundaries. Therefore, in the present invention, boron (B) may be added to achieve such an effect. However, when the content of boron (B) exceeds a certain level, not only excessive characteristic effects, but also an increases in manufacturing cost is induced, so the present invention may limit the content of boron (B) to 0.005% or less.
One or more of calcium (Ca): 0 to 0.05%, Magnesium (Mg): 0 to 0.05%, and rare earth element (REM) excluding yttrium (Y): 0 to 0.05%
Here, the rare earth element (REM) is scandium (Sc), yttrium (Y), and a lanthanide element. Since calcium (Ca), magnesium (Mg), and the rare earth element (REM) excluding yttrium (Y) are elements that contribute to the improvement in ductility of a steel sheet by spheroidizing sulfides, in the present invention, one or more of calcium (Ca), magnesium (Mg), and the rare earth element (REM) excluding yttrium (Y) may be added to achieve such an effect. However, when the content of calcium (Ca), magnesium (Mg), and the rare earth element (REM) excluding yttrium (Y) exceeds a certain level, not only excessive characteristic effects, but also an increase in manufacturing cost are induced, so the present invention may limit the content of calcium (Ca), magnesium (Mg), and the rare earth element (REM) excluding yttrium (Y) to 0.05% or less, respectively.
One or more of tungsten (W): 0 to 0.5% and zirconium (Zr): 0 to 0.5%
Since tungsten (W) and zirconium (Zr) are elements that increase strength of a steel sheet by improving hardenability, in the present invention, one or more of tungsten (W) and zirconium (Zr) may be added to achieve such an effect. However, when the content of tungsten (W) and zirconium (Zr) exceeds a certain level, not only excessive characteristic effects, but also an increase in manufacturing cost are induced, so the present invention may limit the content of tungsten (W) and zirconium (Zr) to 0.5% or less, respectively.
One or more of antimony (Sb): 0 to 0.5% and tin (Sn): 0 to 0.5%
Since antimony (Sb) and tin (Sn) are elements that improve plating wettability and plating adhesion of a steel sheet, in the present invention, one or more of antimony (Sb) and tin (Sn) may be added to achieve such an effect. However, when the content of antimony (Sb) and tin (Sn) exceeds a certain level, brittleness of a steel sheet increases, and thus, cracks may occur during hot working or cold working, so the present invention may limit the content of antimony (Sb) and tin (Sn) to 0.5% or less, respectively.
One or more of yttrium (Y): 0 to 0.2% and hafnium (Hf): 0 to 0.2%
Since yttrium (Y) and hafnium (Hf) are elements that improve corrosion resistance of a steel sheet, in the present invention, one or more of the yttrium (Y) and hafnium (Hf) may be added to achieve such an effect. However, when the content of yttrium (Y) and hafnium (Hf) exceeds a certain level, the ductility of the steel sheet may deteriorate, so the present invention may limit the content of yttrium (Y) and hafnium (Hf) to 0.2% or less, respectively.
Cobalt (Co): 0 to 1.5%
Since cobalt (Co) is an element that promotes bainite transformation to increase a TRIP effect, in the present invention, cobalt (Co) may be added to achieve such an effect. However, when the content of cobalt (Co) exceeds a certain level, since weldability and ductility of a steel sheet may deteriorate, the present invention may limit the content of cobalt (Co) to 1.5% or less.
The high strength steel sheet having excellent workability according to an aspect of the present disclosure may include a balance of Fe and other unavoidable impurities in addition to the components described above. However, in a general manufacturing process, unintended impurities may inevitably be mixed from a raw material or the surrounding environment, and thus, these impurities may not be completely excluded. Since these impurities are known to those skilled in the art, all the contents are not specifically mentioned in the present specification. In addition, additional addition of effective components other than the above-described components is not entirely excluded.
The high strength steel sheet having excellent workability according to an aspect of the present invention may include, as microstructures, ferrite which is a soft structure, and tempered martensite, bainite, and retained austenite which are hard structures. Here, the soft structure and the hard structure may be interpreted as a concept distinguished by a relative hardness difference.
As a preferred example, the microstructure of the high strength steel sheet having excellent workability according to an aspect of the present invention may include, by volume fraction, 30 to 70% of tempered martensite, 10 to 45% of bainite, 10 to 40% of retained austenite, 3 to 20% of ferrite, and an unavoidable structure. As the unavoidable structure of the present invention, fresh martensite, perlite, martensite austenite constituent (M-A), and the like may be included. When the fresh martensite or the pearlite is excessively formed, the workability of the steel sheet may be lowered or the fraction of the retained austenite may be lowered.
The high strength steel sheet having excellent workability according to an aspect of the present invention, as shown in the following [Relational Expression 1], a ratio of an average nanohardness value ([H]F, Hv) of the soft structure (ferrite) to an average nanohardness value ([H]TM+B+γ, Hv) of the hard structure (tempered martensite, bainite, and retained austenite) may satisfy a range of 0.4 to 0.9.
0.4≤[H]F/[H]TM+B+γ≤0.9 [Relational Expression 1]
The nanohardness values of the hard and soft structures may be measured using a nanoindenter (FISCHERSCOPE HM2000). Specifically, after electropolishing the surface of the steel sheet, the hard and soft structures are randomly measured at 20 points or more under the condition of an indentation load of 10,000 μN, and the average nanohardness value of the hard and soft structures may be calculated based on the measured values.
In the high strength steel sheet having excellent workability according to an aspect of the present invention, as shown in the following [Relational Expression 2], a ratio of a fraction of retained austenite (V(1.2 μm, γ), vol %) having an average grain size of 1.2 μm or more to a fraction (V(γ), vol %) of retained austenite of the steel sheet may be 0.1 or more. As shown in the following [Relational Expression 3], the ratio of the fraction (V(lath, γ), vol %) of the retained austenite in lath form to the fraction (V(γ), vol %) of the retained austenite of the steel sheet may be 0.5 or more.
V(1.2 μm, γ)/V(γ)≥0.1 [Relational Expression 2]
V(lath, γ)/V(γ)≥0.5 [Relational Expression 3]
In the high strength steel sheet having excellent workability according to an aspect of the present invention, since a balance BT·E of tensile strength and elongation expressed by the following [Relational Expression 4] is 22,000 (MPa %) or more, a balance BT·H of tensile strength and hole expansion ratio expressed by the following [Relational Expression 5] is 7*106 (MPa2%1/2) or more, and bendability BR expressed by the following [Relational Expression 6] satisfies a range of 0.5 to 3.0, it may have an excellent balance of strength and ductility, a balance of strength and a hole expansion ratio, and superb bending formability.
BT·E=[Tensile Strength (TS, MPa)]*[Elongation (El, %)] [Relational Expression 4]
BT·H=[Tensile Strength (TS, MPa)]2*[Hole Expansion Ratio (HER, %)]1/2 [Relational Expression 5]
BR=R/t [Relational Expression 6]
In the above Relational Expression 6, R is a minimum bending radius (mm) at which cracks do not occur after a 90° bending test, and t is a thickness (mm) of the steel sheet.
In the present invention, it is important to stabilize retained austenite of a steel sheet because it is intended to simultaneously secure superb ductility and bending formability as well as high strength properties. In order to stabilize the retained austenite, it is necessary to concentrate carbon (C) and manganese (Mn) in the ferrite, bainite, and tempered martensite of the steel sheet into austenite. However, when carbon (C) is concentrated into austenite by using ferrite, the strength of the steel sheet may be insufficient due to the low strength characteristics of the ferrite, and the excessive interphase hardness difference may occur, thereby reducing the hole expansion ratio (HER). Therefore, it is intended to concentrate carbon (C) and manganese (Mn) into austenite by using the bainite and tempered martensite.
When the content of silicon (Si) and aluminum (Al) in the retained austenite is limited to a certain range, carbon (C) and manganese (Mn) may be concentrated in large amounts from bainite and tempered martensite into retained austenite, thereby effectively stabilizing the retained austenite. In addition, by limiting the content of silicon (Si) and aluminum (Al) in austenite to a certain range, it is possible to increase the content of silicon (Si) and aluminum (Al) in ferrite. As the content of silicon (Si) and aluminum (Al) in the ferrite increases, the hardness of the ferrite increases, so it is possible to effectively reduce an interphase hardness difference between ferrite which is a soft structure and tempered martensite, bainite, and retained austenite which are hard structures.
When the ratio of the average nanohardness value ([H]F, Hv) of the soft structure (ferrite) to the average nanohardness value ([H]TM+B+γ, Hv) of the hard structure (tempered martensite, bainite, and retained austenite) is greater than a certain level, the interphase hardness difference between the soft structure (ferrite) and the hard structure (tempered martensite, bainite, and retained austenite) is lowered, so it is possible to secure a desired balance (TSXE1) of tensile strength and elongation, a balance (TS2XHER1/2) of tensile strength and hole expansion ratio, and bendability (R/t). On the other hand, when the ratio of the average nanohardness value ([H]F, Hv) of the soft structure (ferrite) to the average nanohardness value ([H]TM+B+γ, Hv) of the hard structure (tempered martensite, bainite, and retained austenite) is excessive, the ferrite is excessively hardened and the workability is rather lowered, so the desired balance (TSXE1) of tensile strength and elongation, the balance of tensile strength and hole expansion ratio (TS2XHER1/2), and the bendability (R/t) may not all be secured. Therefore, the present invention may limit the ratio of the average nanohardness value ([H]F, Hv) of the soft structure to the average nanohardness value ([H]TM+B+γ, Hv) of the hard structure (tempered martensite, bainite, and retained austenite) to a range of 0.4 to 0.9.
In the retained austenite, retained austenite having an average grain size of 1.2 μm or more may be heat-treated at a bainite formation temperature to increase an average size in order to inhibit transformation from austenite to martensite, thereby improving the workability of the steel sheet.
In addition, in the retained austenite, retained austenite in a lath form affects the workability of the steel sheet. The retained austenite is divided into retained austenite in a lath form which is formed between bainite phases and retained austenite in a block form which is formed in a portion without bainite phases. As the retained austenite in the block form is additionally transformed into bainite during the heat treatment, the retained austenite in lath form increases, thereby effectively improving the processing of the steel sheet.
Therefore, in order to improve the ductility and workability of the steel sheet, it is preferable to increase the fraction of the retained austenite having an average grain size of 1.2 μm or more and the fraction of the retained austenite in lath form, in the retained austenite.
In the high strength steel sheet having excellent workability according to an aspect of the present invention, the ratio of the fraction of the retained austenite (V(1.2 μm, γ), vol %) having an average grain size of 1.2 μm or more to the fraction (V(γ), vol %) of the retained austenite of the steel sheet may be limited to 0.1 or more, and the ratio of the fraction (V(lath, γ), vol %) of the retained austenite in lath form to the fraction (V(γ), vol %) of the retained austenite of the steel sheet may be limited to 0.5 or more. When the ratio of the fraction (V(1.2 μm, γ), vol %) of the retained austenite having an average grain size of 1.2 μm or more to the fraction (V(γ), vol %) of the retained austenite of the steel sheet is less than 0.1 or the ratio of the fraction (V(lath, γ), vol %) of the retained austenite in lath form to the fraction (V(γ), vol %) of the retained austenite of the steel sheet is less than 0.5, the bendability (R/t) does not satisfy 0.5 to 3.0, so there is a problem in that the desired workability may not be secured.
A steel sheet including retained austenite has superb ductility and bending formability due to transformation-induced plasticity occurring during transformation from austenite to martensite during processing. When the fraction of the retained austenite is lower than a certain level, the balance (TSXE1) of tensile strength and elongation may be less than 22,000 MPa %, or the bendability (R/t) may exceed 3.0. Meanwhile, when the fraction of the retained austenite exceeds a certain level, local elongation may be lowered. Accordingly, in the present invention, the fraction of the retained austenite may be limited to a range of 10 to 40 vol % in order to obtain a steel sheet having a balance (TSXE1) of tensile strength and elongation and superb bendability (R/t).
Meanwhile, both untempered martensite (fresh martensite) and tempered martensite are microstructures that improve the strength of the steel sheet. However, compared with the tempered martensite, fresh martensite has a characteristic of greatly reducing the ductility and the hole expansion ratio of the steel sheet. This is because the microstructure of the tempered martensite is softened by the tempering heat treatment. Therefore, in the present invention, it is preferable to use tempered martensite to provide a steel sheet having a balance of strength and ductility, a balance of strength and hole expansion ratio, and superb bending formability. When the fraction of the tempered martensite is less than a certain level, it is difficult to secure the balance (TSXE1) of tensile strength and elongation of 22,000 MPa % or more or the balance (TS2XHER1/2) of tensile strength and hole expansion ratio of 7*106 (MPa2%1/2) or more, and when the fraction of the tempered martensite exceeds a certain level, ductility and workability is lowered, and the balance (TSXE1) of tensile strength and elongation is less than 22,000 MPa %, or bendability (R/t) exceeds 3.0, which is not preferable. Therefore, in the present invention, the fraction of the tempered martensite may be limited to 30 to 70 vol % to obtain a steel sheet having the balance (TSXE1) of tensile strength and elongation, the balance (TS2XHER1/2) of tensile strength and hole expansion ratio, and superb bendability (R/t).
In order to improve the balance (TSXE1) of tensile strength and elongation, the balance (TS2XHER1/2) of tensile strength and hole expansion ratio, and the bendability (R/t), it is preferable that bainite is appropriately included as the microstructure. As long as a fraction of bainite is a certain level or more, it is possible to secure the balance (TSXE1) of tensile strength and elongation of 22,000 MPa % or more, the balance (TS2XHER1/2) of tensile strength and hole expansion ratio of 7*106 (MPa2%1/2) or more and the bendability (R/t) of 0.5 to 3.0. On the other hand, when the fraction of bainite is excessive, the decrease in the fraction of tempered martensite is necessarily accompanied, so the present invention may not secure the desired balance (TSXE1) of tensile strength and elongation, the balance (TS2XHER1/2) of tensile strength and hole expansion ratio, and bendability (R/t). Accordingly, the present invention may limit the fraction of bainite to a range of 10 to 45 vol %.
Since ferrite is an element contributing to improvement in ductility, the present invention may secure the desired balance (TSXE1) of tensile strength and elongation, as long as the fraction of ferrite is a certain level or more. However, when the fraction of ferrite is excessive, the interphase hardness difference increases and the hole expansion ratio (HER) may decrease, so the present invention may not secure the desired balance (TS2XHER1/2) of tensile strength and hole expansion ratio. Accordingly, the present invention may limit the fraction of ferrite to a range of 3 to 20 vol %.
Hereinafter, an example of a method for manufacturing a steel sheet of the present invention will be described in detail.
A method for manufacturing a high strength steel sheet having excellent workability according to an aspect of the present invention may include: providing a cold-rolled steel sheet having a predetermined component; heating (primary heating) the cold-rolled steel sheet to a temperature within a range of Ac1 or higher and less than Ac3, and holding (primary holding) the cold-rolled steel sheet for 50 seconds or more; cooling (primary cooling) the cold-rolled steel sheet to a temperature within a range of 600 to 850° C. (primary cooling stop temperature) at an average cooling rate of 1° C./s or more; cooling (secondary cooling) the cold-rolled steel sheet to a temperature within a range of 300 to 500° C. at an average cooling rate of 2° C./s or more, and holding (secondary holding) the cold-rolled steel sheet in the temperature within a range for 5 seconds or more; cooling (tertiary cooling) the cold-rolled steel sheet to a temperature within a range of 100 to 300° C. (secondary cooling stop temperature) at an average cooling rate of 2° C./s or more; heating (secondary heating) the cold-rolled steel sheet to a temperature within a range of 350 to 550° C., and holding (tertiary holding) the cold-rolled steel sheet in the temperature within a range for 10 seconds or more; cooling (quaternary cooling) the cold-rolled steel sheet to a temperature within a range of 250 to 450° C., and holding (quaternary holding) the cold-rolled steel sheet in the temperature within a range for 10 seconds or more; cooling (fifth cooling) the cold-rolled steel sheet to room temperature.
In addition, the cold-rolled steel sheet of the present invention may be provided by heating a steel slab to 1000 to 1350° C.; performing finishing hot rolling in a temperature within a range of 800 to 1000° C.; coiling the hot-rolled steel sheet at a temperature within a range of 300 to 600° C.; performing hot-rolled annealing heat treatment on the coiled steel sheet in a temperature within a range of 650 to 850° C. for 600 to 1700 seconds; and cold rolling the hot-rolled annealing heat-treated steel sheet at a reduction ratio of 30 to 90%.
Preparation and Heating of Steel Slab
A steel slab having a predetermined component is prepared. Since the steel slab according to the present invention includes an alloy composition corresponding to an alloy composition of the steel sheet described above, the description of the alloy compositions of the slab is replaced by the description of the alloy composition of the steel sheet described above.
The prepared steel slab may be heated to a certain temperature within a range, and the heating temperature of the steel slab at this time may be in the range of 1000 to 1350° C. This is because, when the heating temperature of the steel slab is less than 1000° C., the steel slab may be hot rolled in the temperature within a range below the desired finish hot rolling temperature within a range, and when the heating temperature of the steel slab exceeds 1350° C., the temperature reaches a melting point of steel, and thus, the steel slab is melted.
Hot Rolling and Coiling
The heated steel slab may be hot rolled, and thus, provided as a hot-rolled steel sheet. During the hot rolling, the finish hot rolling temperature is preferably in the range of 800 to 1000° C. When the finish hot rolling temperature is less than 800° C., an excessive rolling load may be a problem, and when the finish hot rolling temperature exceeds 1000° C., grains of the hot-rolled steel sheet are coarsely formed, which may cause a deterioration in physical properties of the final steel sheet.
The hot-rolled steel sheet after the hot rolling has been completed may be cooled at an average cooling rate of 10° C./s or more, and may be coiled at a temperature of 300 to 600° C. When the coiling temperature is less than 300° C., the coiling is not easy, and when the coiling temperature exceeds 600° C., a surface scale is formed to the inside of the hot-rolled steel sheet, which may make pickling difficult.
Hot-Rolled Annealing Heat Treatment
It is preferable to perform a hot-rolled annealing heat treatment process in order to facilitate pickling and cold rolling, which are subsequent processes after the coiling. The hot-rolled annealing heat treatment may be performed in a temperature within a range of 650 to 850° C. for 600 to 1700 seconds. When the hot-rolled annealing heat treatment temperature is less than 650° C. or the hot-rolled annealing heat treatment time is less than 600 seconds, the strength of the hot-rolled annealing heat-treated steel sheet increases, and thus, subsequent cold rolling may not be easy. On the other hand, when the hot-rolled annealing heat treatment temperature exceeds 850° C. or the hot-rolled annealing heat treatment time exceeds 1700 seconds, the pickling may not be easy due to a scale formed deep inside the steel sheet.
Pickling and Cold Rolling
After the hot-rolled annealing heat treatment, in order to remove the scale generated on the surface of the steel sheet, the pickling may be performed, and the cold rolling may be performed. Although the conditions of the pickling and cold rolling are not particularly limited in the present invention, the cold rolling is preferably performed at a cumulative reduction ratio of 30 to 90%. When the cumulative reduction ratio of the cold rolling exceeds 90%, it may be difficult to perform the cold rolling in a short time due to the high strength of the steel sheet.
The cold-rolled steel sheet may be manufactured as a non-plated cold-rolled steel sheet through the annealing heat treatment process, or may be manufactured as a plated steel sheet through a plating process to impart corrosion resistance. As the plating, plating methods such as hot-dip galvanizing, electro-galvanizing, and hot-dip aluminum plating may be applied, and the method and type are not particularly limited.
Annealing Heat Treatment
In the present invention, in order to simultaneously secure the strength and workability of the steel sheet, the annealing heat treatment process is performed.
The cold-rolled steel sheet is heated (primarily heated) to a temperature within a range of Ac1 or higher and less than Ac3 (two-phase region), and held (primarily held) in the temperature within a range for 50 seconds or more. The primary heating or primary holding temperature is Ac3 or higher (single-phase region), the desired ferrite structure may not be realized, so the desired level of [H]F/[H]TM+B+γ, and the balance (TS2XHER1/2) of tensile strength and hole expansion ratio may be implemented. In addition, when the primary heating or primary holding temperature is in a temperature within a range less than Ac1, there is a fear that sufficient heating is not made, and thus, the microstructure desired by the present invention may not be implemented even by subsequent heat treatment. The average temperature increase rate of the primary heating may be 5° C./s or more.
When the primary holding time is less than 50 seconds, the structure may not be sufficiently homogenized and the physical properties of the steel sheet may be lowered. The upper limit of the primary holding time is not particularly limited, but the primary heating time is preferably limited to 1200 seconds or less in order to prevent the decrease in toughness due to the coarsening of grains.
After the primary holding, it is preferable to cool (primarily cool) the cold-rolled steel sheet to a temperature within a range (primary cooling stop temperature) of 600 to 850° C. at an average cooling rate of 1° C./s or more. The upper limit of the average cooling rate of the primary cooling does not need to be particularly specified, but is preferably limited to 100° C. or lower. When the primary cooling stop temperature is less than 600° C., the ferrite is excessively formed and the retained austenite is insufficient, and [H]F/[H]TM+B+γ and the balance (TSXE1) between tensile strength and elongation may be lowered. In addition, since it is preferable that the upper limit of the primary cooling stop temperature is 30° C. or lower than the primary holding temperature, the upper limit of the primary cooling stop temperature may be limited to 850° C.
After the primary cooling, it is preferable to cool (secondarily cool) the cold-rolled steel sheet to a temperature within a range of 300 to 500° C. at an average cooling rate of 2° C./s or more, and to hold (secondarily hold) the cold-rolled steel sheet in the temperature within a range for 5 seconds or more. When the average cooling rate of the secondary cooling is less than 2° C./s, the ferrite is excessively formed and the retained austenite is insufficient, so [H]F/[H]TM+B+γ and the balance (TSXE1) of tensile strength and elongation may be lowered. The upper limit of the average cooling rate of the secondary cooling does not need to be particularly specified, but is preferably limited to 100° C./s or less. Meanwhile, when the secondary holding temperature exceeds 500° C., the retained austenite is insufficient, so [H]F/[H]TM+B+γ, V(lath, γ)/V(γ), the balance (TSXE1) of tensile strength and elongation, and the bendability (R/t) may be lowered. In addition, when the secondary holding temperature is less than 300° C., V(1.2 μm, γ)/V(γ) and the bendability (R/t) may be lowered due to the low heat treatment temperature. When the secondary holding time is less than 5 seconds, V(lath, γ)/V(γ), and the bendability (R/t) may be lowered due to the insufficient heat treatment time. On the other hand, the upper limit of the secondary holding time does not need to be particularly specified, but is preferably set to 600 seconds or less.
Meanwhile, it is preferable that the average cooling rate Vc1 of the primary cooling is smaller than the average cooling rate Vc2 of the secondary cooling (Vc1<Vc2).
After the secondary holding, it is preferable to cool (tertiarily cool) the cold-rolled steel sheet to a temperature within a range (secondary cooling stop temperature) of 100 to 300° C. at an average cooling rate of 2° C./s or more. When the average cooling rate of the tertiary cooling is less than 2° C./s, V(1.2 μm, γ)/V(γ) and bendability (R/t) may be lowered due to slow cooling. The upper limit of the average cooling rate of the tertiary cooling does not need to be particularly specified, but is preferably limited to 100° C./s or less. Meanwhile, when the secondary cooling stop temperature exceeds 300° C., the bainite is excessively formed and the tempered martensite is insufficient, so the balance (TSXE1) of tensile strength and elongation may be lowered. On the other hand, when the secondary cooling stop temperature is less than 100° C., the tempered martensite is excessively formed and the retained austenite is insufficient, so [H]F/[H]TM+B+γ, V(1.2 μm, γ)/V(γ), the balance (TSXE1) of tensile strength and elongation, and the bendability (R/t) may be lowered.
After the tertiary cooling, it is preferable to heat (secondarily heat) the cold-rolled steel sheet to a temperature within a range of 350 to 550° C., and hold (tertiarily hold) the cold-rolled steel sheet in the temperature within a range for 10 seconds or more. When the tertiary holding temperature exceeds 500° C., the retained austenite is insufficient, so [H]F/[H]TM+B+γ, V(lath, γ)/V(γ), the balance (TSXE1) of tensile strength and elongation, and the bendability (R/t) may be lowered. On the other hand, when the tertiary holding temperature is less than 350° C., V(1.2 μm, γ)/V(γ) and the bendability (R/t) may be lowered due to the low holding temperature. When the tertiary holding time is less than 10 seconds, V(lath, γ)/V(γ), and the bendability (R/t) may be lowered due to the insufficient holding time. The upper limit of the tertiary holding time is not particularly limited, but a preferred tertiary holding time may be 1800 seconds or less.
After the tertiary holding, it is preferable to cool (quaternary cool) the cold-rolled steel sheet to a temperature within a range of 250 to 450° C. at an average cooling rate of 1° C./s or more, and to hold (quaternarily hold) the cold-rolled steel sheet in the temperature within a range for 10 seconds or more. When the average cooling rate of the quaternary cooling is less than 1° C./s, V(1.2 μm, γ)/V(γ) and the bendability (R/t) may be lowered due to the slow cooling. The upper limit of the average cooling rate of the quaternary cooling does not need to be particularly specified, but is preferably limited to 100° C./s or less. When the quaternary holding temperature exceeds 450° C., V(lath, γ)/V(γ), and the bendability (R/t) may be lowered due to the heat treatment for a long time. On the other hand, when the quaternary holding temperature is less than 250° C., V(lath, γ)/V(γ), and the bendability (R/t) may be lowered due to the low holding temperature. When the quaternary holding time is less than 10 seconds, V(lath, γ)/V(γ), and the bendability (R/t) may be lowered due to the insufficient holding time. The upper limit of the quaternary holding time is not particularly limited, but a preferred quaternary holding time may be 176,000 seconds or less.
After the quaternary holding, it is preferable to cool (fifth cool) the cold-rolled steel sheet to room temperature at an average cooling rate of 1° C./s or more.
The high strength steel sheet having excellent workability manufactured by the above-described manufacturing method may include, as a microstructure, tempered martensite, bainite, retained austenite, and ferrite, and as a preferred example, may include, by the volume fraction, 30 to 70% of tempered martensite, 10 to 45% of bainite, 10 to 40% of retained austenite, 3 to 20% of ferrite, and unavoidable structures.
In the high strength steel sheet having excellent workability manufactured by the above-described manufacturing method, as shown in the following [Relational Expression 1], the ratio of the average nanohardness value ([H]F, Hv) of the soft structure (ferrite) to the average nanohardness value ([H]TM+B+γ, Hv) of the hard structure (tempered martensite, bainite, and retained austenite) may satisfy the range of 0.4 to 0.9, and, as shown in the following [Relational Expression 2], the ratio of the fraction of retained austenite having an average grain size of 1.2 μm or more to the fraction of retained austenite of the steel sheet may satisfy 0.1 or more.
0.4≤[H]F/[H]TM+Bγ≤0.9 [Relational Expression 1]
v(1.2 μm, γ)/V(γ)≥0.1 [Relational Expression 2]
In addition, in the high strength steel sheet having excellent workability manufactured by the above-described manufacturing method, as shown in the following [Relational Expression 3], the fraction (V(lath, γ), vol %) of the retained austenite in lath form to the fraction (V(γ), vol %) of the retained austenite of the steel sheet may be 0.5 or more.
V(lath, γ)/V(γ)≥0.5 [Relational Expression 3]
In the high-strength steel sheet having excellent workability manufactured by the above-described manufacturing method, a balance BT·E of tensile strength and elongation expressed by the following [Relational Expression 4] is 22,000 (MPa %), a balance BT·H of tensile strength and hole expansion ratio expressed by the following [Relational Expression 5] is 7*106 (MPa2%1/2) or more, and bendability BR expressed by the following [Relational Expression 6] may satisfy a range of 0.5 to 3.0.
BT·E=[Tensile Strength (TS, MPa)]*[Elongation (EL, %)] [Relational Expression 4]
BT·H=[Tensile Strength (TS, MPa)]2*[Hole Extension Ratio (HER, %)]1/2 [Relational Expression 5]
BR=R/t [Relational Expression 6]
In the above Relational Expression 6, R is a minimum bending radius (mm) at which cracks do not occur after a 90° bending test, and t is a thickness (mm) of the steel sheet.
Hereinafter, a high strength steel sheet having excellent workability and a method for manufacturing same according to an aspect of the present invention will be described in more detail. It should be noted that the following examples are only for the understanding of the present invention, and are not intended to specify the scope of the present invention. The scope of the present invention is determined by matters described in claims and matters reasonably inferred therefrom.
A steel slab having a thickness of 100 mm having alloy compositions (a balance of Fe and unavoidable impurities) shown in Table 1 below was prepared, heated at 1200° C., and then was subjected to finish hot rolling at 900° C. Thereafter, the steel slab was cooled at an average cooling rate of 30° C./s, and coiled at a coiling temperature of Tables 2 and 3 to manufacture a hot-rolled steel sheet having a thickness of 3 mm. The hot-rolled steel sheet was subjected to hot-rolled annealing heat treatment under the conditions of Tables 2 and 3. Thereafter, after removing a surface scale by pickling, cold rolling was performed to a thickness of 1.5 mm.
Thereafter, the heat treatment was performed under the annealing heat treatment conditions disclosed in Tables 2 to 7 to manufacture the steel sheet.
The microstructure of the thus prepared steel sheet was observed, and the results were shown in Tables 8 and 9. Among the microstructures, ferrite (F), bainite (B), tempered martensite (TM), and pearlite (P) were observed through SEM after nital-etching a polished specimen cross section. The fractions of bainite and tempered martensite, which are difficult to distinguish among them, were calculated using an expansion curve after evaluation of dilatation. Meanwhile, since fresh martensite (FM) and retained austenite (retained γ) are also difficult to distinguish, a value obtained by subtracting the fraction of retained austenite calculated by X-ray diffraction method from the fraction of martensite and retained austenite observed by the SEM was determined as the fraction of the fresh martensite.
Meanwhile, [H]F/[H]TM+B+γ, V(lath, γ)/V(γ), V(1.2 μm, γ)/V(γ), a balance (TSXE1) of tensile strength and elongation, a balance (TS2XHER1/2) of tensile strength and hole expansion ratio, and bendability (R/t) were observed, and the results were shown in Tables 10 and 11.
Nanohardness values of hard and soft structures were measured using the nanoindentation method. Specifically, after electropolishing surfaces of each specimen, the hard and soft structures were randomly measured at 20 points or more under the condition of an indentation load of 10,000 μN using a nanoindenter (FISCHERSCOPE HM2000), and the average nanohardness value of the hard and soft structures was calculated based on the measured values.
The retained austenite fraction (V(1.2 μm, γ)) having an average grain size of 1.2/L or more and the fraction (V(lath, γ)) of the retained austenite in lath form were determined by the area measured within the retained austenite phase using a phase map of EPMA.
Tensile strength (TS) and elongation (El) were evaluated through a tensile test, and the tensile strength (TS) and the elongation (El) were measured by evaluating the specimens collected in accordance with JIS No. 5 standard based on a 90° direction with respect to a rolling direction of a rolled sheet. The bendability (R/t) was evaluated by a V-bending test, and calculated by collecting a specimen based on the 90° direction with respect to the rolling direction of the rolled sheet and was determined as a value obtained by dividing a minimum bending radius R, at which cracks do not occur after a 90° bending test, by a thickness t of a sheet. The hole expansion ratio (HER) was evaluated through the hole expansion test, and was calculated by the following [Relational Expression 7] by, after forming a punching hole (die inner diameter of 10.3 mm, clearance of 12.5%) of 10 mmØ, inserting a conical punch having an apex angle of 60° into a punching hole in a direction in which a burr of a punching hole faces outward, and then compressing and expanding a peripheral portion of the punching hole at a moving speed of 20 mm/min.
Hole Expansion Ratio (HER, %)={(D−D0)/D0}×100 [Relational Expression 7]
In the above Relational Expression 5, D is a hole diameter (mm) when cracks penetrate through the steel plate along the thickness direction, and D0 is the initial hole diameter (mm).
As shown in Tables 1 to 9 above, it could be seen that the specimens satisfying the conditions presented in the present invention simultaneously provide strength and workability since the value of [H]F/[H]TM+B+γ satisfies the range of 0.4 to 0.9, and the value of V (lath, γ)/V(γ) satisfies 0.5 or more, the value of V(1.2 μm, γ)/V(γ) satisfies 0.1 or more, the balance (TSXE1) of tensile strength and elongation is 22,000 MPa % or more, the balance (TS2XHER1/2) of tensile strength and hole expansion ratio is 7*106 (MPa2%1/2) or more, and the bendability (R/t) satisfies the range of 0.5 to 3.0.
It could be seen that, in specimens 2 to 5, since the alloy composition range of the present invention overlaps, but the hot-rolled annealing temperature and time are outside the range of the present invention, the pickling failure occurs or the fracture occurs during the cold rolling.
In specimen 6, the amount of ferrite formed was insufficient because the primary heating or holding temperature in the annealing heat treatment process after the cold rolling exceeded (single-phase region) the range limited by the present invention. As a result, it could be seen that, in specimen 6, [H]F/[H]TM+B+γ was less than 0.4, and the balance of tensile strength and hole expansion ratio (TS2XHER1/2) was less than 7*106 (MPa2%1/2).
In specimen 8, the primary cooling stop temperature in the annealing heat treatment process after the cold rolling is low, so ferrite was excessively formed and retained austenite was formed less. As a result, it could be seen that, in specimen 8, [H]F/[H]TM+B+γ exceeds 0.9, and the balance (TSXE1) of tensile strength and elongation is less than 22,000 MPa %.
In Specimen 9, the average cooling rate of the secondary cooling was low, so ferrite was excessively formed and retained austenite was formed less. As a result, it could be seen that, in specimen 9, [H]F/[H]TM+B+γ exceeds 0.9, and the balance (TSXE1) of tensile strength and elongation is less than 22,000 MPa %.
In specimen 11, the secondary holding temperature is high, so the retained austenite was formed less. As a result, it could be seen that, in specimen 12, [H]F/[H]TM+B+γ exceeds 0.9, V(lath, γ)/V(γ) is less than 0.5, the balance (TSXE1) of tensile strength and elongation is less than 22,000 MPa %, and the bendability (R/t) exceeds 3.0.
It could be seen that, in specimen 12, the secondary holding temperature is low, so V(1.2 μm, γ)/V(γ) is less than 0.1 and the bendability (R/t) exceeds 3.0.
It could be seen that, in specimen 13, the secondary holding time is short, so V(lath, γ)/V(γ) is less than 0.5, V(1.2 μm, γ)/V(γ) is less than 0.1, and the bendability (R/t) exceeds 3.0.
It could be seen that, in specimen 14, the average cooling rate of the tertiary cooling is low, so V(1.2 μm, γ)/V(γ) is less than 0.1 and the bendability (R/t) exceeds 3.0.
In Specimen 15, the secondary cooling stop temperature was high, so bainite was excessively formed and tempered martensite was formed less. As a result, it could be seen that the balance (TSXE1) of tensile strength and elongation is less than 22,000 MPa %.
In specimen 16, the secondary cooling stop temperature is low, so the tempered martensite was excessively formed and the retained austenite was formed less. As a result, it could be seen that [H]F/[H]TM+B+γ exceeds 0.9, V(1.2 μm, γ)/V(γ) is less than 0.1, the balance (TSXE1) of tensile strength and elongation is less than 22,000 MPa %, and the bendability (R/t) exceeds 3.0.
In specimen 17, the tertiary holding temperature is high, so the retained austenite was formed less. It could be seen that [H]F/[H]TM+B+γ exceeds 0.9, V(lath, γ)/V(γ) is less than 0.5, the balance (TSXE1) of tensile strength and elongation is less than 22,000 MPa %, and the bendability (R/t) exceeds 3.0.
It could be seen that, in specimen 18, the tertiary holding temperature is low, so V(1.2 μm, γ)/V(γ) is less than 0.1 and the bendability (R/t) exceeds 3.0.
It could be seen that, in specimen 19, the tertiary holding time is short, so V(lath, γ)/V(γ) is less than 0.5, V(1.2 μm, γ)/V(γ) is less than 0.1, and the bendability (R/t) exceeds 3.0.
It could be seen that, in specimen 20, the quaternary holding temperature is high, so V(lath, γ)/V(γ) is less than 0.5, V(1.2 μm, γ)/V(γ) is less than 0.1, and the bendability (R/t) exceeds 3.0, and in specimen 21, the quaternary holding temperature is high, so V(lath, γ)/V(γ) is less than 0.5, V(1.2 μm, γ)/V(γ) is less than 0.1, and the bendability (R/t) exceeds 3.0.
It could be seen that, in specimen 22, the quaternary holding time is short, so V(lath, γ)/V(γ) is less than 0.5, V(1.2 μm, γ)/V(γ) is less than 0.1, and the bendability (R/t) exceeds 3.0.
Specimens 45 to 53 may satisfy the manufacturing conditions presented in the present invention, but may be outside the alloy composition range. In these cases, it could be seen that the condition of the [H]F/[H]TM+B+γ, the condition of the V(lath, γ)/V(γ), the condition of V(1.2 μm, γ)/V(γ), the condition of the balance (TSXE1) of tensile strength and elongation, the condition of the balance (TS2XHER1/2) of tensile strength and hole expansion ratio, and the condition of bendability (R/t) of the present invention are not all satisfied. Meanwhile, it could be seen that, in specimen 47, when the total content of aluminum (Al) and silicon (Si) is less than 1.0%, the conditions of [H]F/[H]TM+B+γ, the balance (TSXE1) of tensile strength and elongation, and the bendability (R/t) are not satisfied.
While the present invention has been described in detail through exemplary embodiment, other types of exemplary embodiments are also possible. Therefore, the technical spirit and scope of the claims set forth below are not limited to exemplary embodiments.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2019-0169613 | Dec 2019 | KR | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/KR2020/016559 | 11/23/2020 | WO |
| Publishing Document | Publishing Date | Country | Kind |
|---|---|---|---|
| WO2021/125589 | 6/24/2021 | WO | A |
| Number | Name | Date | Kind |
|---|---|---|---|
| 20070190353 | Taniguchi et al. | Aug 2007 | A1 |
| 20110146852 | Matsuda et al. | Jun 2011 | A1 |
| 20130160907 | Kawabe et al. | Jun 2013 | A1 |
| 20130167980 | Kawata et al. | Jul 2013 | A1 |
| 20140170439 | Allain et al. | Jun 2014 | A1 |
| 20150322552 | Takashima et al. | Nov 2015 | A1 |
| 20160160309 | Allain et al. | Jun 2016 | A1 |
| 20160177427 | Takashima et al. | Jun 2016 | A1 |
| 20180298462 | Sano et al. | Oct 2018 | A1 |
| 20190003009 | Kawata et al. | Jan 2019 | A1 |
| 20190071757 | Murata et al. | Mar 2019 | A1 |
| 20190218652 | Haga et al. | Jul 2019 | A1 |
| 20190256945 | Zhou et al. | Aug 2019 | A1 |
| 20200392598 | Tanaka et al. | Dec 2020 | A1 |
| 20200392610 | Tanaka et al. | Dec 2020 | A1 |
| 20210115528 | Magar et al. | Apr 2021 | A1 |
| Number | Date | Country |
|---|---|---|
| 103080357 | May 2013 | CN |
| 103097566 | May 2013 | CN |
| 104870676 | Aug 2015 | CN |
| 2617849 | Jul 2013 | EP |
| 3279363 | Feb 2018 | EP |
| 3438309 | Feb 2019 | EP |
| 3636788 | Apr 2020 | EP |
| 2009019258 | Jan 2009 | JP |
| 2012237042 | Dec 2012 | JP |
| 2015034326 | Feb 2015 | JP |
| 2015105383 | Jun 2015 | JP |
| 2016008310 | Jan 2016 | JP |
| 2017053001 | Mar 2017 | JP |
| 2017186647 | Oct 2017 | JP |
| 2018162495 | Oct 2018 | JP |
| 20060118602 | Nov 2006 | KR |
| 20130032917 | Apr 2013 | KR |
| 101340758 | Dec 2013 | KR |
| 20140012167 | Jan 2014 | KR |
| 20140127570 | Nov 2014 | KR |
| 20180016518 | Feb 2018 | KR |
| 20180088707 | Aug 2018 | KR |
| 20180125560 | Nov 2018 | KR |
| 20190039424 | Apr 2019 | KR |
| 20190040018 | Apr 2019 | KR |
| 20190071755 | Jun 2019 | KR |
| 2016001702 | Jan 2016 | WO |
| 2018221307 | Dec 2018 | WO |
| 2019131188 | Jul 2019 | WO |
| 2019131189 | Jul 2019 | WO |
| Entry |
|---|
| European Search Report—European Application No. 20901119.6 issued on Jul. 13, 2023, citing EP 2617849, JP 2017-053001, EP 3438309, and EP 3279363. |
| Japanese Office Action—Japanese Application No. 2022-536581 issued on Jun. 20, 2023, citing JP 2015-105383, JP 2018-162495, JP 2017-186647, JP 2016-008310, WO 2019/131188, and WO 2019/131189. |
| Chinese Office Action—Chinese Application No. 202080087461.0 issued on Mar. 7, 2023, citing CN 103097566, JP 2012-237042, CN 104870676, and CN 103080357. |
| International Search Report—PCT/KR2020/016559 dated Mar. 2, 2021. |
| European Interview Summary—European Application No. 20901119.6 issued on Sep. 5, 2024, citing WO 2018/221307 and EP 3636788. |
| Number | Date | Country | |
|---|---|---|---|
| 20230025863 A1 | Jan 2023 | US |
