High-strength steel sheet excellent in hole expansion and ductility and method for manufacturing the same

Abstract

The present invention relates to a steel suitable for automotive materials, and more particularly to a high-strength steel sheet having excellent hole expansion and ductility and a method of manufacturing the same. The microstructure of the high-strength steel sheet of the present invention is composed of a hard phase and a soft phase, and by an optimized cold-rolling and annealing process, a martensite phase as the hard phase is uniformly distributed on a recrystallized ferrite matrix, and the crack resistance during processing can be improved by introducing a non-equilibrium ferrite phase at the interface of the hard phase and the soft phase.

Description

Technical Field

[0001] This invention relates to a steel suitable for automotive materials, and more specifically to a high-strength steel sheet with excellent hole expansion and ductility, and a method for manufacturing the same. Background Technology

[0002] In recent years, in the automotive industry, due to environmental regulations and energy use regulations related to CO2 emissions, high-strength steel is needed to improve fuel efficiency or durability.

[0003] In particular, with the expansion of regulations on the impact stability of automobiles, high-strength steel with excellent strength is being used as the material for structural components such as members, seat rails, and pillars, which are used to improve the impact resistance of the vehicle body.

[0004] These automotive parts have complex shapes due to their stability and design, and are primarily formed and manufactured using stamping dies, thus requiring high strength and a high level of formability.

[0005] However, while steel exhibits advantages in absorbing impact energy as its strength increases, elongation typically decreases with increasing strength, leading to reduced processability. Furthermore, excessively high yield strength reduces material introduction into the mold during forming, resulting in poorer formability and increased manufacturing costs.

[0006] Furthermore, automotive parts often have numerous expanded forming areas after machining holes, thus requiring hole expandability (HER) for successful forming. However, high-strength steel has low hole expandability, leading to defects such as cracks during the forming process. As mentioned above, poor hole expandability can cause cracks in the formed parts of the component during a collision, making the component more susceptible to damage and potentially jeopardizing passenger safety.

[0007] In addition, representative high-strength steels used in automobiles include dual-phase steel (DP steel), transformation-induced plasticity steel (TRIP steel), complex-phase steel (CP steel), and ferrite-bainite steel (FB steel).

[0008] DP steel, as an ultra-high tensile steel, has a low yield strength ratio of approximately 0.5 to 0.6, making it easy to process, and boasts a high elongation rate second only to TRIP steel. Therefore, it is mainly used in exterior doors, seat rails, seat belts, suspensions, arms, and wheel hubs.

[0009] TRIP steel has a yield strength ratio ranging from 0.57 to 0.67, thus exhibiting excellent formability (high ductility), making it suitable for components requiring high formability such as frame parts, roofs, seat belts, and bumpers.

[0010] CP steel has a low yield strength ratio, high elongation and bending workability, so it is used in side panels, body bottom reinforcements, etc., while FB steel has excellent hole expansion properties, so it is mainly used in suspension lower arms or wheel discs.

[0011] DP steel is primarily composed of ferrite with excellent ductility and hard phases (martensite and bainite), and may contain trace amounts of retained austenite. This type of DP steel exhibits low yield strength and high tensile strength, resulting in excellent properties such as a low yield ratio (YR), high work hardening rate, high ductility, continuous yield behavior, room temperature aging resistance, and bake hardening properties. Furthermore, by controlling the fraction of each phase, recrystallization degree, and distribution uniformity, high-strength steel with high porosity can be manufactured.

[0012] However, in order to ensure ultra-high strength with a tensile strength of over 1100 MPa, it is necessary to increase the fraction of hard phases such as martensitic phases that are beneficial to improving strength. In this case, the yield strength increases, which leads to defects such as cracks in the stamping process.

[0013] Typically, automotive DP steel is produced by manufacturing slabs through steelmaking and continuous casting processes, followed by heating, rough rolling, and hot finishing rolling to obtain hot-rolled coils, which are then annealed to produce the final product.

[0014] Annealing is a process mainly carried out when manufacturing cold-rolled steel sheets. Cold-rolled steel sheets are manufactured as follows: hot-rolled coils are pickled to remove surface oxide scale, cold-rolled at room temperature with a specified reduction rate, and then annealed and further leveling rolling processes are performed as needed.

[0015] Cold-rolled steel sheets (cold-rolled materials) obtained by cold rolling are in a very hardened state and are not suitable for making parts that require machinability. Therefore, their machinability can be improved by softening them through heat treatment in a continuous annealing furnace as a subsequent process.

[0016] As an example, the annealing process involves heating a steel sheet (cold-rolled material) to approximately 650-850°C in a furnace and holding it for a certain period of time, thereby reducing hardness and improving machinability through recrystallization and phase transformation.

[0017] For unannealed steel sheets, their hardness, especially surface hardness, is high and their machinability is insufficient. However, annealed steel sheets have recrystallized structures, which reduce hardness, yield point, and tensile strength, thus helping to improve machinability.

[0018] As a representative method for reducing the yield strength of DP steel, the heating process during continuous annealing is to make equiaxed crystals by completely recrystallizing ferrite, so that the austenite formed and grown in subsequent processes will also be in equiaxed crystal form, which is beneficial for forming austenite phase with small and uniform particle size.

[0019] Furthermore, as a prior art for improving the workability of high-strength steel, Patent Document 1 proposes a method based on microstructure refinement. Specifically, it discloses a method for dispersing finely precipitated copper particles with a particle size of 1-100 nm within the microstructure of a multiphase steel plate dominated by a martensitic phase. However, this technique requires the addition of 2-5% Cu to obtain well-formed finely precipitated phase particles, which may lead to red-hot brittleness caused by a large amount of Cu, and also results in an excessive increase in manufacturing costs.

[0020] Patent document 2 discloses a steel sheet with a ferrite matrix and a pearlite microstructure comprising 2-10% by area, wherein the strength is improved by precipitation strengthening and grain refinement caused by the addition of carbide-nitride forming elements (e.g., Ti). This steel sheet exhibits good porosity but has limitations in further improving tensile strength, and suffers from cracking during stamping due to its high yield strength and low ductility.

[0021] Patent document 3 discloses a method for manufacturing cold-rolled steel sheets that simultaneously achieve high strength and high ductility by utilizing tempered martensitic phases and have excellent sheet shape after continuous annealing. However, the carbon (C) content in the steel is as high as 0.2% or more. Therefore, in addition to the problem of poor weldability, there is also the possibility of furnace pit defects caused by the addition of a large amount of Si.

[0022] In view of the above-mentioned existing technology, in order to improve the formability of high-strength steel, such as its pore-expanding properties, which meet the physical properties of weldability and other uses, it is necessary to develop a method that can reduce yield strength but improve ductility while forming a uniform microstructure in the steel.

[0023] [Existing technical documents]

[0024] [Patent Literature]

[0025] (Patent Document 1) Japanese Patent Publication No. 2005-264176

[0026] (Patent Document 2) Korean Patent Publication No. 2015-0073844

[0027] (Patent Document 3) Japanese Patent Publication No. 2010-090432 Summary of the Invention

[0028] Technical problems to be solved

[0029] One aspect of the present invention provides a high-strength steel sheet having a low yield strength ratio and high strength, and excellent formability such as hole expansion properties, which is suitable for use as a material for automotive structural components, etc., and a method thereof.

[0030] The technical problem to be solved by this invention is not limited to the above-described content. The technical problem to be solved by this invention can be understood from the overall content of this specification, and those skilled in the art to which this invention pertains can easily understand the additional technical problems of this invention.

[0031] Technical solution

[0032] One aspect of the present invention provides a high-strength steel plate with excellent hole-expanding properties, wherein, by weight percent, the high-strength steel plate comprises: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): less than 1.2% (excluding 0%), chromium (Cr): less than 0.1% (excluding 0%), molybdenum (Mo): less than 0.1% (excluding 0%), niobium (Nb): less than 0.1% (excluding 0%), and titanium (Ti): less than 0.1% (excluding 0%). Except for %, boron (B): less than 0.002% (except 0%), aluminum (sol.Al): 0.02-0.05%, phosphorus (P): less than 0.05% (except 0%), sulfur (S): less than 0.01% (except 0%), nitrogen (N): less than 0.01% (except 0%), iron (Fe) and other unavoidable impurities, the fine structure contains 20-30% ferrite, 5-15% non-equilibrium ferrite and the balance martensite.

[0033] Another aspect of the present invention provides a method for manufacturing a high-strength steel plate with excellent porosity, characterized in that the method comprises the following steps: preparing a steel billet containing the above-mentioned alloy composition; heating the steel billet in a temperature range of 1100-1300°C; hot-rolling the heated steel billet to manufacture a hot-rolled steel plate; coiling the hot-rolled steel plate in a temperature range of 400-700°C; cooling the coiled hot-rolled steel plate to room temperature; cold-rolling the cooled hot-rolled steel plate to manufacture a cold-rolled steel plate; subjecting the cold-rolled steel plate to continuous annealing; after the continuous annealing, performing a first cooling at an average cooling rate of 1-10°C / second to cool to a temperature range of 570-630°C; and after the first cooling, performing a second cooling at an average cooling rate of 5-50°C to cool to a temperature range of 300-400°C, wherein the continuous annealing is performed in an apparatus equipped with a heating zone, a heat soaking zone, and a cooling zone, and the heating zone and the heat soaking zone are controlled at a temperature range of 810-850°C.

[0034] Beneficial effects

[0035] According to the present invention, a steel plate can be provided that, even with high strength, has excellent hole-expanding properties, thereby improving formability and impact resistance.

[0036] As described above, the improved formability of the steel sheet of the present invention prevents processing defects such as cracks or wrinkles during stamping, thus making it suitable for use in components such as structures that require processing into complex shapes. Furthermore, it effectively manufactures a material with improved impact resistance, preventing defects such as cracks from easily forming when a car using this component inevitably experiences a collision. Attached Figure Description

[0037] Figure 1 The thermal history and phase transformation history during continuous annealing according to an embodiment of the present invention are shown.

[0038] Figure 2 (a) illustrates a void formation mechanism in an organization, and (b) illustrates an interface strengthening mechanism in an organization according to an embodiment of the invention.

[0039] Figure 3 Microscopic images of an inventive example and a comparative example according to an embodiment of the present invention are shown.

[0040] Best practice

[0041] The inventors of this invention have conducted in-depth research in order to develop a material with formability that is suitable for use in automotive materials for parts that need to be processed into complex shapes.

[0042] In particular, the inventors derived a microstructure that can eliminate the hardness difference between the soft and hard phases that affect the crack resistance of steel, and confirmed that the objective can be achieved by controlling the micronization of the hard phase and the grain shape, which are conducive to preventing the formation and propagation of pores, thus completing the present invention.

[0043] In particular, the present invention introduces an intermediate phase, preferably a non-equilibrium ferrite phase, to eliminate the hardness difference between the soft and hard phases, and the technical significance of forming such a structure lies in optimizing the alloy composition and manufacturing conditions.

[0044] The present invention will now be described in detail.

[0045] According to one aspect of the present invention, a high-strength steel plate with excellent porosity and ductility may contain, by weight percent: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): less than 1.2% (except 0%), chromium (Cr): less than 0.1% (except 0%), molybdenum (Mo): less than 0.1% (except 0%), niobium (Nb): less than 0.1% (except 0%), titanium (Ti): less than 0.1% (except 0%), boron (B): less than 0.002% (except 0%), aluminum (sol.Al): 0.02-0.05%, phosphorus (P): less than 0.05% (except 0%), sulfur (S): less than 0.01% (except 0%), and nitrogen (N): less than 0.01% (except 0%).

[0046] The reasons for limiting the alloy composition of the steel plate provided in this invention as described above will be explained in detail below.

[0047] In addition, unless otherwise specified, the content of each element in this invention is based on weight, and the proportion of the tissue is based on area.

[0048] Carbon (C): 0.05-0.12%

[0049] Carbon (C) is an important element added for solid solution strengthening. This C combines with precipitating elements to form fine precipitates, thus helping to improve the strength of steel.

[0050] When the carbon content exceeds 0.12%, the increased hardenability leads to the formation of martensite during cooling in the steel manufacturing process, resulting in excessively increased strength and a decrease in elongation. Furthermore, the reduced weldability increases the possibility of welding defects when machining components. Conversely, when the carbon content is less than 0.05%, it is difficult to ensure the target strength level.

[0051] Therefore, the content of C can be 0.05-0.12%. More preferably, the content of C can be 0.06% or more, and can be less than 0.10%.

[0052] Manganese (Mn): 2.5-3.0%

[0053] Manganese (Mn) is an element that causes sulfur (S) in steel to precipitate as MnS, thereby preventing hot brittleness caused by the formation of FeS, and is beneficial to the solid solution strengthening of steel.

[0054] When the Mn content is less than 2.5%, not only is the aforementioned effect unattainable, but it is also difficult to ensure the target strength level. On the other hand, when the Mn content exceeds 3.0%, the likelihood of problems such as weldability and hot rollability increases, and martensite formation is more likely due to increased hardenability, which may reduce ductility. Furthermore, the formation of excessive Mn bands (Mn oxide bands) within the microstructure increases the risk of defects such as processing cracks. Additionally, Mn oxides dissolve from the surface during annealing, significantly hindering plating performance.

[0055] Therefore, the content of Mn can be 2.5-3.0%.

[0056] Silicon (Si): Less than 1.2% (except 0%)

[0057] Silicon (Si) is a ferrite stabilizing element that helps ensure a target ferrite fraction by promoting ferrite phase transformation. Furthermore, due to its excellent solid solution strengthening ability, it is effective in improving the strength of ferrite and is a useful element that ensures strength without reducing the ductility of steel.

[0058] When the Si content exceeds 1.2%, the solid solution strengthening effect becomes excessive, which can reduce ductility and cause surface oxide scale defects, thus adversely affecting the quality of the plated surface. Furthermore, it hinders chemical processing.

[0059] Therefore, the Si content can be less than 1.2%, and may exclude 0%. More preferably, the Si content can be 0.1% or more.

[0060] Chromium (Cr): less than 0.1% (except 0%)

[0061] Chromium (Cr) is an element that contributes to the desired microstructure of this invention. It inhibits the formation of martensite and bainite phases during annealing heat treatment, and on the other hand, forms fine carbides, thereby contributing to increased strength. Specifically, Cr has the effect of suppressing the formation of bainite that competes with non-equilibrium ferrite; therefore, when an appropriate level of Cr is present, the formation of a non-equilibrium ferrite phase is favored at high temperatures.

[0062] When the Cr content exceeds 0.1%, a non-equilibrium ferrite phase cannot be formed, thus reducing the ductility and porosity of the steel. Furthermore, when carbides form at grain boundaries, strength and elongation may deteriorate. Additionally, there is an increase in manufacturing costs.

[0063] Therefore, the Cr content can be less than 0.1%, and 0% can be excluded. More preferably, the Cr content can be more than 0.01%.

[0064] Molybdenum (Mo): Less than 0.1% (except 0%)

[0065] Molybdenum (Mo) is an element that helps improve strength by inhibiting the phase transformation of pearlite, promoting the formation of non-equilibrium ferrite phase, suppressing the formation of martensite phase during annealing heat treatment, and forming fine carbides.

[0066] When the Mo content exceeds 0.1%, the hardenability becomes excessive, making it impossible to form a non-equilibrium ferrite phase. As a result, the ductility and porosity of the steel may decrease, and there is an issue of increased manufacturing costs.

[0067] Therefore, the content of Mo can be less than 0.1%, and 0% can be excluded. More preferably, the content of Mo can be less than 0.01%.

[0068] Niobium (Nb): Less than 0.1% (except 0%)

[0069] Niobium (Nb) is an element that helps improve strength by inhibiting the coarsening of austenite grains during annealing heat treatment due to austenite grain boundary segregation and the formation of fine carbides.

[0070] When the Nb content exceeds 0.1%, coarse carbides precipitate. Due to the reduced carbon content in the steel, the strength and elongation may be poor, and there is an increase in manufacturing costs.

[0071] Therefore, the Nb content can be less than 0.1%, and 0% can be excluded. More preferably, the Nb content can be less than 0.01%.

[0072] Titanium (Ti): Less than 0.1% (except 0%)

[0073] Titanium (Ti) is an element that forms fine carbides, helping to ensure yield strength and tensile strength. Furthermore, Ti causes nitrogen in steel to precipitate as TiN, thereby suppressing the formation of AlN caused by the unavoidable presence of Al in the steel, thus reducing the likelihood of cracking during continuous casting.

[0074] When the Ti content exceeds 0.1%, coarse carbides precipitate, and due to the reduced carbon content in the steel, there is a possibility of decreased strength and elongation. Furthermore, there is a possibility of nozzle clogging during continuous casting, and there is an increase in manufacturing costs.

[0075] Therefore, the Ti content can be less than 0.1%, and 0% can be excluded. More preferably, the Ti content can be less than 0.01%.

[0076] Boron (B): less than 0.002% (except 0%)

[0077] Boron (B) is an element that delays the transformation of austenite into pearlite during the cooling process after annealing heat treatment. However, when the content of B exceeds 0.002%, excessive B accumulates on the surface, which may lead to poor coating adhesion.

[0078] Therefore, the content of B can be less than 0.002%, and 0% can be excluded.

[0079] Aluminum (sol.Al): 0.02-0.05%

[0080] Aluminum (sol.Al) is added to refine the grain size of steel and for deoxidation. When the aluminum (sol.Al) content is less than 0.02%, aluminum-killed steel cannot be manufactured in a stable state. On the other hand, when the aluminum (sol.Al) content exceeds 0.05%, grain refinement occurs, thus improving strength. However, this also leads to the formation of excessive inclusions during continuous casting operations, increasing the likelihood of surface defects in the coated steel sheet.

[0081] Therefore, the content of the acid-soluble aluminum (sol.Al) can be 0.02-0.05%.

[0082] Phosphorus (P): less than 0.05% (except 0%)

[0083] Phosphorus (P) is the substitutional element with the greatest solid solution strengthening effect, and it is also an element that improves in-plane anisotropy without significantly reducing formability while ensuring strength. However, when too much P is added, the possibility of brittle fracture increases significantly, leading to an increased likelihood of slab breakage during hot rolling, and it also hinders the properties of the coated surface.

[0084] Therefore, in this invention, the content of P can be controlled below 0.05%, except for 0% considering the level that is inevitably added.

[0085] Sulfur (S): less than 0.01% (except 0%)

[0086] Sulfur (S) is an impurity element in steel and is an unavoidable addition. Since it hinders ductility, it is preferable to control the sulfur content as low as possible. In particular, S can increase the likelihood of developing red-hot brittleness, so it is preferable to control the sulfur content to below 0.01%. However, considering the level that is unavoidably added during manufacturing, 0% is an exception.

[0087] Nitrogen (N): less than 0.01% (except 0%)

[0088] Nitrogen (N) is a solid solution strengthening element, but when the nitrogen (N) content exceeds 0.01%, the risk of brittleness increases, and it combines with Al in the steel to precipitate excessive AlN, which may hinder the quality of continuous casting.

[0089] Therefore, the content of N can be less than 0.01%, and considering the level that will inevitably be added, 0% can be excluded.

[0090] The remaining component of this invention is iron (Fe). However, undesirable impurities may inevitably be introduced from the raw materials or the surrounding environment during conventional manufacturing processes, and therefore such impurities cannot be excluded. These impurities are well known to those skilled in the art of conventional manufacturing processes, and therefore their contents are not specifically described herein.

[0091] The fine microstructure of the steel plate of the present invention having the above alloy composition can be composed of ferrite as a soft phase, martensite as a hard phase, and non-equilibrium ferrite phase formed at their interfaces.

[0092] Specifically, the steel plate of the present invention comprises 20-30% ferrite phase and 5-15% non-equilibrium ferrite phase by area, and may contain martensite phase as the balance. Additionally, it may contain trace amounts of retained austenite phase.

[0093] In this invention, the non-equilibrium ferrite phase is a microstructure that minimizes the hardness difference between the soft and hard phases, and is different from the existing equilibrium ferrite (polygonal ferrite). The non-equilibrium ferrite can be acicular ferrite or bainitic ferrite. Furthermore, depending on the cooling conditions, it may include Widmanstätten ferrite, massive ferrite, etc. Specifically, the non-equilibrium ferrite is influenced by the composition of the mother phase and contains relatively higher levels of C and Mn compared to equilibrium ferrite. For example, in the case of equilibrium ferrite, assuming a C concentration of 0.02%, the non-equilibrium ferrite has a higher C content of 0.03-0.04%.

[0094] Therefore, the C and Mn concentrations of the hard phase formed near (around) the non-equilibrium ferrite are relatively reduced, thus decreasing the hardness difference between the soft and hard phases and improving porosity. Furthermore, when the Si concentration in the non-equilibrium is less than 1%, the stacking fault energy increases, cross slip becomes difficult, and void formation due to deformation is hindered. Figure 2 ).

[0095] When the fraction of this non-equilibrium ferrite phase is too high, the fraction of the hard phase is relatively reduced, thus failing to ensure the target strength level. With this in mind, the non-equilibrium ferrite phase can be contained in amounts of less than 15%. On the other hand, when the fraction of the non-equilibrium ferrite phase is less than 5%, the aforementioned effect (minimizing the hardness difference between the hard and soft phases) cannot be sufficiently obtained, thus resulting in poor porosity.

[0096] When the fraction of the ferrite phase is less than 20%, it is not conducive to ensuring the ductility of the steel. However, when the fraction of the ferrite phase exceeds 30%, the fraction of the hard phase decreases relatively, making it difficult to ensure the target level of strength.

[0097] In the microstructure other than the ferrite and non-equilibrium ferrite phases, the fraction of martensite phase is not specifically limited, but to ensure ultra-high strength of 1100 MPa or more, a martensite phase with an area fraction of 50% or more may be included. However, when the fraction of martensite phase exceeds 75%, ductility decreases, making it difficult to ensure the target level of formability.

[0098] Furthermore, it is advantageous for the residual austenite phase to have a fraction of no more than 3%, and even if the fraction is 0%, there is no problem in ensuring the desired physical properties.

[0099] The steel plate of the present invention having the above-mentioned fine structure has a tensile strength of 1100 MPa or more, a yield strength of 550-700 MPa, and an elongation (total elongation) of 12% or more, thus possessing the characteristics of high strength and high ductility.

[0100] Furthermore, the steel plate has a hole expansion ratio (HER) of more than 25%, which results in excellent resistance to cracks that may occur during processing and resistance to impact fracture.

[0101] The following describes in detail a method for manufacturing a high-strength steel sheet with excellent hole expansion and ductility according to another aspect of the present invention.

[0102] In short, the present invention can manufacture the desired steel plate through the process of [steel billet heating-hot rolling-coiling-cold rolling-continuous annealing], and the process is described in detail below.

[0103] Heating of steel billets

[0104] First, a steel billet that meets the above alloy composition can be prepared, and then it can be heated.

[0105] This process is performed to ensure the smooth execution of subsequent hot rolling processes and to fully obtain the desired physical properties of the steel sheet. In this invention, the conditions for this heating process are not particularly limited; any normal conditions are acceptable. As an example, the heating process can be carried out in a temperature range of 1100-1300℃.

[0106] [Hot Rolled]

[0107] The steel billet heated as described above can be hot-rolled to produce hot-rolled steel plates. In this case, hot finishing rolling can be performed at an exit-side temperature of Ar3 or higher up to 1000°C or lower.

[0108] When the exit temperature during hot finishing rolling is below Ar3, the resistance to hot deformation increases sharply, and the top, bottom, and edge portions of the hot-rolled coil become single-phase regions, thus increasing in-plane anisotropy and potentially leading to poorer formability. Furthermore, when the exit temperature during hot finishing rolling exceeds 1000°C, the rolling load is relatively reduced, which is beneficial for productivity, but there is a possibility of forming a thick oxide scale.

[0109] More specifically, the hot finishing rolling can be carried out in a temperature range of 760-940°C.

[0110] [Collection]

[0111] The hot-rolled steel sheet manufactured as described above can be rolled into a coil shape.

[0112] The coiling process can be carried out within a temperature range of 400-700℃. When the coiling temperature is below 400℃, excessive martensitic or non-equilibrium phases are formed, causing an excessive increase in the strength of the hot-rolled steel sheet. Therefore, problems such as shape defects due to load may occur during subsequent cold rolling. On the other hand, when the coiling temperature exceeds 700℃, the pickling performance deteriorates due to the increase in surface oxide scale.

[0113] [cool down]

[0114] Preferably, the coiled hot-rolled steel sheet is cooled to room temperature at an average cooling rate of 0.1°C / s or less (except 0°C / s). In this case, the coiled hot-rolled steel sheet can be cooled after processes such as conveying and placing; the processes prior to cooling are not limited to this.

[0115] As described above, by cooling the coiled hot-rolled steel sheet at a specified speed, a hot-rolled steel sheet with finely dispersed carbides serving as austenite nucleation sites can be obtained.

[0116] [Cold Rolled]

[0117] The hot-rolled steel sheet that has been coiled as described above can be cold-rolled to produce cold-rolled steel sheet.

[0118] In this invention, the cold rolling can be performed with a cold rolling reduction rate of 55-70%. When the cold rolling reduction rate is less than 55%, the recrystallization driving force weakens, making it difficult to obtain good recrystallized grains. On the other hand, when the cold rolling reduction rate exceeds 70%, the risk of cracks at the edge of the steel sheet increases, and there is a possibility of a sharp increase in rolling load.

[0119] In this invention, by applying an appropriate level of cold rolling reduction during cold rolling, the recrystallization of ferrite can be further promoted in the heating zone during subsequent continuous annealing, thereby inducing the formation of fine ferrite and resulting in the formation of small and uniform austenite at the ferrite grain boundaries. This affects the size or distribution of the non-equilibrium structure during cooling and is beneficial for maintaining the strength of the final product while simultaneously improving elongation, porosity, and other processability.

[0120] Furthermore, the cold rolling reduction rate can be achieved in just one cold rolling cycle, that is, in just one stand. As mentioned above, the steel can be reduced, thus having an economical effect.

[0121] However, for hot-rolled steel plates with a thickness of 6 mm or more before cold rolling, the target reduction rate can be achieved by repeatedly rolling using a reversing mill. In this case, the total number of passes for repeated rolling can be set to one stand. A reversing mill is a mill used for rolling thin sheet steel, which refers to a mill that reciprocates and rolls the material between a pair of rolls, and each reciprocating pass of the material can be set as one pass.

[0122] The present invention allows for pickling of hot-rolled steel sheets prior to cold rolling, and the pickling process can be carried out using conventional methods.

[0123] [Continuous annealing]

[0124] Preferably, the cold-rolled steel sheet manufactured as described above is subjected to continuous annealing. As an example, the continuous annealing process can be carried out in a continuous annealing furnace (CAL).

[0125] Typically, a continuous annealing furnace (CAL) can consist of [heating zone - soaking zone - cooling zone (slow cooling zone and rapid cooling zone) - (over-aging zone if necessary)] and undergoes the following process: cold-rolled steel sheets are loaded into the continuous annealing furnace as described above, then heated at a specific temperature in the heating zone, and held in the soaking zone for a specified time after reaching the target temperature.

[0126] In this invention, the temperatures of the heating zone and the soaking zone during continuous annealing can be controlled to be the same, which means that the ending temperature of the heating zone and the starting temperature of the soaking zone are controlled to be the same. Figure 1 ).

[0127] Specifically, the temperature of the heating zone and the heat exchange zone can be controlled between 810-850℃.

[0128] When the temperature of the heating band is below 810°C, sufficient heat input for recrystallization cannot be applied. On the other hand, when the temperature of the heating band exceeds 850°C, productivity decreases and excessive austenite phase is formed, resulting in a significant increase in the fraction of hard phase after subsequent cooling, thus potentially leading to a decrease in the ductility of the steel.

[0129] Furthermore, when the temperature of the heat exchanger is below 810°C, excessive cooling is required at the termination temperature of the heating zone, which is uneconomical, and the heat used for recrystallization may be insufficient. On the other hand, when the temperature of the heat exchanger exceeds 850°C, the austenite fraction is too large, and the hard phase increases during cooling, which may reduce formability.

[0130] Increasing the temperature of the heat exchanger within the aforementioned temperature range can reduce the stability of austenite, thereby promoting the formation of non-equilibrium ferrite phases during subsequent cooling.

[0131] Although detailed explanations follow, this invention employs staged cooling during the cooling process after passing through the heating and soaking zones. The non-equilibrium ferrite phase is introduced after the first cooling step, and the final microstructure can consist of a certain percentage of soft phase, hard phase, and non-equilibrium ferrite phase. Therefore, in the steel plate of this invention, not only are strength and ductility improved, but also machinability is enhanced due to the interfacial strengthening effect of the non-equilibrium ferrite phase.

[0132] Therefore, in order to obtain the fine microstructure desired by the present invention, it is preferable to control the heat input applied to the steel plate during the continuous annealing in the heating zone consisting of the heating zone and the soaking zone.

[0133] [Stage Cooling]

[0134] As described above, the desired microstructure can be formed by cooling the cold-rolled steel sheet that has undergone the heat treatment as described above, and stepwise cooling is preferred in this case.

[0135] In this invention, the staged cooling can consist of primary cooling and secondary cooling. Specifically, after the continuous annealing, primary cooling can be performed at an average cooling rate of 1-10°C / second to cool to a temperature range of 570-630°C, and then secondary cooling can be performed at an average cooling rate of 5-50°C / second to cool to a temperature range of 300-400°C.

[0136] At this point, the primary cooling proceeds more slowly than the secondary cooling, thus suppressing plate shape defects caused by the rapid temperature drop during the secondary cooling, which is the subsequent relatively rapid cooling zone.

[0137] When the termination temperature during the first cooling is below 570°C, the low temperature results in low carbon diffusion activity and an increased carbon concentration in the ferrite. Conversely, the low carbon concentration in the austenite leads to an excessively high fraction of hard phases, increasing the yield strength ratio and thus increasing the tendency for cracking during processing. Furthermore, the excessively rapid cooling rate between the soaking zone and the cooling zone (slow cooling zone) causes uneven plate shape. When the termination temperature exceeds 630°C, subsequent cooling (secondary cooling) requires excessively high cooling rates, making it difficult to introduce non-equilibrium ferrite phases.

[0138] Furthermore, when the average cooling rate during the first cooling exceeds 10°C / second, carbon diffusion cannot be sufficiently achieved. Additionally, considering productivity, the first cooling can be performed at an average cooling rate of 1°C / second or higher.

[0139] As described above, after primary cooling, rapid cooling (secondary cooling) can be performed at a certain or higher cooling rate. However, when the secondary cooling termination temperature is below 300°C, cooling deviations occur in the width and length directions of the steel plate, potentially leading to a deterioration in plate shape. On the other hand, when the secondary cooling termination temperature exceeds 400°C, the hard phase cannot be adequately secured, potentially resulting in reduced strength. Furthermore, the formation of bainite may lead to an increase in yield strength and a decrease in elongation.

[0140] Furthermore, when the average cooling rate during the secondary cooling is less than 5°C / second, the fraction of the hard phase may be too high, but when the average cooling rate during the secondary cooling exceeds 50°C / second, the hard phase may become insufficient.

[0141] Additionally, if necessary, an aging process can be performed after the phased cooling is completed.

[0142] The over-aging treatment is a process in which the material is held at the secondary cooling termination temperature for a certain period of time, and uniform heat treatment is performed along the width and length directions of the roll material, thereby improving the shape quality. Therefore, the over-aging treatment can be performed for 200-800 seconds.

[0143] The over-aging treatment can be performed after the secondary cooling is terminated. Therefore, the temperature of the over-aging treatment can be the same as the temperature at which the secondary cooling is terminated, or the over-aging treatment can be performed within the range of the temperature at which the secondary cooling is terminated.

[0144] The high-strength steel sheet of the present invention, manufactured as described above, has a microstructure composed of hard and soft phases. In particular, through optimized cold rolling and annealing processes, the recrystallization of ferrite is maximized, resulting in a microstructure in which martensite, as a hard phase, is uniformly distributed on the ultimately recrystallized ferrite matrix. Furthermore, by introducing a non-equilibrium ferrite phase at the interface between the hard and soft phases, it has the effect of improving crack resistance during processing.

[0145] Therefore, even though the steel plate of the present invention has a high strength of 1100 MPa or more, it can ensure excellent formability such as hole expansion by ensuring a low yield strength ratio and high ductility.

[0146] The present invention will now be described in more detail through embodiments. However, these embodiments are merely illustrative of implementation of the invention, and the invention is not limited to these embodiments. This is because the scope of the invention is determined by the contents of the claims and the contents reasonably inferred therefrom. Detailed Implementation

[0147] (Example)

[0148] Steel billets with the alloy compositions shown in Table 1 are manufactured, and then each billet is heated at 1200°C for 1 hour, followed by hot finishing at a finishing temperature of 880-920°C to produce hot-rolled steel sheets. Subsequently, each hot-rolled steel sheet is coiled at 650°C and cooled to room temperature at a cooling rate of 0.1°C / second. Then, the coiled hot-rolled steel sheets are cold-rolled and continuously annealed according to the conditions shown in Table 2, followed by staged cooling (primary cooling-secondary cooling), and then over-aging at 360°C for 520 seconds to produce the final steel sheet.

[0149] At this point, the first cooling in the staged cooling process is carried out at an average cooling rate of 3°C / second, and the second cooling is carried out at an average cooling rate of 20°C / second. In addition, cold rolling is carried out on a single stand.

[0150] The microstructure of each steel plate manufactured as described above was observed, and the physical property indices used in the processing technology, such as tensile and processing characteristics and hole expansion rate, were evaluated. The results are shown in Table 3 below.

[0151] At this point, the tensile test on each specimen is carried out by taking tensile specimens of size JIS5 in a direction perpendicular to the rolling direction and then performing the tensile test at a strain rate of 0.01 / second.

[0152] In addition, the hole enlargement property (HER, %) measurement test was conducted according to ISO 16630 standard. Specifically, when a circular hole is punched in the specimen and then enlarged using a conical punch, the enlargement is expressed as the ratio of the enlargement amount until a crack at the hole edge penetrates along the thickness direction to the initial hole size. In this case, the specimen size was 120 mm × 120 mm, the clearance was 12%, the punch diameter was 10 mm, the punch holding load was 20 tons, and the test speed was set to 12 mm / min.

[0153] Furthermore, the martensitic and non-equilibrium phases corresponding to the hard phases in the microstructure were observed by SEM at magnifications of 2000x and 5000x after etching with nitric acid. The size and fraction of each observed phase were then measured. Additionally, the fraction of each phase was measured using SEM and an image analyzer program after etching with nitric acid.

[0154] [Table 1]

[0155]

[0156] [Table 2]

[0157]

[0158] [Table 3]

[0159]

[0160] As shown in Tables 1 to 3, it can be seen that the steel alloy composition and manufacturing conditions meet the conditions proposed in this invention. In particular, the cold rolling and continuous annealing processes in Examples 1 to 6 both meet the conditions proposed in this invention. During the annealing process after cold rolling, sufficient recrystallization of ferrite is achieved, forming a fine hard phase. Moreover, non-equilibrium ferrite structures are connected at the interface, resulting in high strength while maintaining appropriate yield strength for sheet metal processing and excellent elongation. Furthermore, due to excellent hole expansion properties, it can be confirmed that the target level of formability can be ensured.

[0161] On the other hand, in Comparative Examples 1 to 6, where the continuous annealing temperature during the steel sheet manufacturing process was low, recrystallization did not occur sufficiently, and the appropriate fraction of austenite formed in the soaking temperature exhibited high stability, thus preventing the sufficient introduction of non-equilibrium ferrite during cooling. As a result, poor ductility and / or porosity were observed.

[0162] In addition, although Comparative Examples 7 to 10 were heated at appropriate temperatures during continuous annealing, their ductility and / or porosity were poor due to the high termination temperature during the first cooling and insufficient time for the introduction of non-equilibrium ferrite during the cooling process.

[0163] Furthermore, in Comparative Examples 11 to 14, which contain excessive amounts of Cr as a hardenability element, the yield strength is too high, which poses a risk of cracking during processing. Moreover, due to the low temperature of the soaking tub, the non-equilibrium phase cannot be introduced, resulting in poor ductility in some comparative examples.

[0164] Figure 3 Microscopic images of Comparative Examples 4 to 7 and Invention Example 1 are shown.

[0165] like Figure 3 As shown, in Invention Example 1, a uniform and fine non-equilibrium phase is introduced into a sufficient fraction of recrystallized ferrite matrix during a first cooling process, and a certain fraction of martensite phase is formed during a second cooling process.

[0166] On the other hand, in Comparative Examples 4 to 7, due to the conditions of leaving the soaking zone or the primary cooling termination temperature during continuous annealing, it can be confirmed that a small amount of non-equilibrium ferrite is introduced. Among them, in Comparative Example 4, where the soaking zone temperature did not reach 800°C and the primary cooling termination temperature was relatively high, and in Comparative Example 7, where the primary cooling termination temperature was relatively high, it can be seen that the non-equilibrium ferrite content was less than 1% and was almost unobservable.

Claims

1. A high-strength steel plate with excellent hole-expanding properties, comprising, by weight percent: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): less than 1.2% and excluding 0%, chromium (Cr): less than 0.1% and excluding 0%, molybdenum (Mo): less than 0.1% and excluding 0%, niobium (Nb): less than 0.1% and excluding 0%, titanium (Ti): less than 0.1% and excluding 0%, and boron (B): less than 0.002%. Except for 0%, aluminum (sol.Al): 0.02-0.05%, phosphorus (P): less than 0.05% and except for 0%, sulfur (S): less than 0.01% and except for 0%, nitrogen (N): less than 0.01% and except for 0%, iron (Fe) and other unavoidable impurities, the microstructure contains 20-30% ferrite, 5-15% non-equilibrium ferrite and the balance martensite, wherein the non-equilibrium ferrite is one or more of acicular ferrite and bainitic ferrite.

2. The high-strength steel plate with excellent hole-expanding properties according to claim 1, wherein, The steel plate further comprises an area fraction of less than 3% and includes 0% residual austenite phase.

3. The high-strength steel plate with excellent hole-expanding properties according to claim 1, wherein, The steel plate has a tensile strength of 1100 MPa or higher, a yield strength of 550-700 MPa, and a total elongation of 12% or higher.

4. The high-strength steel plate with excellent hole-expanding properties according to claim 1, wherein, The hole expansion ratio (HER) of the steel plate is 25% or more.

5. A method for manufacturing a high-strength steel plate with excellent hole-expanding properties, characterized in that, The method includes the following steps: Prepare a steel billet, which, by weight percent, comprises: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): less than 1.2% and excluding 0%, chromium (Cr): less than 0.1% and excluding 0%, molybdenum (Mo): less than 0.1% and excluding 0%, niobium (Nb): less than 0.1% and excluding 0%, titanium (Ti): less than 0.1% and excluding 0%, boron (B): less than 0.002% and excluding 0%, aluminum (sol.Al): 0.02-0.05%, phosphorus (P): less than 0.05% and excluding 0%, sulfur (S): less than 0.01% and excluding 0%, nitrogen (N): less than 0.01% and excluding 0%, iron (Fe), and other unavoidable impurities; The steel billet is heated within a temperature range of 1100-1300℃; The heated steel billet is hot-rolled to produce hot-rolled steel sheet; The hot-rolled steel sheet is coiled within a temperature range of 400-700℃; The hot-rolled steel sheet is cooled to room temperature after being coiled. The cooled hot-rolled steel sheet is cold-rolled to produce cold-rolled steel sheet; The cold-rolled steel sheet is subjected to continuous annealing treatment; Following the continuous annealing, a single cooling process is performed at an average cooling rate of 1-10°C / second, cooling to a temperature range of 570-630°C; and After the first cooling, a second cooling process is performed at an average cooling rate of 5-50°C / second, cooling the temperature to a range of 300-400°C. The continuous annealing is carried out in a device equipped with a heating zone, a heat soaking zone, and a cooling zone, wherein the heating zone and the heat soaking zone are controlled within a temperature range of 810-850°C.

6. The method for manufacturing a high-strength steel plate with excellent hole-expanding properties according to claim 5, wherein, The hot rolling process involves hot finishing at an exit-side temperature of Ar3 or higher but below 1000°C.

7. The method for manufacturing a high-strength steel plate with excellent hole-expanding properties according to claim 5, wherein, The cooling after winding is carried out at a cooling rate of less than 0.1°C / second and except 0°C / second.

8. The method for manufacturing a high-strength steel plate with excellent hole-expanding properties according to claim 5, characterized in that, The cold rolling is carried out in one stand, and the total reduction rate is 55-70%.

9. The method for manufacturing a high-strength steel plate with excellent hole-expanding properties according to claim 5, wherein, After the secondary cooling, the process further includes an over-aging treatment step, which is performed for 200-800 seconds.

Images