JP5456026B2 - High-strength steel sheet, hot-dip galvanized steel sheet with excellent ductility and no cracks at the edge, and manufacturing method thereof - Google Patents

Description

The present invention relates to a steel sheet for automobiles excellent in strength and workability mainly used in various structural members of automobiles, and more specifically, a tensile strength of 780 to 980 MPa and an elongation of at least 28% or more. The present invention relates to a method for producing a high-strength cold-rolled steel sheet and a hot-dip galvanized steel sheet that does not crack at the edge (Edge) portion of the rolled sheet during cold rolling.

  Steel plates used as structural members in automobile components play a role in improving passenger safety in the event of a vehicle collision by well absorbing external impacts. Recently, as such parts, steel sheets having excellent strength and formability such as tensile strength of 780 MPa or more and elongation rate of 25% or more are mainly used. Such a steel plate needs to have high tensile strength, high elongation, and low yield ratio (yield strength / tensile strength ratio).

  Furthermore, in recent years, environmental pollution problems due to automobile exhaust gas have attracted attention, and in order to further reduce fuel consumption, research for reducing the weight of automobiles using ultra-high strength steel with a tensile strength of 780 MPa or more has been conducted. It is increasing.

  High-strength steel for automobiles generally includes transformation induced plasticity (TRIP) steel and dual phase (DP) steel.

The manufacturing process of ultra-high-strength steel excellent in workability is roughly divided into a slab manufacturing process, a hot rolling process, a process of cooling and winding a hot-rolled plate material, a cold rolling process, and an annealing process. A plate having a hot rolled ferrite and pearlite processed by cold rolling, after the cold rolled sheet was annealed at a temperature between A ₁ transformation point to A ₃ transformation point, the cooling rate of the steel plate during cooling The steel obtained by transforming the austenite formed in the annealing process to martensite by adjusting the temperature is called a duplex steel. The strength of the duplex steel is determined by the fraction of martensite and ferrite contained in the duplex steel. As the proportion of martensite in the overall structure increases, the strength increases and the ductility decreases. For this reason, the duplex stainless steel must have the proper martensite ratio. On the other hand, by forming austenite in the annealing process as in the above-described method for producing the duplex steel, by controlling the cooling rate and the end cooling temperature in the cooling process, a part of the austenite remains at room temperature, thereby making the steel material There is a method of simultaneously increasing the strength and ductility of the steel. Residual austenite generated in this manner is transformed into martensite during plastic deformation in order to increase the strength of the steel material, and at the same time, ductility is increased by relaxing stress concentration by plastic-induced transformation. This is a transformation induced plasticity (TRIP) steel used as a high strength steel having high strength and high ductility.

  In order to increase the strength and ductility of the above transformation-induced plastic steel, it is very important to suppress metastable retained austenite to martensite at room temperature and maintain retained austenite above a certain fraction It is. Therefore, the stability of retained austenite is increased by adding Si, Al, etc. to increase the activity of carbon in the ferrite to suppress the formation of carbides and increase the concentration of carbon in the austenite phase. Secure.

  Known techniques for ensuring such high strength and elongation include the technique disclosed in Patent Document 1. According to this technique, C: 0.6 to 0.22 wt (wt)%, Si: 0.05 to 1.0 wt%, Mn: 0.5 to 2.0 wt%, and Al: 0.25 to 1 By adding Mo: 0.03 to 0.3 wt% to steel containing 0.5 wt%, a steel having a strength of 490 MPa or more and an elongation of 35% or more is provided.

  In Patent Document 2, C: 0.15 to 0.30 wt%, Si: 1.5 to 2.5 wt%, Mn: 0.5 to 2.0 wt%, and Al: 0.02 to 0.1 wt% And steel having a strength of 780 MPa or more and an elongation of 30% or more are disclosed.

  In Patent Document 3, C: 0.06 to 0.6 wt%, [Si + Al]: 0.5 to 3.0 wt%, Mn: 0.5 to 3.0 wt%, and a strength of about 800 MPa and 40% Steels having an elongation of about 5 are disclosed. However, when the actual content of the component currently disclosed by patent document 3 is seen, it turns out that the quantity of [Al + Si] is 1.5 wt% or less. For example, when 1 wt% of Al is added, 0.5 wt% of Si is added. Therefore, the component range of [Si + Al] is not 0.5 to 3.0 wt%, but 1.5 wt% is an accurate expression.

  Moreover, although the manufacturing process currently disclosed by patent document 3 is divided into two, one is a manufacturing process which performs annealing and plating after hot rolling, and the other is hot rolling and cold rolling. Thereafter, it is a manufacturing process for performing two-stage annealing. In the two-stage annealing, martensite is formed in the matrix structure by the first-stage annealing after the cold rolling, and the second-stage annealing is performed by an ordinary transformation-induced plastic steel. This results in economic losses and is less applicable.

  The above-described conventional techniques each have a strength of only about 490 MPa to 780 MPa, and the manufacturing process is complicated. Therefore, in order to simultaneously secure a tensile strength of 780 MPa or more and an elongation of 28% or more, a manufacturing technique that is economically advantageous and highly applicable is required.

Japanese Patent Publication No. 6-1458920 Korean Patent Publication No. 2002-0045212 Japanese Patent Publication No. 2005-336526

  The object of the present invention is to control the composition components and adjust the cooling stage in hot rolling so as to include a martensite fraction of 30 to 70%, thereby providing a tensile strength of 780 to 980 MPa and 28% or more. It is to provide a high-strength steel plate, a hot-dip galvanized steel plate, and a method for producing the same, in which an edge portion has no cracks.

In one embodiment of the present invention, C: 0.1 to 0.25%, Si: 1.0 to 1.9%, Mn: 1.5 to 2.5%, Al: 0.5 to 1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and inevitable impurities, and 1.75 ≦ Si + Al A high-strength hot-rolled steel sheet satisfying ≦ 3.25%, wherein the microstructure of the steel sheet after hot rolling is composed of 30-70% martensite and the balance ferrite. And a galvanized steel sheet.

In another embodiment of the present invention, by weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5 ~1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and unavoidable impurities, 1.75 ≦ Si + Al the steel slab satisfying ≦ 3.25%, comprising the steps of hot rolling at _{a 3} or higher, the step of primary cooling in the temperature range of 600 to 800 ° C. at a cooling rate of 30 to 200 ° C. / s (sec) And air cooling the primary-cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C, and the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C / s at a temperature range of room temperature to 300 ° C. And a step of winding the secondary-cooled hot-rolled steel sheet in a temperature range of room temperature to 300 ° C. A method for producing a high-strength hot-rolled steel sheet is provided.

In another embodiment of the present invention, by weight, C: 0.1 to 0.25%, Si: 1.0 to 1.9%, Mn: 1.5 to 2.5%, Al: 0 .5~1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and unavoidable impurities, and 1. A steel slab satisfying 75 ≦ Si + Al ≦ 3.25% is subjected to primary rolling in a temperature range of 600 to 800 ° C. at a stage of hot rolling in a temperature range of A ₃ or higher and a cooling rate of 30 to 200 ° C./s. Air cooling the primary-cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C., and air-cooling the hot-rolled steel sheet in a temperature range of room temperature to 300 ° C. at a cooling rate of 50 to 200 ° C./s. A stage of secondary cooling, a stage of winding the secondary-cooled hot-rolled steel sheet in a temperature range of room temperature to 300 ° C .; Manufacturing a high-strength cold-rolled steel sheet, comprising: cold-rolling the rolled hot-rolled steel sheet at a rolling reduction of 30 to 50%; and annealing the cold-rolled steel sheet. A method is provided.

  In another aspect of the present invention, there is provided a method for producing a high-strength hot-dip galvanized steel sheet, further comprising the step of hot-dip galvanizing in addition to the above-described method for producing a high-strength cold-rolled steel sheet.

  The above as well as other aspects, features, and other advantages of the present invention will become apparent from the following detailed description of the invention when taken in conjunction with the accompanying drawings.

The test result of tensile strength by the [Si] + [Al] content in the preliminary test of the present invention versus the martensite volume fraction after hot rolling is shown. The test result of the elongation rate by the [Mar] site volume fraction after hot-rolling [Si] + [Al] content in the preliminary test of this invention is shown. The crack length of the edge part after the cold rolling by the [Si] + [Al] content in the preliminary test of the present invention versus the martensite volume fraction after hot rolling is shown. [Si] + [Al] content of inventive steel and comparative steel produced according to the present invention vs. tensile strength, elongation by hot-rolled martensite volume fraction, and crack length of edge part after cold rolling .

  Embodiments of the present invention will be described in detail with reference to the drawings. However, the present invention can be embodied in various forms and should not be construed as 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 thickness of layers and regions are exaggerated for clarity. Since the reference numerals in the drawings represent the components, these descriptions are omitted.

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

  The content of carbon (C) is 0.1 to 0.25%. C is the most important component and has a close relationship with all physical and chemical properties such as strength and ductility. In the steel sheet of the present invention, if the amount of carbon is less than 0.1%, the fraction and stability of retained austenite decrease, and if it exceeds 0.25%, the weldability decreases and the second phase fraction is excessive. However, there is a disadvantage that the workability is reduced due to a large increase. Therefore, the composition range of C is limited to 0.1 to 0.25%.

  The content of silicon (Si) is 1.0 to 1.9%. Si is a component that is dissolved in ferrite and stabilizes the ferrite. In the cold-rolled steel sheet of the present invention, the concentration of carbon in the austenite phase is increased by increasing the activity of carbon by being dissolved in ferrite, thereby suppressing the formation of carbide phase and increasing the stability of retained austenite. Play a role. Moreover, there exists an effect which the intensity | strength of a steel plate increases by solid solution. If the Si content is less than 1.0%, the strength of the steel is reduced, and the effect of suppressing the formation of carbides such as carbide phase is reduced. If it exceeds 1.9%, a hot rolling scale is induced and plating is performed. The weldability is deteriorated and the weldability is also deteriorated. Therefore, the composition range of Si is limited to 1.0 to 1.9%.

  The content of manganese (Mn) is 1.5 to 2.5%. Mn is a component that increases the strength by increasing the curability and easily generating low-temperature transformation phases such as acicular ferrite and bainite. Mn also stabilizes austenite. Therefore, when Mn is added at less than 1.5%, the above effect cannot be expected. When Mn is added at more than 2.5%, the weldability is lowered, and a segregation zone is formed in the center of the plate during hot rolling. In addition, the formation of inclusions causes hydrogen embrittlement. Therefore, the Mn content is limited to 1.5 to 2.5%.

  The content of aluminum (Al) is 0.5 to 1.6%. Al is less effective in solid solution strengthening than Si, but as a ferrite stabilizing element, it is a ferrite stabilizing element that shows the effect of solid solution strengthening. It suppresses the formation of carbides such as carbides, and reduces the carbon concentration in residual austenite. It is a component that increases the stability and increases the stability. If the Al content is less than 0.5%, the stability of austenite decreases, so it becomes difficult to suppress the formation of carbides. If it exceeds 1.6%, the austenite fraction decreases, so the ductility of the steel sheet. Is relatively lowered and the surface properties are deteriorated. Therefore, the Al content is limited to 0.5 to 1.6%.

  The content of titanium (Ti) is 0.005 to 0.03%. Ti is a component that forms TiN so as to suppress the formation of AlN nitride caused by the bonding of Al and N and to allow Al to function. If Ti is less than 0.005%, such a role cannot be expected, and even if added in excess of 0.03%, no further effect can be expected. Therefore, the Ti content is limited to 0.005 to 0.03%.

  The content of boron (B) is 5 to 30 ppm. B is a component that improves the curability even if a small amount is added to the steel. When 5 ppm or more is added, it segregates at the austenite grain boundary at a high temperature, thereby suppressing ferrite formation and improving the curability. When added over 30 ppm, the recrystallization temperature of the steel sheet is increased, so that the drawability of the steel sheet is lowered and the weldability is deteriorated. Therefore, the B content is limited to 5-30 ppm.

  The content of antimony (Sb) is 0.01 to 0.03%. When Sb is added in an appropriate amount of 0.01 to 0.03%, the surface properties of the steel sheet are improved. However, when Sb is added in an amount exceeding 0.03%, Sb concentration occurs on the surface of the steel plate, resulting in poor surface properties. Become. Sb is a kind of impurity inevitably contained in steel materials, and an Sb content of less than 0.01% is found in steel produced without a specific purpose. When the Sb content is less than 0.01%, concentration occurs on the surface. However, since the amount is extremely small to cause changes in surface characteristics, the content is made 0.01% or more. Therefore, the Sb content is limited to 0.01 to 0.03%.

  In the present invention, 1.75 << Si + Al << 3.25 is satisfied. Both Si and Al play a role in suppressing the carbide formation of the steel, and increase the content of solute carbon in the retained austenite, thereby improving the stability of the retained austenite. Therefore, when considering the strength and elongation rate of TRIP steel, it is necessary to control the contents of these two components at the same time. However, if the total content of the two components exceeds 3.25%, the formation of surface oxides As a result, the surface quality such as plating property is deteriorated, and the strength and ductility may be reduced due to the reduction of the austenite fraction occurring during the two-phase annealing treatment. Further, if the total content of the two components is less than 1.75%, the effect of basic solid solution strengthening necessary for producing a TRIP steel having a target tensile strength of 780 MPa or more is reduced, and the stability of retained austenite is reduced. It becomes difficult to secure. Therefore, the content of Si + Al is limited to 1.75 to 3.25%.

  The present invention includes the above composition, the balance Fe and inevitable impurities.

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

The present invention, a steel slab satisfying the above composition after hot rolling at A ₃ above the austenite region, to primary cooling at a cooling rate of 30 to 200 ° C. / s. At this time, when the cooling rate is less than 30 ° C./s, a pearlite structure is formed, and it is difficult to secure a target material. When the cooling rate exceeds 200 ° C./s, a temperature deviation occurs in the steel sheet, and the material is distorted by residual stress. May occur. For this reason, a cooling rate is restrict | limited to 30-200 degreeC / sec.

  After the primary cooling, the steel sheet is maintained in a temperature range of 600 to 800 ° C. This means that the steel sheet is cooled by natural convection in a normal temperature atmosphere without forced cooling. If the temperature is less than 600 ° C., it is difficult to secure the ferrite phase formation fraction, and if it exceeds 800 ° C., excessive ferrite may be formed or a pearlite structure may be formed.

  After the air cooling, the steel sheet is secondarily cooled at a cooling rate of 50 to 200 ° C./s, and then wound at room temperature to 300 ° C. When the cooling rate is less than 50 ° C./s, it is difficult to obtain a target structure due to the formation of the bainite phase, and when it exceeds 200 ° C./s, an excessive martensite phase is formed, and the shape distortion of the rolled sheet may occur. is there. Therefore, the secondary cooling rate is limited to 50 to 200 ° C./s.

  In addition, when the coiling temperature exceeds 300 ° C, it is difficult to secure a target structure due to the formation of a bainite phase, and when coiling at room temperature to 300 ° C, the base structure has a fine lath-shaped martensite structure. However, it has a high dislocation density after hot rolling and a uniform solute carbon distribution. Therefore, the coiling temperature is limited to room temperature to 300 ° C.

  After winding, cold rolling is performed at a reduction rate of 30 to 50%. Usually, cold rolling is performed at a reduction rate of 60%, but in the present invention, it is performed at a reduction rate of 30 to 50%. As a result, among the martensite structure and ferrite structure generated after coiling by cold rolling, the dislocation density can be sufficiently increased in the ferrite structure, and the re-dissolution of carbon and the formation of the austenite phase during annealing It occurs uniformly. If the cold rolling reduction is less than 30%, the dislocation density in the ferrite structure is not sufficient, so that the target structure cannot be obtained, and if it exceeds 50%, fine cracks occur at the boundary between the martensite phase and the ferrite phase. It tends to occur, and crack defects occur particularly at the rolled plate edge. Therefore, the cold reduction rate is limited to 30 to 50%.

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

  In the present invention, hot dip galvanization or alloying hot dip galvanization is performed after cold rolling. After the annealing step, the steel sheet is passed through a plating bath containing molten zinc, and the plating layer is adhered to the surface layer with a certain thickness. At this time, the temperature of the plating bath is preferably 450 to 500 ° C., and a hot-dip galvanized steel sheet is produced by a step of slow cooling at a rate of 30 ° C./s or less.

  In addition, after the plated steel sheet that has passed through the plating bath is directly heated in a temperature range of 500 to 600 ° C. to perform alloying heat treatment of the hot dip galvanized layer, it is gradually cooled at a rate of 30 ° C./s or less, Produces galvannealed steel sheet.

  Hereinafter, the structure of the steel sheet produced by this method will be described in detail.

  In the present invention, the hot-rolled steel sheet after hot rolling has a martensite fraction of 30 to 70%. After hot rolling, the lath-shaped fine martensite structure produced in the primary and secondary cooling processes induces uniform austenite transformation in the annealing process after cold rolling, and thereby stabilized residual austenite Increase the fraction. In the conventional method, a hot-rolled steel sheet is produced by only one cooling process. In this case, the hot-rolled steel sheet contains a pearlite structure in which coarse carbides are present. Coarse carbide is re-dissolved in the annealing process after cold rolling, but because the size of the carbide is coarse, it remains easily without being re-dissolved even at a high temperature of 700 ° C. or higher. Austenite is hardly formed in a region where there is relatively no carbide, starting to form mainly in the vicinity of a carbide having a high carbon concentration. Accordingly, the austenite fraction becomes lower during the two-phase annealing and a local deviation occurs, and the residual austenite fraction generated in the cooling step after annealing decreases. The martensite phase in the hot-rolled steel sheet can improve such disadvantages.

  In the present invention, when the martensite phase is formed through the primary and secondary cooling steps, the hot-rolled steel sheet can be produced in a state where carbides are hardly formed. Such a hot-rolled steel sheet containing a martensite phase is re-dissolved immediately when fine carbides are formed in a martensite structure having a high dislocation density during annealing heat treatment. For this reason, the deviation of the carbon concentration in the steel sheet of the present invention is considerably reduced as compared with the steel sheet having a pearlite structure. Therefore, since the austenite phase formed during annealing is uniformly distributed and formed in the vicinity of the grain boundaries and the lath martensite structure, the retained austenite fraction and stability can be improved.

  In addition, by utilizing a relatively low cold reduction rate by utilizing the dislocation density that is higher due to the martensite phase, it is possible to suppress the formation of fine cracks that are likely to occur at the edge portion after cold rolling. At this time, if the martensite fraction of the hot-rolled steel sheet is less than 30%, the effect of increasing the retained austenite fraction is weak, and if it is 70% or more, fine cracks are generated at the edge portion in cold rolling. To do. Therefore, in the present invention, the martensite structure fraction of the hot-rolled steel sheet is set to 30 to 70%.

  In the cold-rolled steel sheet that has undergone the cold rolling and annealing processes of the present invention, retained austenite is generated in the ferrite and bainite structure. At this time, the fraction of retained austenite in the steel sheet of the present invention is 5 to 15%. The retained austenite phase appears in the lath shape due to the influence of the lath-shaped martensite structure in the hot-rolled steel sheet, and is more stable than the remaining austenite in other shapes. Moreover, the fraction of a bainite structure is 20 to 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 and a tensile strength of 780 to 980 MPa and an elongation of 28% or more can be produced.

  Hereinafter, the present invention will be described more specifically with reference to examples.

(Example)
Table 1 shows the composition range of the steel used in the examples. In the present invention, in order to obtain the target tensile strength and elongation, appropriate adjustment of the components is important.

  The steel slab having the composition shown in Table 1 is cooled and wound through a hot rolling process, and then cold-rolled and annealed. The cooling conditions, the cold reduction rate, the tensile test results, and the like are shown in Table 2. .

  The tensile properties in the mechanical properties of the material vary depending on the ingredients and the microstructure production conditions. The steel sheet of the present invention induces uniform austenite transformation in the annealing process due to the fine martensitic structure of lath formed immediately after hot rolling, thereby increasing the fraction of stabilized retained austenite and causing deformation. When transformed, it transforms to obtain excellent mechanical properties.

  In addition, by utilizing a relatively low cold reduction rate by utilizing the dislocation density that has become higher due to the martensite phase, an attempt was made to suppress the formation of fine cracks that are likely to occur at the edge after cold rolling. In order to obtain such characteristics, the components proposed in the present invention must be basically satisfied, and the cooling process immediately after hot rolling, the martensite fraction, and the cold reduction rate must be adjusted optimally. Don't be.

  Table 2 shows the results of the measurement of the cooling process immediately after hot rolling, the martensite fraction, the cold reduction ratio, the mechanical properties, and the degree of occurrence of cracks at the edge portion.

  As shown in Table 2, the comparative materials 9 to 15 do not satisfy the basic component requirements, and even if the cooling conditions, the martensite fraction, and the cold rolling reduction conditions are appropriately adjusted during the manufacturing process, It can be seen that the target strength and ductility could not be obtained in the invention.

  Moreover, 2-2, 2-3, 3-2, 3-3, 4-2, 4-3, 5-2, 7-2, 7-3, and 8-2 which are comparative materials satisfy the component requirements. However, it can be seen that the requirements such as the cooling condition, the martensite fraction, the cold reduction rate, etc. are not satisfied, and the target mechanical properties cannot be obtained, or fine cracks are generated at the edge portion.

  In Table 2, the crack length at the edge portion was measured by observing the martensite fraction with an optical microscope of 200 magnification.

  In particular, this is a result obtained by measuring a fine structure at a thickness of t / 4 with a 2% Nital etching solution and then measuring with an image analyzer. The fine cracks generated in the edge part were selected as an average value after randomly selecting the cold-rolled rolled sheet edge part and selecting 30 or more longest cracks generated within a length of 100 mm.

  The steel sheet of the present invention has a tensile strength of 780 to 980 MPa and an elongation of 28% or more by controlling the component structure and production conditions. Therefore, the steel sheet of the present invention can be used for structural parts having high strength and workability. Moreover, the steel plate of this invention has an effect which improves economical efficiency by preventing that a crack generate | occur | produces in an edge part.

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

Claims (7)

By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and unavoidable impurities, and satisfies the 1.75 ≦ Si + Al ≦ 3.25% A high-strength hot-rolled steel sheet, wherein the microstructure of the steel sheet after hot rolling consists of 30 to 70% martensite and the remaining ferrite. By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005~0.03%, B: 5~30ppm, and Sb: 0.01 to 0.03%, and the balance Fe and unavoidable impurities, and satisfies the 1.75 ≦ Si + Al ≦ 3.25% A high-strength cold-rolled steel sheet, wherein the microstructure of the steel sheet after cold rolling using the hot-rolled steel sheet according to claim 1 is 5 to 15% residual austenite, 20 to 40% bainite, and the remaining ferrite A high-strength cold-rolled steel sheet characterized by comprising:   The high-strength cold-rolled steel sheet according to claim 2, wherein the cold-rolled steel sheet has a tensile strength of 780 to 980 MPa and an elongation of 28% or more.   A high-strength hot-dip galvanized steel sheet comprising a hot-dip galvanized layer on the high-strength cold-rolled steel sheet according to claim 2 or 3. By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005 to 0.03%, B: 5 to 30 ppm, and Sb: 0.01 to 0.03%, balance Fe and inevitable impurities, and satisfy 1.75 ≦ Si + Al ≦ 3.25% the steel slab comprising the steps of hot rolling at _{a 3} or more temperature and primary cooling in a temperature range of 600 to 800 ° C. at a cooling rate of 30 to 200 ° C. / s,
Air-cooling the first cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C .;
Secondary cooling the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C./s in a temperature range of room temperature to 300 ° C .;
The method for producing a high-strength hot-rolled steel sheet according to claim 1, comprising a step of winding the secondary-cooled hot-rolled steel sheet in a temperature range of room temperature to 300 ° C. By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005 to 0.03%, B: 5 to 30 ppm, and Sb: 0.01 to 0.03%, balance Fe and inevitable impurities, and satisfy 1.75 ≦ Si + Al ≦ 3.25% the steel slab was hot rolled at _{a 3} or higher, at a cooling rate of 30 to 200 ° C. / s, the method comprising the primary cooling in the temperature range of 600 to 800 ° C.,
A step of air-cooling the first cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C;
Secondary cooling the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C./s in a temperature range of room temperature to 300 ° C .;
Winding the secondary-cooled hot-rolled steel sheet in a temperature range from room temperature to 300 ° C .;
The method for producing a high-strength cold-rolled steel sheet according to claim 2, comprising cold rolling and annealing the rolled hot-rolled steel sheet at a rolling reduction of 30 to 50%. By weight, C: 0.1-0.25%, Si: 1.0-1.9%, Mn: 1.5-2.5%, Al: 0.5-1.6%, Ti: 0.005-0.03%, B: 5-30 ppm, and Sb: 0.01-0.03%, balance Fe and inevitable impurities, and satisfying 1.75 ≦ Si + Al ≦ 3.25% the slab was hot rolled at _{a 3} or higher, at a cooling rate of 30 to 200 ° C. / s, the method comprising the primary cooling in the temperature range of 600 to 800 ° C.,
Air-cooling the first cooled hot-rolled steel sheet in a temperature range of 600 to 800 ° C .;
Secondary cooling the air-cooled hot-rolled steel sheet at a cooling rate of 50 to 200 ° C./s in a temperature range of room temperature to 300 ° C .;
Winding the secondary-cooled hot-rolled steel sheet in a temperature range from room temperature to 300 ° C .;
Cold rolling and annealing the rolled hot-rolled steel sheet at a rolling reduction of 30 to 50%,
A method for producing a high-strength hot-dip galvanized steel sheet according to claim 4, comprising the step of hot-dip galvanizing the annealed cold-rolled steel sheet.

Images