High-strength thin steel sheet having excellent formability and method for producing same

Abstract

Disclosed are a high-strength thin steel sheet and a method for manufacturing the same, the thin steel sheet comprising, in wt%, 0.001-0.004% of C, 0.5% or less (excluding 0%) of Si, 1.2% or less (excluding 0%) of Mn, 0.005-0.12% of P, 0.01% or less of S, 0.01% or less of N, 0.1% or less (excluding 0%) of acid-soluble aluminum, 0.01-0.04% of Ti, and the balance Fe and unavoidable impurities, the contents of Ti, N, and S satisfying the following relational expression 1, the ratio (b / a) of the average random strength ratio (b) of (111) [1-10]-(111) [-1-12] orientation to the average random strength ratio (a) of (001) [1-10]-(110) [1-10] orientation at a t / 4 (t: thickness of the thin steel sheet) position in the sheet thickness direction is 2.3 or more, and the bake hardening amount (BH) is 4 MPa or more. [relational expression 1]-0.02 < = [Ti]-(24 / 7) [N]-(3 / 2) [S] < = 0.025 (wherein [Ti], [N], and [S] each represent the content (wt%) of the element).

Description

[0001] This application is a divisional application of a Chinese patent application with an application date of August 4, 2017, a Chinese patent application number of 201780046048.8, and an invention name of “High-strength thin steel sheet with excellent formability and its manufacturing method”, and this application claims priority to the Korean application with application number 10-2016-0102946. Technical Field

[0002] The present invention relates to a high-strength thin steel sheet and a method for producing the same, and more particularly, to a high-strength thin steel sheet having excellent formability and which can be preferably used as a material for an automobile exterior plate or the like and a method for producing the same. Background Art

[0003] Steel used as materials for the interior and exterior panels of automobiles (doors, hoods, fenders, floors, etc.) is required to have not only high strength but also excellent formability. This is to ensure the safety of passengers in accidents and to improve fuel efficiency by reducing the weight of the vehicle body.

[0004] However, the increase in steel sheet strength leads to poor formability, making it difficult to satisfy both factors (strength and formability) at the same time. In particular, in parts that require higher formability, such as door inners or rear floors, poor forming such as cracks often occur during processing, so high-strength steel is rarely used in such parts.

[0005] So far, the known steel sheets with excellent strength and formability developed include so-called interstitial free steel (IF steel). This steel is obtained by removing solid solution elements such as carbon (C), nitrogen (N), and sulfur (S) by adding strong carbonitride-forming elements such as titanium (Ti) and / or niobium (Nb), thereby ensuring both strength and formability. Representative examples are disclosed in Patent Documents 1 to 4. However, the IF steel shows an average plastic anisotropy coefficient (Lankford value, r value) of 1.5 to 1.8, which is far from sufficient to replace existing parts using soft cold-rolled steel sheets of deep drawing quality (DDQ) grade.

[0006] (Patent Document 1) Japanese Patent Publication No. 1992-280943

[0007] (Patent Document 2) Japanese Patent Publication No. 1993-070836

[0008] (Patent Document 3) Japanese Patent Publication No. 1993-263184

[0009] (Patent Document 4) Japanese Patent Publication No. 1998-096051 Summary of the invention

[0010] Technical issues to be solved

[0011] One of the objects of the present invention is to provide a high-strength thin steel sheet having excellent formability and a method for producing the same.

[0012] Technical Solution

[0013] One aspect of the present invention provides a high-strength steel sheet, which comprises, by weight%, C: 0.001-0.004%, Si: 0.5% or less (except 0%), Mn: 1.2% or less (except 0%), P: 0.005-0.12%, S: 0.01% or less, N: 0.01% or less, acid-soluble aluminum: 0.1% or less (except 0%), Ti: 0.01-0.04%, and the balance of Fe and inevitable impurities. The content of Ti, N and S satisfies the following relationship 1, and the ratio (b / a) of the average random intensity ratio (b) of (111) [1-10] to (111) [-1-12] orientation to the average random intensity ratio (a) of (001) [1-10] to (110) [1-10] orientation at the t / 4 position (t: thickness of the thin steel plate) in the plate thickness direction is 2.3 or more, and the bake hardening amount (BH) is 4 MPa or more.

[0014] [Relationship 1] -0.02≤[Ti]-(24 / 7)[N]-(3 / 2)[S]≤0.025

[0015] ([Ti], [N] and [S] respectively represent the content (weight %) of the element.)

[0016] Another aspect of the present invention provides a method for manufacturing a high-strength thin steel plate, characterized in that it includes the following steps: hot rolling a steel billet to obtain a hot-rolled steel plate, wherein the steel billet contains, by weight%, C: 0.001-0.004%, Si: 0.5% or less (except 0%), Mn: 1.2% or less (except 0%), P: 0.005-0.12%, S: 0.01% or less, N: 0.01% or less, acid-soluble aluminum: 0.1% or less (except 0%), Ti: 0.01-0.04%, and the balance of Fe and inevitable impurities; hot rolling the hot-rolled steel plate at a temperature of 450-750°C, Coiling; cold rolling the coiled hot-rolled steel sheet at a reduction rate of more than 75% to obtain a cold-rolled steel sheet; heating the cold-rolled steel sheet to an annealing temperature of 830 to 880°C, and then maintaining the annealing time at the annealing temperature for 30 to 80 seconds to perform continuous annealing; cooling the continuously annealed cold-rolled steel sheet to a temperature below 650°C at a rate of 2 to 10°C / second; level rolling the cooled cold-rolled steel sheet at a reduction rate of 0.3 to 1.6%; when heating the cold-rolled steel sheet, the average heating rate from (recrystallization starting temperature + 20)°C to the annealing temperature is less than 5°C / second.

[0017] Beneficial Effects

[0018] One of the effects of the present invention is that the thin steel sheet of the present invention is excellent in strength and formability and can be preferably used as a material for automobile exterior panels and the like. BRIEF DESCRIPTION OF THE DRAWINGS

[0019] Figure 1 This is a graph analyzing the degree of development of the texture of Inventive Example 1.

[0020] Best Mode for Carrying Out the Invention

[0021] As a result of in-depth research conducted by the inventors to solve the above-mentioned problems of the prior art, it has been confirmed that by removing solid solution elements such as carbon (C), nitrogen (N), and sulfur (S) by adding titanium (Ti) alone, which is a strong carbonitride-forming element in steel, or by adding titanium (Ti) and niobium (Nb) in combination, and at the same time appropriately controlling the positional distribution of carbides and the like produced as a result of the removal of the solid solution elements, and controlling the texture, the strength and drawability can be significantly improved, and the solid solution carbon re-melted during the residual annealing process at an appropriate level can be significantly improved. Bake hardenability can thus be completed, thereby completing the present invention.

[0022] Next, a high-strength thin steel sheet having excellent formability according to one aspect of the present invention will be described in detail.

[0023] First, the alloy components and preferred content ranges of the high-strength steel sheet are described in detail. Unless otherwise specified, the content of each component below is based on weight.

[0024] C: 0.001~0.004%

[0025] Carbon is an interstitial solid solution element, and has a great influence on the formation of the texture of the steel sheet during cold rolling and annealing. In particular, when the content of solid solution carbon in the steel increases, the growth of grains with {111} texture that is beneficial to drawing workability is suppressed, and the growth of grains with {110} and {100} textures is promoted, thereby deteriorating the drawability of the thin steel sheet. Furthermore, when the carbon content is too much, the content of Ti required to precipitate carbon into carbides is too much, which is not only disadvantageous in terms of economy, but also a large amount of fine TiC carbides are distributed in the steel, which has the problem of sharply reducing the drawability. Therefore, in the present invention, the upper limit of the carbon content is controlled to 0.004%, preferably 0.0035%. In addition, although the lower the carbon content, the more conducive it is to improve the drawability, when the carbon content is too low, there is a problem of sharply reducing the bake hardenability of the thin steel sheet. Therefore, in the present invention, the lower limit of the carbon content is controlled to 0.001%, preferably 0.0012%.

[0026] Si: 0.5% or less (except 0%)

[0027] Silicon helps to improve the strength of thin steel sheets through solid solution strengthening. However, when the silicon content is too high, it will cause surface scale defects, thereby reducing the surface properties of the plated surface. Therefore, in the present invention, the upper limit of the silicon content is controlled to 0.5%, preferably 0.05%. In addition, the present invention does not specifically limit the lower limit of the silicon content, but it can be preferably 0.001%, and more preferably 0.002%.

[0028] Mn: 1.2% or less (except 0%)

[0029] Manganese is a solid solution strengthening element, which not only helps to improve the strength of steel, but also precipitates S in steel as MnS, thereby inhibiting the occurrence of plate fracture caused by S during hot rolling and high temperature embrittlement. However, when the manganese content is too much, the excess Mn will solid dissolve, thereby causing the problem of poor drawability. In the present invention, the upper limit of the manganese content is controlled to be less than 1.2%, preferably controlled to be less than 1.0%, and more preferably controlled to be less than 0.8%. In addition, the present invention does not specifically limit the lower limit of the manganese content, but it can be preferably 0.01%, and more preferably 0.1%.

[0030] P: 0.005~0.12%

[0031] Phosphorus is an element that is very effective in improving the strength of steel, and has a very excellent solid solution effect and little effect on drawability. In the present invention, the lower limit of the phosphorus content is controlled to 0.005%, preferably 0.008%, and more preferably 0.010%. However, when the phosphorus content is too high, the excess P will precipitate as FeTiP, which has the problem of deteriorating the drawability. In the present invention, the upper limit of the phosphorus content is controlled to 0.12%, preferably 0.10%, and more preferably 0.08%.

[0032] S: 0.01% or less, N: 0.01% or less

[0033] Sulfur and nitrogen are impurities that inevitably exist in steel. To ensure excellent welding properties, it is preferable to control the content of sulfur and nitrogen to be as low as possible. In the present invention, in order to ensure appropriate welding properties, the upper limits of the content of sulfur and nitrogen are controlled to 0.01% or less.

[0034] Acid soluble aluminum: less than 0.1% (except 0%)

[0035] Acid-soluble aluminum helps improve the drawability and ductility of thin steel sheets by precipitating AlN. However, when the content of acid-soluble aluminum is too high, excessive Al-based inclusions are formed during steelmaking operations, which has the problem of generating internal defects in the steel sheet. In the present invention, the upper limit of the acid-soluble aluminum content is controlled to 0.1%, preferably controlled to 0.08%, and more preferably controlled to 0.05%. In addition, the present invention does not specifically limit the lower limit of the acid-soluble aluminum content, but it can be preferably 0.01%, and more preferably 0.02%.

[0036] Ti: 0.01~0.04%

[0037] Titanium reacts with dissolved carbon and dissolved nitrogen during hot rolling to precipitate Ti-based carbonitrides, so titanium is an element that is very helpful in improving the drawability of thin steel sheets. In the present invention, the lower limit of the titanium content is controlled to be above 0.01%, preferably above 0.012%, and more preferably above 0.015%. However, when the titanium content is too high, the remaining Ti after reacting with dissolved carbon and dissolved nitrogen combines with P to form excessive FeTiP precipitates, which may deteriorate the drawability, and because a large amount of TiC or TiN precipitates are distributed in the steel, the dissolved carbon content becomes too low, which may deteriorate the bake hardenability of the thin steel sheet. In the present invention, the upper limit of the titanium content is controlled to be 0.04%, preferably 0.03%.

[0038] In addition, the remaining amount of Fe and inevitable impurities are included. However, in the usual manufacturing process, unnecessary impurities are inevitably mixed from the raw materials or the surrounding environment, so these impurities cannot be excluded. For technicians with general knowledge in this technical field, these impurities are well known, so all of them are not specifically mentioned in this specification. In addition, the addition of effective components other than the above composition is not excluded. In particular, in order to further improve the mechanical and physical properties of the steel plate, the following components may be further included.

[0039] Nb: 0.005~0.04%

[0040] During hot rolling, niobium causes dissolved carbon to precipitate in the form of (Ti, Nb)C complex carbides, thereby facilitating the formation of texture during annealing. In particular, when an appropriate amount of Nb is added, it has the effect of improving the plastic anisotropy in each direction (0°, 45°, 90°). Therefore, compared with the 90° direction, the plastic deformation anisotropy (r value) in the 0° direction and the 45° direction increases, resulting in the material's planar anisotropy (△r, Planar anisotropy) reaching near zero (0), and showing the characteristic that the r value is evenly distributed on the plate surface, thereby having the advantage of preventing the material from having an ear (earing) shape during forming. In order to obtain the above-mentioned effect in the present invention, it is preferred that the lower limit of the niobium content be controlled to 0.005% or more, and more preferably to 0.008% or more. However, when the niobium content is too high, most of the dissolved carbon in the steel precipitates as fine NbC, and even after annealing, the dissolved carbon is hardly dissolved again, so that the bake hardenability deteriorates, and since the amount of precipitation of fine (Ti, Nb) C complex carbides is relatively small, not only the drawability (r value) deteriorates, but also there is a problem of material deterioration caused by the increase in recrystallization temperature. The upper limit of the niobium content is preferably 0.04%, more preferably 0.03%, and further preferably 0.025%.

[0041] B: 0.002% or less (except 0%)

[0042] Boron suppresses secondary processing brittleness caused by P in steel. However, when the boron content is too high, the ductility of the steel sheet decreases. Therefore, in the present invention, the upper limit of the boron content is controlled to be 0.002% or less, preferably 0.0015% or less. In addition, the present invention does not specifically limit the lower limit of the boron content, but it can be preferably 0.0003%, and more preferably 0.0005%.

[0043] In addition, when designing an alloy of a thin steel sheet having the above-mentioned composition range, the contents of Ti, N and S preferably satisfy the following relational expression 1. If the value of [Ti]-(24 / 7)[N]-(3 / 2)[S] is less than -0.02, the content of Ti for precipitating C in the steel as TiC is absolutely insufficient, and the r value as a workability evaluation index is significantly reduced. On the other hand, when the value of [Ti]-(24 / 7)[N]-(3 / 2)[S] exceeds 0.025, in addition to the formation of TiC precipitates that are advantageous for workability, FeTiP precipitates are also formed, which significantly hinders the development of {111} orientation during annealing. More preferably, the value of [Ti]-(24 / 7)[N]-(3 / 2)[S] is controlled to -0.01 to 0.01.

[0044] [Relationship 1] -0.02≤[Ti]-(24 / 7)[N]-(3 / 2)[S]≤0.025

[0045] ([Ti], [N] and [S] respectively represent the content (weight %) of the element.)

[0046] Next, the structure, precipitates, etc. of the high-strength thin steel sheet will be described in detail.

[0047] The arrangement of certain planes and orientations generated inside the crystal is called texture. The shape of these textures developed in a certain direction is called fiber texture. Texture is closely related to drawing properties. It is known that the higher the surface strength value of the γ-fiber texture that forms the {111} plane perpendicular to the rolling surface, the better the drawing processability. Usually, α-fiber texture is defined as RD / / <110> , the γ-fiber texture is defined as ND / / <111> .

[0048] In addition, the inventors of the present invention have learned that in order to form the γ-fiber texture as described above, the ratio of the average random intensity ratio (b) of the γ-fiber texture ((111)[1-10] to (111)[-1-12] orientation) to the average random intensity ratio (a) of the α-fiber texture ((001)[1-10] to (110)[1-10] orientation) at the t / 4 position (t: thickness of the steel plate) from the surface of the steel plate in the thickness direction of the plate is very important. More specifically, the present inventors have confirmed that when the ratio (b / a) of the average random intensity ratio (b) of the (111) [1-10] to (111) [-1-12] orientation at the position t / 4 (t: thickness of the thin steel plate) from the steel plate surface to the plate thickness direction to the average random intensity ratio (a) of the (001) [1-10] to (110) [1-10] orientation is ensured to be 2.3 or more, the average plastic anisotropy coefficient (Lankford value, r value) can be ensured to be 1.9 or more, thereby ensuring excellent drawability. In addition, the higher the average random intensity ratio of the γ-fiber texture ((111) [1-10] to (111) [-1-12] orientation), the more advantageous it is for drawability, and therefore, the upper limit thereof is not particularly limited in the present invention.

[0049] In particular, the present invention has confirmed that when molding automobile parts, it is necessary to ensure excellent drawability in multiple directions in order to mold a complete part without cracks, rather than ensuring excellent drawability in a specific direction, and it has been found that the development degree of all γ-fiber textures in the range of 0 to 90° and the value expressed can indicate complete moldability. That is, for all directions of the γ-fiber texture at 0° ((111) [1-10]), 30° ((111) [1-21]), 60° ((111) [0-11]), and 90° ((111) [-1-12]), the higher the development of the average random intensity ratio is overall, the more advantageous it is.

[0050] The average plastic anisotropy coefficient (Lankford value, r value) obtained from the plastic anisotropy coefficients measured in various directions with respect to the rolling direction is a representative material property value indicating drawability, and this value is calculated by the following formula 1.

[0051] r value = (r0 + r90 + 2r45) / 4 (Formula 1)

[0052] (where ri represents the plastic anisotropy coefficient measured in the test piece collected at an angle i° to the rolling direction.)

[0053] In the above formula, the larger the r value is, the deeper the molded cup can be during drawing, which indicates excellent drawability. A thin steel sheet according to one embodiment of the present invention has an r value of 1.9 or more, and thus exhibits excellent drawability.

[0054] According to one example, the average grain size of the high-strength thin steel plate can be 5 μm or more, preferably 7 μm or more. Wherein, the average grain size refers to the average equivalent circular diameter of the grains. In the present invention, the coarser the grain size, the more favorable it is for formability, so it is advantageous to ensure coarse grains as much as possible. To this end, by controlling the composition, the C content is reduced to the level of extremely low carbon steel below 40 ppm, and the precipitation of carbides is controlled as effectively as possible, thereby ensuring the growth of grains during annealing. This is because the coarser the size of the grains, the easier it is for carbides in the grains to precipitate relative to the grain boundaries, thereby significantly reducing the possibility of cracks during processing. In addition, the larger the average grain size, the more favorable it is for formability, so the upper limit of the average grain size is not particularly limited in the present invention, but considering that annealing at a high temperature of more than 860°C for the growth of grains has the possibility of causing damage to the refractory bricks in the annealing furnace, the upper limit of the average grain size can be limited to 20 μm.

[0055] According to one example, the P of the high-strength steel sheet of the present invention is defined by the following mathematical formula 1: in It can be 80% or more, preferably 82% or more. in ) is less than 80%, that is, when a large amount of carbides are precipitated in the grain boundaries, the possibility of cracking during processing is significantly increased, thereby having the possibility of deterioration of ductility and drawability. in ) is higher, the more conducive to improving ductility and drawability. Therefore, in the present invention, the ratio (P in The upper limit of ) is not particularly limited. In addition, the above-mentioned carbide refers to TiC single carbide, NbC single carbide or (Ti, Nb)C composite carbide.

[0056] [Mathematical formula 1]

[0057] P in (%) = {N in / (N in +N gb )}×100

[0058] (Where N in N is the number of carbides with a circle equivalent diameter of 20 nm or less existing in the grains, gb is the number of carbides with a circle equivalent diameter of 20 nm or less existing at the grain boundary.

[0059] According to one example, the high-strength steel sheet of the present invention has a 2) may contain less than 0.2 FeTip precipitates, preferably less than 0.1 FeTip precipitates. The FeTiP precipitates are mainly precipitated in needle shape, thereby reducing the development of {111} orientation during annealing. When the FeTiP precipitates formed exceed 0.2 / μm 2 In addition, the smaller the number of FeTiP precipitates per unit area, the more advantageous it is for improving the drawability. Therefore, the lower limit of the number of FeTiP precipitates is not particularly limited in the present invention.

[0060] According to one example, the high-strength thin steel sheet of the present invention has a bake hardening density (BH) of 4 MPa or more, more preferably 10 MPa or more, and further preferably 15 MPa or more, thereby showing excellent bake hardenability.

[0061] According to one example, the high-strength thin steel sheet of the present invention has a thickness of 0.8 mm or less, and has a product value of a yield strength (YS, MPa) and an average plastic anisotropy coefficient (Lankford value, r value) of 290 MPa or more, and thus has very excellent dent resistance and formability indicating resistance to external physical forces, and can therefore be preferably used as a material for automobile outer panels.

[0062] The high-strength thin steel sheet of the present invention described above can be manufactured by various methods, and the manufacturing method is not particularly limited. However, as a preferred example, it can be manufactured by the following method.

[0063] Next, a method for producing a high-strength thin steel sheet having excellent formability according to another aspect of the present invention will be described in detail.

[0064] First, a steel slab having the above-mentioned component system is hot rolled to obtain a hot-rolled steel sheet.

[0065] According to one example, hot finishing rolling can be performed at a temperature in the austenite single phase region (temperature above Ar3°C). If the hot finishing rolling temperature is lower than Ar3°C, the possibility of two-phase region rolling is high, which may cause material inhomogeneity. For reference, Ar3°C can be calculated by the following formula 2.

[0066] [Formula 2]

[0067] Ar3(℃)=910-310[C]-80[Mn]-20[Cu]-15[Cr]-55[Ni]-80[Mo]

[0068] (Wherein, [C], [Mn], [Cu], [Cr], [Ni], and [Mo] each represent the weight % of the element.)

[0069] Next, the hot rolled steel sheet is coiled.

[0070] At this time, the coiling temperature is preferably 450 to 750° C., more preferably 500 to 700° C. When the coiling temperature is lower than 450° C., a large amount of FeTiP precipitates are precipitated, thereby reducing the drawability and there is a possibility of causing the sheet shape distortion. On the other hand, when the coiling temperature exceeds 750° C., the precipitates are coarsened and the dissolved carbon in the annealing process is difficult to re-melt, which may deteriorate the bake hardening amount (BH).

[0071] According to one example, the average cooling rate from the hot finishing rolling temperature to the coiling temperature may be 10 to 200°C / sec. If the average cooling rate is less than 10°C / sec, the growth of ferrite grains is uneven, and FeTiP precipitates are formed, making it difficult to ensure the desired formability in the present invention. On the other hand, when the average cooling rate from the hot finishing rolling temperature to the coiling temperature exceeds 200°C / sec, the temperature of the hot rolled steel sheet is uneven due to excessive cooling, and thus the shape of the hot rolled steel sheet is deteriorated.

[0072] Next, the coiled hot-rolled steel sheet is cold-rolled to obtain a cold-rolled steel sheet.

[0073] At this time, the cold rolling reduction rate is preferably 75% or more. If the cold rolling reduction rate is less than 75%, the growth of the γ-fiber texture is insufficient, resulting in a problem of poor drawability. In addition, the higher the cold rolling reduction rate, the more conducive to the growth of the γ-fiber texture. Therefore, the upper limit of the cold rolling reduction rate is not particularly limited in the present invention. However, when the cold rolling reduction rate is too high, the roll load during rolling increases, so that the shape of the steel plate deteriorates. Therefore, considering this point, the upper limit of the cold rolling reduction rate can be limited to 85%.

[0074] Next, the cold rolled steel sheet is subjected to continuous annealing.

[0075] At this time, the annealing temperature (T) is preferably 830 to 880°C, more preferably 840 to 870°C. If the annealing temperature (T) is lower than 830°C, the γ-fiber texture that is beneficial to workability cannot be fully grown, and there is a possibility that the drawability is deteriorated, and the precipitates cannot be melted again during the annealing process, and there is a possibility that the bake hardening amount (BH) is deteriorated. On the other hand, when the annealing temperature (T) exceeds 880°C, although it is beneficial to workability, the shape of the steel plate is deteriorated due to the deviation of the grain size, and there is a possibility that problems occur in the annealing heating furnace equipment.

[0076] In addition, the annealing time (t), that is, the holding time at the annealing temperature is preferably 30 to 80 seconds, and more preferably 40 to 70 seconds. When the annealing time is sufficiently ensured after the γ-fiber texture is fully developed, part of the carbides will be re-melted into solid solution carbon. When cooling is performed in the state of the presence of such solid solution carbon, an appropriate level of solid solution carbon remains in the thin steel plate, thereby showing an excellent bake hardening amount (BH). If the annealing time (t) is less than 30 seconds, due to insufficient remelting time, solid solution carbon will not remain in the thin steel plate, or the remaining solid solution carbon is insufficient, so that the displayed bake hardening amount (BH) is poor. On the other hand, when the annealing time (t) exceeds 80 seconds, the grains are coarsened due to the excessively long holding time, and the shape of the steel plate will deteriorate due to the deviation of the grain size, which is also disadvantageous in terms of economic efficiency.

[0077] According to one example, the annealing temperature (T, ° C.) and annealing time (t, seconds) during continuous annealing may satisfy the following relational expression 2. If the value of 0.001*T*t is less than 30, the drawability and bake hardenability may deteriorate, while if the value of 0.001*T*t exceeds 70, the shape of the steel sheet may deteriorate due to coarsening of the grains and deviation of the grain size.

[0078] [Equation 2]30≤0.001*T*t≤70

[0079] In addition, during continuous annealing, the average heating rate from the recrystallization start temperature + 20°C to the annealing temperature is preferably 5°C / second or less, more preferably 4.5°C / second or less, and further preferably 3.8°C / second or less. The recrystallization start temperature is defined as the temperature at which the rolled structure elongated by cold rolling begins to form new recrystallized grains during annealing, and more specifically, it is defined as the temperature at which the area fraction of new recrystallized grains accounts for 50% of the total grains. In the initial stage of recrystallization, new grains will be accompanied by nucleation and growth. The slower the heating rate in this stage, the more the nucleation of the {111} texture that is beneficial to processability will increase, and ultimately a high r value can be ensured. When the heating rate in the temperature range exceeds 5°C / second, the nucleation of the {111} texture is insufficient during recrystallization, and the grains will also be refined, which may not fully ensure the processability required in the present invention. In addition, the slower the heating rate in the temperature range, the more favorable the nucleation and growth of the {111} texture which is advantageous for workability. Therefore, the lower limit of the heating rate is not particularly limited in the present invention.

[0080] Next, the continuously annealed cold-rolled steel sheet is cooled to a temperature of 650° C. or less.

[0081] At this time, the average cooling rate is preferably 2 to 10°C / second, more preferably 3 to 8°C / second. If the average cooling rate is less than 2°C / second, the dissolved carbon that has been re-melted during annealing will precipitate as carbides again, which may result in poor bake hardenability. On the other hand, when the average cooling rate exceeds 10°C / second, there is a possibility of plate shape distortion. In addition, 650°C is the temperature at which most of the precipitation and diffusion of carbides are completed, so there is no particular limitation on the cooling conditions thereafter.

[0082] Next, the cooled cold-rolled steel sheet is temper-rolled to obtain a high-strength thin steel sheet.

[0083] At this time, the flattening reduction rate is preferably 0.3-1.6%. Flattening increases the yield strength of steel, increases the aging resistance through the large number of movable dislocations introduced during the rolling process, and increases the bake hardenability through the interaction between the solid solution carbon and the dislocations. If the flattening reduction rate is less than 0.3%, it is not only not conducive to controlling the plate shape, but also cannot fully ensure the movable dislocation, so the possibility of tensile strain defects is high. On the other hand, when the flattening reduction rate exceeds 1.6%, when forming the parts of the customer company, not only the possibility of cracks increases, but also there is a tendency for the formability index r value to decrease.

[0084] Next, as needed, hot-dip galvanizing can be performed on the surface of the high-strength thin steel sheet to obtain a hot-dip galvanized steel sheet, or alloying heat treatment can be performed after hot-dip galvanizing to obtain an alloyed hot-dip galvanized steel sheet. At this time, the alloying heat treatment temperature is preferably 450 to 600°C. When the alloying heat treatment temperature is lower than 450°C, alloying is insufficient, which will cause a reduction in sacrificial corrosion protection or a reduction in plating adhesion. On the other hand, when the alloying heat treatment temperature exceeds 600°C, excessive alloying is performed, which will cause a reduction in powdering. In addition, after the alloying heat treatment, the Fe concentration in the coating is preferably 8 to 12 weight %. DETAILED DESCRIPTION

[0085] The present invention will be described in more detail below by way of examples. However, the following examples are merely illustrative for describing the present invention in more detail and do not limit the scope of the present invention.

[0086] A steel billet (220 mm thick) having the alloy composition shown in Table 1 below was heated to 1200°C and hot rolled to produce a hot rolled steel sheet (3.2 mm thick). At this time, the finishing rolling temperature was uniformly set to about 930°C above Ar3. The hot rolled steel sheet was then coiled, cold rolled, continuously annealed, cooled, and temper rolled under the conditions shown in Table 2 below to produce a thin steel sheet.

[0087] Afterwards, the number, distribution, texture, etc. of the precipitates of each manufactured thin steel plate were observed and measured, and the results are shown in Table 3 below. More specifically, the number ratio of carbides and the number of FeTiP precipitates are the number of precipitates per unit length (μm) at 5 positions after observing the precipitates through a replica using a transmission electron microscope (TEM) and calculating the average value. The texture is based on the ND direction crystal orientation degree under the conditions of R (rolling), T (transverse), and V (vertical) at the 1 / 4t position of the steel plate, and the intensity ratio of each orientation is calculated and analyzed (using ODF) using electron backscatter diffraction (EBSD). In addition, Figure 1 This is a graph analyzing the degree of development of the texture of Inventive Example 1. All the Inventive Examples show a tendency similar to that of Inventive Example 1.

[0088] Afterwards, the r value and bake hardening amount (BH) of each manufactured thin steel plate were measured. The test piece was taken according to JIS No. 5 specification, and the r value was measured using the ASTM STD test piece. The bake hardening amount was evaluated by the difference between the yield strength value after 2% pre-strain and the yield strength value after the test piece was kept at 170°C for 20 minutes.

[0089] [Table 1]

[0090]

[0091]

[0092] [Table 2]

[0093]

[0094]

[0095] [Table 3]

[0096]

[0097] Referring to Table 3, it can be seen that for Invention Examples 1 to 7 that meet the alloy composition and manufacturing conditions proposed by the present invention, the number of FeTiP precipitates per unit area, the ratio of carbides with a size of less than 20 nm present in the ferrite grains, and the average random strength ratio (b / a) all meet the range controlled by the present invention, and basically can ensure not only an r value of more than 1.9, but also a (yield strength * r value) of more than 290 MPa, and a BH property of more than 4 MPa.

[0098] However, although the alloy compositions of Comparative Examples 1 to 7 satisfy the ranges proposed in the present invention, at least one of the manufacturing conditions fails to satisfy the ranges proposed in the present invention, and thus exhibits poor drawability and bake hardenability. In addition, since the alloy compositions of Comparative Examples 8 to 11 fail to satisfy the ranges proposed in the present invention, poor drawability and bake hardenability are exhibited.

Claims

1. A high-strength steel sheet, which is composed, by weight, of C: 0.001-0.004%, Si: 0.5% or less and 0% or less, Mn: 1.2% or less and 0% or less, P: 0.005-0.12%, S: 0.01% or less, N: 0.01% or less, acid-soluble aluminum: 0.1% or less and 0% or less, Ti: 0.015-0.04%, the remainder of Fe and unavoidable impurities, wherein the contents of Ti, N and S satisfy the following relational expression 1, and the steel sheet contains 0.2 / μm 2 The following FeTiP precipitates are present, and the ratio (b / a) of the average random intensity ratio (b) of the (111) [1-10] to (111) [-1-12] orientation to the average random intensity ratio (a) of the (001) [1-10] to (110) [1-10] orientation at the t / 4 position in the plate thickness direction is 2.3 or more, the bake hardening amount (BH) is 4 MPa or more, and the average plastic anisotropy coefficient r value is 2.06 or more, in, t is the thickness of the thin steel plate, [Relationship 1] -0.02≤[Ti]-(24 / 7)[N]-(3 / 2)[S]≤0.025, Herein, [Ti], [N] and [S] respectively represent the weight % content of the element.

2. The high-strength steel sheet according to claim 1, wherein: The thin steel sheet further includes, in weight%, one or more selected from the group consisting of Nb: 0.005 to 0.04% and B: 0.002% or less and excluding 0%.

3. The high-strength steel sheet according to claim 1, wherein: The steel sheet has a P defined by the following mathematical formula 1. in More than 80%, [Mathematical formula 1] P in (%) = {N in / (N in +N gb )}×100, Among them, N in N is the number of carbides with a circle equivalent diameter of 20 nm or less existing in the grains, gb It is the number of carbides with a circle equivalent diameter of 20 nm or less existing in the grain boundary.

4. The high-strength steel sheet according to claim 1, wherein: The product of the yield strength (YS) and the average plastic anisotropy coefficient (r value) of the thin steel plate is 290 MPa or more.

5. A method for manufacturing a high-strength thin steel plate, characterized in that: The method is used to manufacture the high-strength thin steel plate according to claim 1, and the method comprises the following steps: The steel slab is hot rolled at a finishing temperature of Ar3°C or higher to obtain a hot rolled steel sheet, wherein the steel slab is composed of, by weight%, C: 0.001-0.004%, Si: 0.5% or less and 0% or less, Mn: 1.2% or less and 0% or less, P: 0.005-0.12%, S: 0.01% or less, N: 0.01% or less, acid soluble aluminum: 0.1% or less and 0% or less, Ti: 0.015-0.04%, and the balance of Fe and inevitable impurities; Cooling the hot rolled steel sheet from the finishing temperature to a temperature of 450 to 750° C. at an average cooling rate of 10 to 200° C. / s and coiling the hot rolled steel sheet; Cold rolling the coiled hot-rolled steel sheet at a reduction rate of 75% or more to obtain a cold-rolled steel sheet; The cold-rolled steel sheet is heated to an annealing temperature of 830 to 880° C., and then maintained at the annealing temperature for 30 to 80 seconds, thereby performing continuous annealing; Cooling the continuously annealed cold-rolled steel sheet to a temperature below 650° C. at a rate of 2 to 10° C. / second; Temper rolling the cooled cold-rolled steel sheet at a reduction rate of 0.3 to 1.6%; The contents of Ti, N and S in the steel billet satisfy the following relation 1: When the temperature of the cold-rolled steel sheet is increased, the average temperature increase rate from (recrystallization start temperature + 20)°C to the annealing temperature is 5°C / sec or less. The annealing temperature (T, ° C.) and annealing time (t, seconds) during the continuous annealing satisfy the following relationship 2: [Relationship 1] -0.02≤[Ti]-(24 / 7)[N]-(3 / 2)[S]≤0.025, Wherein, [Ti], [N] and [S] represent the weight % content of the element respectively. [Relationship 2]43≤0.001*T*t≤70.

6. The method for manufacturing a high-strength thin steel plate according to claim 5, wherein: The steel slab further includes, in terms of weight%, one or more selected from the group consisting of Nb: 0.005 to 0.04% and B: 0.002% or less and excluding 0%.

7. The method for manufacturing a high-strength thin steel plate according to claim 5, wherein: The method further includes the step of hot-dip galvanizing on the surface of the cold-rolled steel sheet that has been temper-rolled.

8. The method for manufacturing a high-strength thin steel plate according to claim 7, wherein: The method further comprises the step of performing an alloying heat treatment at 450-600° C. after the hot-dip galvanizing.

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