High-Strength Steel Sheet with Excellent Ductility and Crackless Edge Portion, Hot-Dip Galvanized Steel Sheet, and Manufacturing Method Thereof

Info

Publication number: 20110064968
Type: Application
Filed: May 27, 2009
Publication Date: Mar 17, 2011
Applicant: POSCO (Pohang)
Inventors: Sung-Il Kim (Gwangyang), Young-Hoon Jin (Gwangyang), Jai-Hyun Kwak (Gwangyang), Kwang-Geun Chin (Gwangyang)
Application Number: 12/991,003

Abstract

There are provided steel sheets which have a tensile strength of 780-980 MPa and an elongation of 28% or higher and are free of edge cracks. The high-strength steel sheet includes, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfying 1.75%≦Si+Al≦3.25%.

Description

Description

TECHNICAL FIELD

The present invention relates to steel sheets having excellent strength and workability, which are mainly used in various structural members of automobiles, and manufacturing methods thereof, and more particularly, to a high-strength steel sheet and a hot-dipped galvanized steel sheet, which have a tensile strength of 780-980 MPa and an elongation of at least 28% and are free of edge cracks during cold rolling, and methods for manufacturing these steel sheets.

BACKGROUND ART

Steel sheets which are used as structural members among automotive components must be able to effectively absorb external impacts, thereby improving the safety of passengers when vehicles collide. Recently, steel sheets having excellent strength and formability, such as a tensile strength of 780 MPa higher and an elongation of 25% or higher, have been mainly used as such structural members. Also, such steel sheets need to have high tensile strength, high elongation and low yield ratio (ratio of yield strength to tensile strength). In recent years, as environmental pollution problems caused by automobile exhaust gases have been in the spotlight, an increasing amount of studies have focused on reducing the weight of automobiles using advanced high strength steels having a tensile strength of 780 MPa or higher in order to further reduce fuel consumption.

High-strength steels for automobiles typically include transformation-induced plasticity (TRIP) steels and dual phase (DP) steels.

Processes for manufacturing advanced high-strength steels having excellent workability are largely divided into a slab manufacturing process, a hot rolling process, a process for cooling and coiling a hot-rolled strip, a cold rolling process and an annealing process. A steel obtained by cold-rolling a hot-rolled strip having ferrite and pearlite structures, annealing the cold-rolled strip at a temperature between the A₁transformation point and the A₃transformation point and then transforming austenite, formed in the annealing process, to martensite by controlling a cooling rate of the steel during cooling, is referred to as a dual phase steel. The strength of the dual phase steel is determined by the fractions of martensite and ferrite contained therein. As the ratio of martensite in the total structure increases, strength increases and ductility decreases. For this reason, the dual phase steel should have a suitable martensite ratio. Meanwhile, there is a method in which austenite is formed in the annealing process, as in the above-described method for manufacturing the dual phase steel, and then some of the austenite is retained at room temperature by controlling the cooling rate of the steel and the end cooling temperature during the cooling process, thereby increasing both the strength and ductility of the steel. The retained austenite thus produced is transformed into martensite during plastic deformation to increase the strength of the steel, while increasing the ductility of the steel by relieving stress concentration by plasticity-induced transformation. This steel is referred to as transformation-induced plasticity (TRIP) steel and is used as a high-strength steel having high strength and ductility.

In order to increase the strength and ductility of the transformation-induced plasticity steel, it is very important to suppress the transformation of the retained metastable austenite to martensite at room temperature so as to maintain the retained austenite at a given fraction or higher. For this purpose, Si, Al and the like are added to the steel to increase the activity of carbon in ferrite so as to suppress the production of carbides, so that the concentration of carbon in the austenite phase is increased, thereby ensuring the stability of the retained austenite.

Known technologies for ensuring such high strength and elongation include technology disclosed in Japanese Patent Laid-Open Publication No. Hei 6-1458920. According to this technology, a steel having a strength of 490 MPa or higher and an elongation of 35% or higher is provided by adding 0.03-0.3 wt % of Mo to a steel containing 0.6-0.22 wt % of C, 0.05-1.0 wt % of Si, 0.5-2.0 wt % of Mn and 0.25-1.5 wt % of Al.

Also, Korean Patent Laid-Open Publication No. 2002-0045212 discloses a steel containing 0.15-0.30 wt % C, 1.5-2.5 wt % Si, 0.5-2.0 wt % Mn and 0.02-0.1 wt % Al and having a strength of 780 MPa or higher and an elongation of 30% or higher.

Japanese Patent Laid-Open Publication No. 2005-336526 discloses a steel containing 0.06-0.6 wt % C, 0.5-3.0 wt % Si+Al and 0.5-3.0 wt % Mn and having a strength of about 800 MPa and an elongation of 40%. However, when examining the actual content of the components described in this Japanese patent publication, it can be seen that the amount of Al+Si is 1.5 wt % or less. For example, when Al is added in an amount of 1 wt %, Si is added in an amount of 0.5 wt %. Therefore, accurately speaking, the content of Si+Al is not 0.5-3.0 wt %, but 1.5 wt % or less.

Moreover, the manufacturing processes disclosed in the Japanese patent publication are divided into two: a process of performing annealing or plating after hot rolling; and a process of performing two-step annealing after hot rolling and cold rolling. In the two-step annealing, martensite is formed in the matrix structure by the first-step annealing after cold rolling, and the second-step annealing is performed during the annealing process which is performed on a conventional transformation-induced plasticity steel. This is disadvantageous in economic terms and has low applicability.

The above-described prior technologies have strengths of only about 490 MPa and 780 MPa, respectively, and employ complex manufacture processes. Thus, in order to ensure both a tensile strength of at least 780 MPa and an elongation of at least 28%, technology that is advantageous in economic terms and has high applicability is required.

DISCLOSURE Technical Problem

An aspect of the present invention provides a high-strength steel sheet and a hot-dipped galvanized steel sheet, which have a tensile strength of 780-980 MPa and an elongation of 28% or higher and, at the same time, are free of edge cracks, as a result of controlling the composition thereof and adjusting a cooling step during hot rolling so as to have a martensite fraction of 30-70%, as well as methods for manufacturing these steel sheets.

Technical Solution

According to an aspect of the present invention, there are provided a high-strength steel sheet and a hot-dipped galvanized steel sheet, which include, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and the balance of Fe and inevitable impurities, and satisfy 1.75%≦Si+Al≦3.25%.

According to another aspect of the present invention, there is provided a method for manufacturing a high-strength hot-rolled steel sheet, the method including: hot-rolling a steel slab, which includes, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and the balance of Fe and inevitable impurities, and satisfies 1.75≦Si+Al≦3.25%, at a temperature higher than the A₃transformation point, and primarily cooling the hot-rolled steel sheet to the temperature ranging from 600° C. to 800° C. at a cooling rate of 30˜200° C./sec; air-cooling the primarily cooled hot-rolled steel sheet at a temperature ranging from 600° C. to 800° C.; secondarily cooling the air-cooled hot-rolled steel sheet to a temperature ranging from room temperature to 300° C. at a cooling rate of 50˜200° C./sec; and coiling the secondarily cooled hot-rolled steel sheet at a temperature ranging from room temperature to 300° C.

According to another aspect of the present invention, there is provided a method for manufacturing a high-strength cold-rolled steel sheet, the method including: hot-rolling a steel slab, which includes, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and the balance of Fe and inevitable impurities, and satisfies 1.75≦Si+Al≦3.25%, at a temperature higher than the A₃transformation point, and primarily cooling the hot-rolled steel to a temperature ranging from 600° C. to 800° C. at a cooling rate of 30˜200° C./sec; air-cooling the primarily cooled hot-rolled steel sheet at a temperature ranging from 600° C. to 800° C.; secondarily cooling the air-cooled hot-rolled steel sheet at a temperature ranging from room temperature to 300° C. at a cooling rate of 50˜200° C./sec; coiling the secondarily cooled hot-rolled steel sheet at a temperature ranging from room temperature to 300° C.; and cold-rolling the coiled hot-rolled steel sheet at a reduction ratio of 30-50% and annealing the cold-rolled steel sheet.

According to another aspect of the present invention, there is provided a method for manufacturing a high-strength, hot-dipped galvanized steel sheet, the method including, in addition to the method for manufacturing the high-strength cold-rolled steel sheet, hot-dip galvanizing the cold-rolled and annealed steel sheet.

DESCRIPTION OF DRAWINGS

The above and other aspects, features and other advantages of the present invention will be more clearly understood from the following detailed description taken in conjunction with the accompanying drawings, in which:

FIG. 1 shows the results of testing tensile strength according to the [Si]+[Al] content vs. the martensite volume fraction after hot rolling in the present invention in a preliminarily test of the present invention;

FIG. 2 shows the results of testing elongation according to the [Si]+[Al] content vs. the martensite volume fraction after hot rolling in a preliminarily test of the present invention;

FIG. 3 shows the edge crack length according to the [Si]+[Al] content vs. the martensite volume fraction after hot rolling, measured in a preliminarily test of the present invention; and

FIG. 4 shows the tensile strength, elongation and edge crack length after the cold rolling of inventive steels manufactured according to the present invention and comparative steels according to the [Si]+[Al] content vs. the martensite volume fraction after hot rolling.

BEST MODE

Exemplary embodiments of the present invention will now be described in detail with reference to the accompanying drawings. The invention may, however, be embodied in many different forms and should not be construed as being limited to the embodiments set forth herein. Rather, these embodiments are provided so that this disclosure will be thorough and complete, and will fully convey the scope of the invention to those skilled in the art. In the drawings, the thicknesses of layers and regions are exaggerated for clarity. Like reference numerals in the drawings denote like elements, and thus their description will be omitted.

Hereinafter, the composition range of the present invention will be described in detail (hereinafter, wt %).

The content of carbon (C) is 0.1-0.25%. C is the most important component and has a close connection with all physical and chemical properties, including strength and ductility. In the steel sheet of the present invention, if the content of carbon is less than 0.1%, the fraction and stability of retained austenite will be decreased, and if it is more than 0.25%, the weldability of the steel sheet will be reduced and the workability thereof will be deteriorated due to an excessive increase in the fraction of the second phase. For these reasons, the content range of C is limited to 0.1-0.25%.

The content of silicon (Si) is 1.0-1.9%. Si is dissolved in ferrite to form a solid solution to thereby stabilize ferrite. In the cold-rolled steel sheet of the present invention, Si is dissolved in ferrite to form a solid solution, so that it serves to enhance the activity of carbon so as to increase the concentration of carbon in the austenite phase and suppress the production of carbides, thereby increasing the stability of retained austenite. Also, it is dissolved to form a solid solution, thereby increasing the strength of the steel sheet. If the content of Si is less than 1.0%, the strength of the steel will be decreased and the effect of suppressing the production of carbides such as a carbide phase will be reduced, and if it is more than 1.9%, it will cause hot-rolling scales and deteriorate the platability and weldability of the steel sheet. For these reasons, the content of Si is limited to 1.0-1.9%. The content of manganese (Mn) is 1.5-2.5%. Mn serves to increase the hardenability of the steel so as to facilitate the production of low-temperature transformation phases such as ferrite and bainite and increase the strength of the steel. It also stabilizes austenite. Thus, if it is added in an amount of less than 1.5%, the above effects cannot be expected, and if it is added an amount of more than 2.5%, it will deteriorate the weldability of the steel sheet, form a segregation region in the center of the strip during hot rolling and form inclusions to cause hydrogen embrittlement. For these reasons, the content of Mn is limited to 1.5-2.5%.

The content of aluminum (Al) is limited to 0.5-1.6%. Although Al has a solid solution-strengthening effect inferior to that of Si, it is a ferrite-stabilizing element that shows a solid solution-strengthening effect, suppresses the production of carbides and increases the carbon concentration of retained austenite to thereby increase the stability of the retained austenite. If the content of Al is less than 0.5%, the stability of austenite will be reduced and it will difficult to suppress the production of carbides, and if it is more than 1.6%, the fraction of austenite will be reduced, so that the ductility of the steel sheet will be reduced and the surface characteristics thereof will be deteriorated. For these reasons, the content of Al is limited to 0.5-1.6%. The content of titanium (Ti) is 0.005-0.03%. Ti serves to suppress the formation of AlN nitrides resulting from the bonding of Al with N to form TiN, so as to allow the Al content of the steel to perform its functions. If it is added in an amount of less than 0.005%, it will hardly perform such functions, and if it is added in an amount of more than 0.03%, an increase in effects resulting from the increase in the added amount thereof cannot be expected. For these reasons, the content of Ti is limited to 0.005-0.03%.

The content of boron (B) is 5-30 ppm. B is a component that improves the hardenability of the steel even when it is added in a small amount. When it is added in an amount of more than 5 ppm, it is segregated in the austenite grain boundary at high temperatures to suppress the formation of ferrite, thereby contributing to an increase in the hardenability of the steel, but if it is added in an amount of more than 30 ppm, it will increase the recrystallization temperature of the steel sheet, thereby reducing the drawability of the steel sheet and deteriorating the weldability thereof. For these reasons, the content of B is limited to 5-30 ppm.

The content of antimony (Sb) is 0.01-0.03%. If Sb is added in a suitable amount of 0.01-0.03%, it will improve the surface characteristics of the steel sheet, but if it is added in an amount of more than 0.03%, it will cause Sb enrichment on the surface of the steel sheet to deteriorate the surface characteristics thereof. Sb is an impurity that may be inevitably contained in steel materials, and an Sb content of less than 0.01% can be found in steels manufactured without a special object. Also, an Sb content of less than 0.01% is too small to cause Sb enrichment on the surface to change the surface characteristics. Thus, the content of Sb is limited to 0.01-0.03%.

In the present invention, 1.75≦Si+Al≦3.25% is satisfied. Si and Al all serve to suppress the formation of carbides in the steel to increase the content of solid solution carbon in retained austenite and improve the stability of retained austenite. Thus, when considering the strength and elongation of the TRIP steel, it is preferable to control the content of Si together with the content of Al. If the content of the two components is more than 3.25%, the surface qualities (e.g., platability) of the steel can be deteriorated due to the excessive formation of surface oxides, and the strength and ductility of the steel can be reduced because the fraction of austenite will be decreased during two-phase region annealing. Also, if the content of the two components is less than 1.75%, the solid solution-strengthening effect required for manufacturing a TRIP steel having a desired tensile strength of 780 MPa or higher will be reduced and it will be difficult to ensure the stability of retained austenite. For these reasons, the content of Si+Al is limited to 1.75-3.25%.

The steel sheet of the present invention comprises the above-described composition and the balance of Fe and inevitable impurities.

Hereinafter, the manufacturing method of the present invention will be described in detail.

In the present invention, a steel slab satisfying the above-described composition is hot-rolled at a temperature higher than the A₃transformation point, and then primarily cooled at a cooling rate of 30-200° C./sec. If the cooling rate is less than 30° C./sec, a pearlite structure can be formed to make it difficult to realize the desired material, and if it is more than 200° C./sec, the material can be distorted due to the occurrence of residual stress caused by the occurrence of temperature deviation of the steel sheet. For these reasons, the cooling rate is limited to 30˜200° C./sec.

After the primary cooling, the steel sheet is maintained in the temperature range of 600˜800° C. This means that the steel sheet is cooled by natural convection in the air at room temperature without forced cooling. At a temperature lower than 600° C., it is difficult to ensure a ferrite phase-forming fraction, and at higher than 800° C., excessive ferrite can be formed or a pearlite structure can be formed.

After the air cooling, the steel sheet is secondarily cooled at a cooling rate of 50˜200° C./sec, and then coiled at a temperature ranging from room temperature to 300° C. At a cooling rate of less than 50° C./sec, a bainite phase can be formed to make it difficult to realize the desired structure, and at a cooling rate of more than 200° C./sec, excessive martensite can be formed and the shape of the hot-rolled sheet can be distorted. For these reasons, the secondary cooling rate is limited to 50˜200° C./sec.

Also, if the coiling temperature is higher than 300° C., a bainite phase will be formed to make it difficult to realize the desired structure, and if the coiling is performed at a temperature ranging from room temperature to 300° C., the matrix structure will have a lath-shaped fine martensite structure, and thus the steel sheet will have a high dislocation density and a uniform solid solution carbon distribution after hot rolling. For these reasons, the coiling temperature is limited to a temperature ranging from room temperature to 300° C.

After the coiling, the steel sheet is cold-rolled at a reduction ratio of 30-500. In conventional methods, the coiled steel sheet is cold-rolled at a reduction rate of 60%, whereas, in the present invention, the steel sheet is cold-rolled at a reduction rate of 30-500. By doing so, dislocation density in a ferrite structure among the martensite and ferrite structures formed after the coiling can be sufficiently increased and the re-dissolution of carbon and the formation of an austenite phase will uniformly occur. If the cold-rolling reduction ratio is less than 30%, dislocation density in the ferrite structure will be insufficient to make it difficult to obtain the desired structure, and if it is more than 50%, fine cracks at the boundary between the martensite phase and the ferrite phase will easily occur so that crack defects will occur, particularly at the edge of the rolled sheet. For these reasons, the cold-rolling reduction ratio is limited to 30-50%.

After the cold rolling, a conventional annealing process is performed.

In the embodiment of the present invention, after the cold rolling, the steel sheet is hot-dipped galvanized or galvannealed. After the annealing process, the steel sheet is passed through a plating bath containing molten galvanizing material, whereby a plating layer is attached to the surface layer to a given thickness. Herein, the temperature of the plating bath is preferably 450 to 500° C., and the steel sheet is cooled slowly at a rate of 30° C./sec or lower, thereby manufacturing a hot-dipped galvanized steel sheet.

Also, immediately after the steel sheet has been passed through the plating bath, it is galvannealed by being heated to a temperature ranging from 500° C. to 600° C., and then cooled slowly at a rate of 30° C./sec, thereby manufacturing a galvannealed steel sheet.

Hereinafter, the structure of the steel sheet that is manufactured according to the method of the present invention will be described in detail.

In the present invention, the hot-rolled steel sheet after the hot rolling is characterized in that the fraction of martensite is 30-70%. The lath-shaped fine martensite structure resulting from the primary and secondary cooling processes after the hot rolling induces uniform martensite transformation in the annealing process following the cold rolling, thereby increasing the fraction of stabilized retained austenite. In conventional methods, a hot-rolled steel sheet is manufactured by a one-step cooling process, and in this case, the hot-rolled steel sheet contains a pearlite structure in which coarse carbides are present. The coarse carbides are re-dissolved to a form a solid solution in the annealing process following cold rolling; however, because the carbides are coarse in size, they will be likely to remain without being re-dissolved even at a high temperature of 700° C. or higher, and austenite during annealing will start to be formed mainly in the vicinity of carbides having a high carbon concentration and will not be easily formed in a region in which the amount of carbides is relatively small. Thus, during two-phase region annealing, the fraction of austenite will be reduced and a local deviation in the austenite fraction will also occur, so that the fraction of retained austenite that occurs in the cooling process following annealing will be reduced. The martensite phase in a hot-rolled steel sheet can solve this shortcoming.

In the present invention, when a martensite phase is formed through the primary and secondary cooling processes, a hot-rolled steel sheet can be formed without substantially forming carbides. In the hot-rolled steel sheet containing this martensite phase, fine carbides are formed in the martensite structure having a high dislocation density during annealing, and then immediately re-dissolved to form solid solution. Thus, the deviation of carbon concentration in this steel sheet will be significantly reduced as compared to a steel sheet having a pearlite structure. Accordingly, the austenite phase that is formed during annealing is evenly distributed around the grain boundary and around a region that was a lath martensite structure, and thus the fraction and stability of retained austenite can be increased.

Moreover, by applying a relatively low cold-rolling reduction ratio by virtue of a dislocation density that was increased by the martensite phase, the formation of fine cracks that easily occur at the edge of the steel sheet after cold rolling can be suppressed. Herein, if the fraction of martensite in the hot-rolled steel sheet is less than 30%, the effect of increasing the fraction of retained austenite will be insignificant, and it is more than 70%, fine cracks will occur at the edge during cold rolling. For these reasons, in the present invention, the fraction of the martensite structure in the hot-rolled steel sheet is limited to 30-70%.

In the cold-rolled steel sheet that has undergone the cold rolling and the annealing process in the present invention, retained austenite is produced in the ferrite and bainite structures. Herein, the fraction of the retained austenite in the present invention is 5-15%. Due to the influence of the martensite structure having a lath shape in the hot-rolled steel sheet, the retained austenite phase shows a lath shape and is more stable than retained austenite of other shapes. In addition, the fraction of the bainite is 20-40%, and the remainder consists of ferrite.

According to the method of the present invention, a high-strength, cold-rolled steel sheet having excellent workability, a tensile strength of 780-980 MPa and an elongation of 28% or higher can be manufactured.

Hereinafter, the present invention will be described in detail with reference to examples.

Mode for Invention EMBODIMENTS AND COMPARATIVE EXAMPLES EXAMPLES

Table 1 below shows the composition ranges of steels used in Examples. In the present invention, it is important to suitably adjust the components of the steels in order to obtain tensile strength and elongation which are sought in the present invention.

TABLE 1 No. C Mn Si P S Al Ti B Sb N Si + Al Remarks 1 0.11 1.95 1.51 ≦0.02 ≦0.005 1.05 0.018 0.001 0.022 ≦0.003 2.56 Inventive steel 2 0.15 1.98 1.45 ≦0.02 ≦0.005 1.23 0.016 0.001 0.017 ≦0.003 2.68 Inventive steel 3 0.19 1.52 1.47 ≦0.02 ≦0.005 0.55 0.014 0.001 0.02 ≦0.003 2.02 Inventive steel 4 0.21 1.55 1.49 ≦0.02 ≦0.005 1.03 0.011 0.001 0.015 ≦0.003 2.52 Inventive steel 5 0.19 1.51 1.5 ≦0.02 ≦0.005 1.45 0.008 0.001 0.027 ≦0.003 2.95 Inventive steel 6 0.25 1.89 1.44 ≦0.02 ≦0.005 0.82 0.019 0.001 0.02 ≦0.003 2.26 Inventive steel 7 0.24 2.44 1.56 ≦0.02 ≦0.005 1.25 0.015 0.001 0.022 ≦0.003 2.81 Inventive steel 8 0.23 1.62 1.65 ≦0.02 ≦0.005 1.39 0.014 0.001 0.015 ≦0.003 3.04 Inventive steel 9 0.08 1.88 0.05 ≦0.02 ≦0.005 0.03 0.017 0.001 0.02 ≦0.003 0.08 Comparative steel 10 0.21 1.56 1.83 ≦0.02 ≦0.005 0.03 0.012 0.001 0.022 ≦0.003 1.86 Comparative steel 11 0.18 1.78 0.06 ≦0.02 ≦0.005 0.56 0.015 0.001 0.019 ≦0.003 0.62 Comparative steel 12 0.21 1.82 0.52 ≦0.02 ≦0.005 0.62 0.014 0.001 0.02 ≦0.003 1.14 Comparative steel 13 0.19 1.65 1.07 ≦0.02 ≦0.005 0.58 0.01 0.001 0.021 ≦0.003 1.65 Comparative steel 14 0.18 1.48 1.52 ≦0.02 ≦0.005 0.07 0.015 0.001 0.022 ≦0.003 1.59 Comparative steel 15 0.23 2.35 1.63 ≦0.02 ≦0.005 1.78 0.015 0.001 0.022 ≦0.003 3.41 Comparative steel

The steel slabs having the compositions of Table 1 were hot-rolled, cooled, coiled, cold-rolled and then annealed, and the cooling conditions and cold-rolling reduction ratios thereof and the results of tensile tests therefor are shown in Table 2 below.

The tensile properties among the mechanical properties of materials vary depending on the components of the materials and the manufacturing conditions of fine structures. In the steel sheets of the present invention, a fine martensite structure of lath structure produced immediately after hot rolling induces uniform austenite transformation in an annealing process to increase the fraction of stabilized retained austenite, whereby the excellent mechanical properties of the steel sheet are obtained through transformation when it undergoes deformation.

Also, by applying a relatively low cold-rolling reduction ratio by virtue of a dislocation density that was increased by the martensite phase, the formation of fine cracks that easily occur at the edge of the steel sheet after cold rolling is inhibited. In order to obtain such characteristics, the components proposed in the present invention should be satisfied, and the cooling process immediately after hot-rolling, the fraction of martensite and the cold-rolling reduction ratio should be optimally adjusted.

TABLE 2 Edge Primary Intermediate Air-cooling Secondary Cold-rolling crack cooling rate temperature temperature cooling rate Coiling reduction Martensite TS T-El length (° C./s) (° C.) (sec) (° C./s) temperature (° C.) ratio (%) fraction (%) MPa (%) (mm) 1 Inventive 100 680 5 150 200 50 70 740 30 0.02 material 2-1 Inventive 100 700 5 150 200 45 70 783 32 0.05 material 2-2 Comparative 100 600 1 200 200 55 70 785 31 12 material 2-3 Comparative 120 650 5 250 200 60 70 778 32 32 material 3-1 Inventive 60 680 8 120 200 40 55 787 35 0 material 3-2 Comparative 80 680 8 120 220 60 55 790 33 25 material 3-3 Comparative 100 700 2 250 200 40 90 792 31 34 material 4-1 Inventive 80 650 8 100 200 30 60 925 32 0 material 4-2 Comparative 80 650 1 200 200 30 75 930 30 22 material 4-3 Comparative 80 650 2 250 200 30 90 840 30 4 material 5-1 Inventive 100 680 5 150 200 40 65 840 31 0.02 material 5-2 Comparative 30 700 5 80 200 40 10 773 27 0 material 6 Inventive 80 650 8 120 200 40 50 945 30 0.05 material 7-1 Inventive 80 650 8 120 200 50 45 990 28 0.07 material 7-2 Comparative 50 650 9 120 200 60 45 988 28 3 material 7-3 Comparative 30 — — — 650 50 0 985 26 0 material 8-1 Inventive 80 680 8 100 200 40 60 935 30 0.08 material 8-2 Comparative 30 — — — 650 40 0 920 27 0 material 9 Comparative 80 700 8 100 100 40 60 680 20 0.02 material 10-1 Comparative 30 650 6 100 200 40 40 930 27 0 material 10-2 Comparative 30 — — — 650 40 0 825 28 0 material 11 Comparative 80 980 8 100 200 40 50 790 18 0 material 12-1 Comparative 100 700 8 100 200 40 50 800 22 0.05 material 12-2 Comparative 120 700 10 100 200 60 50 815 22 02 material 12-3 Comparative 50 — — — 630 40 0 790 20 0 material 13 Comparative 80 720 15 100 250 40 40 785 27 0 material 14 Comparative 80 710 8 100 200 40 60 740 35 0 material 15 Comparative 50 680 8 100 200 40 40 960 23 0 material

Table 2 above shows the cooling process immediately after hot rolling, the martensite fraction, the cold-rolling reduction ratio and the results of measuring the mechanical properties and the degree of cracking at the edge.

As can be seen in Table 2, in comparative materials 9 to 15, the fundamental component requirements were not satisfied, and thus the strength and ductility sought in the present invention were not obtained even when the cooling condition, the martensite fraction and the cold-rolling reduction rate were suitably controlled during the manufactured.

Also, in comparative materials 2-2, 2-3, 3-2, 3-3, 4-2, 4-3, 5-2, 7-2, 7-3 and 8-2, the component requirements were satisfied, but requirements such as the martensite fraction and the cold-rolling reduction rate were not satisfied, and thus the desired mechanical properties were not obtained or fine cracks occurred at the edge.

The edge crack length in Table 2 was measured by observing the martensite fraction with a 200× optical microscope.

Specifically, the measurement results were obtained by etching a fine structure at the t/4 thickness position with a 2% Nital etching solution and observing the etched portion with an image analyzer. Also, to measure the length of fine cracks that occurred at the edge portions, the edge portions of the cold-rolled sheets were randomly selected, and 30 or more cracks having the longest length within a length of 100 mm were selected from the selected edge portions, and then the lengths thereof were averaged.

The steel sheets according to the present invention have a tensile strength of 780-980 MPa and an elongation of 28% or higher as a result of controlling the composition and manufacture conditions thereof. Thus, the steel sheets of the present invention can be used for structural parts requiring high strength and workability. In addition, the present invention can prevent cracks from occurring at the edge of the steel sheets, thereby increasing economic efficiency.

While the present invention has been shown and described in connection with the exemplary embodiments, it will be apparent to those skilled in the art that modifications and variations can be made without departing from the spirit and scope of the invention as defined by the appended claims.

Claims

1. A high-strength steel sheet comprising, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfying 1.75%≦Si+Al≦3.25%.

2. A high-strength, hot-rolled steel sheet comprising, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfying 1.75%≦Si+Al≦3.25%, wherein the microstructure of the steel sheet after hot rolling comprises 30-70% martensite and a balance of ferrite.

3. A high-strength, cold-rolled steel sheet comprising by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfying 1.75≦Si+Al≦3.25%, wherein the microstructure of the steel sheet after cold rolling comprises 5-15% retained austenite, 20-40% bainite and a balance of ferrite.

4. The high-strength steel sheet of claim 3, wherein the cold-rolled steel sheet has a tensile strength of 780-980 MPa and an elongation of 28% or higher.

5. A high-strength, hot-dipped galvanized steel sheet comprising, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfying 1.75≦Si+Al≦3.25%, the steel sheet having a galvanized layer.

6. A method for manufacturing a high-strength, hot-rolled steel sheet, the method comprising:

hot-rolling a steel slab, which comprises, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfies 1.75%≦Si+Al≦3.25%, at a temperature higher than an A3 transformation point, and primarily cooling the hot-rolled steel sheet to a temperature ranging from 600° C. to 800° C. at a cooling rate of 30-200° C./sec;

air-cooling the primarily cooled hot-rolled steel sheet at a temperature ranging from 600° C. to 800° C.;

secondarily cooling the air-dried hot-rolled steel sheet to a temperature ranging from room temperature to 300° C. at a cooling rate of 50-200° C./sec; and

coiling the secondarily cooled hot-rolled steel sheet at a temperature ranging from room temperature to 300° C.

7. A method for manufacturing a high-strength, cold-rolled steel sheet, the method comprising:

hot-rolling a steel slab, which comprises, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfies 1.75≦Si+Al≦3.25%, at a temperature higher than an A3 transformation point or higher, and primarily cooling the hot-rolled steel sheet to a temperature ranging from 600° C. to 800° C. at a cooling rate of 30˜200° C./sec;

air-cooling the primarily cooled hot-rolled steel sheet at a temperature ranging from 600° C. to 800° C.;

secondarily cooling the air-dried hot-rolled steel sheet to a temperature ranging from room temperature to 300° C. at a cooling rate of 50-200° C./sec;

coiling the secondarily cooled hot-rolled steel sheet at a temperature ranging from room temperature to 300° C.; and

cold-rolling the coiled hot-rolled steel sheet at a reduction ratio of 30-50% and annealing the cold-rolled steel sheet.

8. A method for manufacturing a high-strength, hot-dipped, galvanized steel sheet, the method comprising:

hot-rolling a steel slab, which comprises, by wt %, 0.1-0.25% C, 1.0-1.9% Si, 1.5-2.5% Mn, 0.5-1.6% Al, 0.005-0.03% Ti, 5-30 ppm B, 0.01-0.03% Sb and a balance of Fe and inevitable impurities, and satisfies 1.75%≦Si+Al≦3.25%, at a temperature of an A3 transformation point or higher, and primarily cooling the hot-rolled steel sheet to a temperature ranging from 600° C. to 800° C. at a cooling rate of 30˜200° C./sec;

air-cooling the primarily cooled hot-rolled steel sheet at a temperature ranging from 600° C. to 800° C.;

secondarily cooling the air-dried hot-rolled steel sheet to a temperature ranging from room temperature to 300° C. at a cooling rate of 50-200° C./sec;

coiling the secondarily cooled hot-rolled steel sheet at a temperature ranging from room temperature to 300° C.;

cold-rolling the coiled hot-rolled steel sheet at a reduction ratio of 30-50% and annealing the cold-rolled steel sheet; and

galvanizing the annealed cold-rolled steel sheet.