High-strength thick steel plate with excellent hole-expanding properties and ductility, and method for manufacturing the same.

Abstract

The present invention relates to a steel suitable as a material for automobiles, and more particularly to a high-strength thick steel plate having excellent hole expandability and ductility, and a method for producing the same.

Description

[Technical Field]

[0001] The present invention relates to steel suitable as a material for automobiles, and specifically relates to a high-strength thick steel sheet excellent in hole expandability and ductility and a method for producing the same. [Background Art]

[0002] Recently, in the automobile industry, due to environmental regulations on CO₂ emissions and regulations on energy use, the use of high-strength steel is required for improving fuel efficiency or durability.

[0003] In particular, as regulations on the impact stability of automobiles are expanded, high-strength steel with excellent strength is adopted as a material for structural members such as members, seat rails, pillars and the like for improving the impact resistance of vehicle bodies.

[0004] Such automobile parts have complicated shapes depending on stability and design, and are mainly manufactured by forming with press dies, so high strength and a high level of formability are required.

[0005] However, while higher strength of steel is advantageous for absorbing impact energy, there is generally a problem that as strength increases, the elongation decreases and formability decreases. Furthermore, when the yield strength is excessively high, the inflow of the material into the die during forming is reduced, resulting in decreased formability and increased manufacturing cost.

[0006] Furthermore, since automotive parts have numerous molded areas where holes are expanded after machining, hole expandability (HER) is required for smooth molding. However, high-strength steel has poor hole expandability, which can lead to defects such as cracks occurring during molding. When hole expandability is reduced in this way, cracks can easily occur in the molded parts during a car collision, potentially causing the parts to break and jeopardizing passenger safety. Moreover, as standards for passenger safety become higher, the adoption of thicker materials to ensure rigidity is steadily increasing, particularly among some automobile manufacturers.

[0007] On the other hand, typical high-strength steels used in automotive materials 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, a type of ultra-high-strength steel, has a low yield ratio of approximately 0.5 to 0.6, making it easy to process, and it also has the advantage of having the second highest elongation rate after TRIP steel. Therefore, it is mainly applied to door outers, seat rails, seat belts, suspensions, arms, and wheel discs.

[0009] TRIP steel has a yield ratio in the range of 0.57 to 0.67, which gives it excellent formability (high ductility). Therefore, it is suitable for parts that require high formability, such as members, loops, seat belts, and bumper rails.

[0010] CP steel, with its low yield ratio, high elongation, and bendability, is used for side panels, underbody reinforcements, and other applications, while FB steel, with its excellent hole-expanding properties, is mainly used for suspension lower arms and wheel discs.

[0011] Among these, DP steel is mainly composed of ferrite with excellent ductility and hard phases with high strength (martensite phase, bainite phase), and may contain trace amounts of retained austenite. Such DP steel has low yield strength and high tensile strength, resulting in a low yield ratio (YR), and excellent properties such as 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, the degree of recrystallization, and the uniformity of distribution, it can be manufactured as a high-strength steel with high hole-expanding properties.

[0012] However, in order to achieve ultra-high strength of 980 MPa or more, the proportion of hard phases such as martensite, which are advantageous for improving strength, must be increased. In this case, however, the yield strength increases, which leads to the problem of defects such as cracks occurring during press forming.

[0013] Generally, DP steel for automobiles is manufactured by first producing slabs through steelmaking and continuous casting processes, then applying heating, rough rolling, and finish hot rolling to these slabs to obtain hot-rolled coils, and finally going through an annealing process to produce the final product.

[0014] Here, the annealing process is mainly performed during the manufacturing of cold-rolled steel sheets. Cold-rolled steel sheets are manufactured by acid-washing hot-rolled coils to remove surface scale, cold-rolling them at room temperature with a constant reduction ratio, and then undergoing an annealing process and, if necessary, a further temper-rolling process.

[0015] Cold-rolled steel sheets (cold-rolled material) obtained by cold rolling are themselves in a very hard state and are unsuitable for manufacturing parts that require workability. Therefore, in a subsequent process, they can be softened by heat treatment in a continuous annealing furnace to improve workability.

[0016] For example, in the annealing process, steel sheets (cold-rolled material) are heated to approximately 650-850°C in a heating furnace and then maintained at that temperature for a certain period of time. This process reduces hardness through recrystallization and phase transformation, thereby improving workability.

[0017] Steel sheets that have not undergone the annealing process have high hardness, especially surface hardness, resulting in poor workability. In contrast, steel sheets that have undergone the annealing process have a recrystallized structure, which lowers their hardness, yield point, and tensile strength, thus improving workability.

[0018] A typical method for lowering the yield strength of DP steel is to completely recrystallize the ferrite during the heating process in continuous annealing to produce an equiaxed crystal structure. This is advantageous because it ensures that the austenite formation and growth in subsequent processes also occur in an equiaxed crystal structure, resulting in the formation of a small-grained, uniform austenite phase.

[0019] Furthermore, in the case of thick materials, a relatively thick hot-rolled thickness must be ensured to maintain a certain reduction ratio, which leads to a problem of increased load and reduced operability during subsequent cold rolling. When manufacturing thick materials, if the reduction ratio is low, the non-uniformity of the microstructure due to unrecrystallized ferrite increases during annealing, resulting in a higher yield strength, and the direction of cold rolling is maintained in the microstructure, leading to a problem of reduced workability. Therefore, in the case of thick materials, due to the dimensional characteristics, material variation in the thickness direction is unavoidable, so techniques to homogenize the material as much as possible are required to improve workability and usability.

[0020] On the other hand, Patent Document 1 discloses that by utilizing Ti, Mo, etc. to form fine precipitates and composing the microstructure with ferrite, bainite, and martensite phases, it is possible to ensure hole-expanding properties and elongation.

[0021] However, this paper suffers from problems with weldability and liquid metal embrittlement (LME) due to the excessive addition of silicon to introduce carbon and bainite for the formation of fine precipitates. Furthermore, problems still exist due to the hardness difference between the soft and hard phases, and the formation of an excessive fraction of the bainite phase for high hole-expanding properties results in high yield strength, making it difficult to process and resulting in poor elongation.

[0022] Judging from the above-mentioned prior art, in order to simultaneously improve formability properties such as elongation percentage and hole expandability of thick high-strength steel, development of a method that can reduce yield strength and improve workability while forming a uniform structure within the steel is required.

Prior Art Literature

Patent Literature

[0023]

Patent Literature 1

Summary of the Invention

Problem to be Solved by the Invention

[0024] One aspect of the present invention aims to provide, as a material suitable for applications such as automotive structural members, a high-strength thick steel sheet that has a low yield ratio, high strength, and is excellent in formability such as hole expandability due to improved ductility, and a method for producing the same.

[0025] The problems of the present invention are not limited to the contents described above. The problems of the present invention can be understood from the entire contents of the present specification, and a person having ordinary knowledge in the technical field to which the present invention pertains will have no difficulty in understanding further problems of the present invention.

Means for Solving the Problem

[0026] One aspect of the present invention comprises, by weight%, carbon (C): 0.05 to 0.12%, manganese (Mn): 2.0 to 3.0%, silicon (Si): 0.5% or less (excluding 0%), chromium (Cr): 1.0% or less (excluding 0%), niobium (Nb): 0.1% or less (excluding 0%), titanium (Ti): 0.1% or less (excluding 0%), boron (B): 0.003% or less (excluding 0%), aluminum (sol.Al): 0.02 to 0.05%, phosphorus (P): 0.05% or less (excluding 0%), sulfur (S): 0.01% or less (excluding 0%), nitrogen (N): 0.01% or less (excluding 0%), iron (Fe) and other unavoidable impurities, Provided is a high-strength thick steel sheet excellent in hole expandability and ductility, comprising, as a microstructure, 10 to 30% area fraction of ferrite, 10 to 25% of recrystallized ferrite bridges, 20 to 30% of bainite, and the balance of martensite.

[0027] Another aspect of the present invention provides a method for producing a high-strength thick steel sheet excellent in hole expandability and ductility, comprising the steps of: preparing a steel slab; heating the steel slab in a temperature range of 1100 to 1300°C; hot-rolling the heated steel slab to produce a hot-rolled steel sheet; coiling the hot-rolled steel sheet in a temperature range of 400 to 700°C; after the coiling, cooling the hot-rolled steel sheet to room temperature; cold-rolling the cooled hot-rolled steel sheet at a cold reduction rate of 55 to 80% to produce a cold-rolled steel sheet; subjecting the cold-rolled steel sheet to continuous annealing; after the continuous annealing, performing primary cooling to a temperature range of 650 to 700°C at an average cooling rate of 1 to 10°C / s; and after the primary cooling, performing secondary cooling to a temperature range of 450 to 500°C at an average cooling rate of 5 to 50°C / s, wherein the continuous annealing is performed in an equipment provided with a heating zone, a soaking zone and a cooling zone, and when the cold-rolled steel sheet is heated to a temperature in the heating zone, the method passes through a recrystallization zone where the temperature is maintained at 600 to 700°C for 1 to 3 minutes. [Effect of the Invention]

[0028] According to the present invention, despite having high strength, a thick steel sheet that is excellent in hole expandability and has improved formability and collision resistance can be provided.

[0029] As described above, the steel sheet of the present invention with improved formability can prevent processing defects such as cracks and wrinkles during press forming, and thus has the effect of being suitably applicable to structural parts and the like that require processing into complicated shapes. Furthermore, when an automobile equipped with such parts inevitably collides, the present invention is also effective in producing a material with improved collision resistance such that defects such as cracks are less likely to occur. [Brief Description of Drawings]

[0030] [Figure 1]This shows the thermal history and phase transformation history during continuous annealing according to one embodiment of the present invention. In Figure 1, the dashed line shows the thermal history during conventional continuous annealing, and the solid line shows the thermal history during continuous annealing according to the present invention. [Figure 2] (a) shows the void formation mechanism within the tissue, and (b) shows the interface strengthening mechanism within the tissue in an example of an invention according to one embodiment of the present invention. [Figure 3] This shows a microstructural photograph of a comparative example according to one embodiment of the present invention. [Figure 4] This figure shows a microstructural photograph of an example of the present invention according to one embodiment of the present invention (the arrows in Figure 4 represent the recrystallized ferrite bridge structure). [Modes for carrying out the invention]

[0031] The inventors of this invention have diligently conducted research to develop a material with a level of moldability suitable for use in automotive parts and other applications that require processing into complex shapes.

[0032] In particular, the inventors have derived a microstructure that can improve crack resistance between hard phases in thick steel sheets for automobiles, where the cold reduction ratio is inevitably relatively low. They have also confirmed that the target conditions can be achieved by refining the hard phase and controlling the crystal grain shape, which is advantageous for preventing the formation and propagation of voids, and have thus completed the present invention.

[0033] In particular, the present invention has technical significance in introducing recrystallized ferrite bridges having a structure that connects the hard phases to each other so as to eliminate the unidirectionality of the hard phases, and in optimizing the alloy composition and manufacturing conditions in forming such a structure.

[0034] The present invention will be described in detail below.

[0035] A high-strength thick steel sheet with excellent hole-expandability and ductility according to one aspect of the present invention may contain, by weight percent, carbon (C): 0.05-0.12%, manganese (Mn): 2.0-3.0%, silicon (Si): 0.5% or less (excluding 0%), chromium (Cr): 1.0% or less (excluding 0%), niobium (Nb): 0.1% or less (excluding 0%), titanium (Ti): 0.1% or less (excluding 0%), boron (B): 0.003% or less (excluding 0%), aluminum (sol.Al): 0.02-0.05%, phosphorus (P): 0.05% or less (excluding 0%), sulfur (S): 0.01% or less (excluding 0%), and nitrogen (N): 0.01% or less (excluding 0%).

[0036] The following will explain in detail the reasons for limiting the alloy composition of the high-strength thick steel sheet provided in the present invention as described above.

[0037] On the other hand, unless otherwise specified in this invention, the content of each element is based on weight, and the proportion of tissue is based on area.

[0038] Carbon (C): 0.05~0.12% Carbon (C) is an important element added for solid solution strengthening. Such C contributes to improving the strength of steel by combining with precipitate elements to form fine precipitates.

[0039] If the C content exceeds 0.12%, the hardening ability increases, and during steel manufacturing, martensite is formed during cooling, leading to an excessive increase in strength, but also a decrease in elongation. Furthermore, weldability is poor, and welding defects may occur during processing into parts. On the other hand, if the C content is less than 0.05%, it becomes difficult to achieve the target level of strength.

[0040] Therefore, the above C can include 0.05 to 0.12%. More favorably, it can include 0.06% or more and 0.10% or less.

[0041] Manganese (Mn): 2.0-3.0% Manganese (Mn) is an element that is advantageous for solid solution strengthening of steel because it precipitates sulfur (S) in steel as MnS, preventing hot brittleness caused by the formation of FeS.

[0042] If the Mn content is less than 2.0%, not only will the above-mentioned effects not be obtained, but it will also be difficult to secure the target level of strength. On the other hand, if the content exceeds 3.0%, there is a high possibility of problems occurring in weldability and hot rolling properties, and the increased hardening ability may lead to easier martensite formation, potentially reducing ductility. Furthermore, there is a problem that excessive Mn-band (Mn oxide band) formation in the microstructure increases the risk of defects such as processing cracks. In addition, there is a problem that Mn oxide dissolves to the surface during annealing, significantly hindering plating properties.

[0043] Therefore, the above Mn can be included in an amount of 2.0-3.0%. More favorably, it can be included in an amount of 2.2% or more and 2.8% or less.

[0044] Silicon (Si): 0.5% or less (excluding 0%) Silicon (Si) is a ferrite-stabilizing element, which is advantageous in promoting ferrite transformation and ensuring a target level of ferrite fraction. Furthermore, it has good solid solution strengthening ability, effectively increasing the strength of ferrite, and is a useful element in ensuring strength without reducing the ductility of steel.

[0045] When the Si content exceeds 0.5%, the solid solution strengthening effect becomes excessive, which actually reduces ductility and induces surface scale defects, negatively impacting the surface quality of the plating. Furthermore, it hinders chemical conversion treatment.

[0046] Therefore, the above Si can be included in an amount of 0.5% or less, and 0% can be excluded. More favorably, it can be included in an amount of 0.1% or more.

[0047] Chromium (Cr): 1.0% or less (excluding 0%) Chromium (Cr) is an element that exhibits a hardening effect during cooling, facilitating the formation of the bainite phase, and suppressing the formation of the martensite phase during annealing heat treatment, while also contributing to improved strength by forming fine carbides.

[0048] Furthermore, in this invention, by including the above-mentioned Cr at an appropriate level, it acts as a ferrite-stabilizing element during heating, delaying the austenite phase transformation reaction, while the phase transformation starts at a higher temperature. As a result, it remains in the region where only recrystallization occurs (Trex~A1) for a longer period during heating. Consequently, a recrystallized ferrite bridge structure can be secured.

[0049] If the Cr content exceeds 1.0%, the intended recrystallized ferrite bridges may not form, reducing the ductility and hole-expanding properties of the steel. If carbides form at the grain boundaries, the strength and elongation may be inferior. Furthermore, there is the problem of increased manufacturing costs.

[0050] Therefore, the above Cr can be included at a concentration of 1.0% or less, and 0% can be excluded. More favorably, it can be included at a concentration of 0.01% or more.

[0051] Niobium (Nb): 0.1% or less (excluding 0%) Niobium (Nb) is an element that segregates at austenite grain boundaries, suppresses the coarsening of austenite grains during annealing heat treatment, and contributes to improved strength by forming fine carbides.

[0052] When the Nb content exceeds 0.1%, coarse carbides precipitate, potentially reducing the carbon content in the steel, which can lead to decreased strength and elongation, and also increases manufacturing costs.

[0053] Therefore, the above Nb can be included in amounts of 0.1% or less, and can be excluded if it is 0%.

[0054] Titanium (Ti): 0.1% or less (excluding 0%) Titanium (Ti) is an element that forms fine carbides, contributing to the securing of yield strength and tensile strength. Furthermore, Ti causes nitrogen in steel to precipitate as TiN, and has the effect of suppressing the formation of AlN by Al, which is inevitably present in steel, thus reducing the possibility of crack formation during continuous casting.

[0055] When the Ti content exceeds 0.1%, coarse carbides may precipitate, potentially reducing the carbon content in the steel and decreasing its strength and elongation. Furthermore, it may induce nozzle clogging during continuous casting, leading to increased manufacturing costs.

[0056] Therefore, the above Ti can be included in amounts of 0.1% or less, and 0% can be excluded.

[0057] Boron (B): 0.003% or less (excluding 0%) Boron (B) is an element that delays the transformation of austenite into pearlite during the cooling process after annealing heat treatment. However, if its content exceeds 0.003%, B may become excessively concentrated on the surface, potentially leading to a deterioration in plating adhesion.

[0058] Therefore, the above B can be included in an amount of 0.003% or less, and 0% can be excluded considering the level that will inevitably be added.

[0059] Aluminum (sol.Al): 0.02~0.05% Aluminum (sol.Al) is an element added to steel for grain refinement and deoxidation. If its content is less than 0.02%, it is not possible to produce aluminum-killed steel in a stable state. On the other hand, if its content exceeds 0.05%, it has the effect of refining the crystal grain and improving strength, but the excessive formation of inclusions during continuous casting operations in steelmaking increases the risk of surface defects in plated steel sheets.

[0060] Therefore, the above sol.Al can be included in an amount of 0.02 to 0.05%.

[0061] Phosphorus (P): 0.05% or less (excluding 0%) Phosphorus (P) is the substitutional element with the greatest solid solution strengthening effect, and is advantageous for improving in-plane anisotropy and ensuring strength without significantly reducing formability. However, when P is added in excess, the possibility of brittle fracture increases significantly, increasing the likelihood of slab fracture during hot rolling and hindering the surface properties of the plated material.

[0062] Therefore, in the present invention, the content of P can be controlled to 0.05% or less, and 0% can be excluded considering the level that is inevitably added.

[0063] Sulfur (S): 0.01% or less (excluding 0%) Sulfur (S) is an element that is inevitably added as an impurity element in steel, and since it inhibits ductility, it is preferable to control its content to be as low as possible. In particular, since S increases the possibility of red-hot brittleness, it is preferable to control its content to 0.01% or less. However, considering the level that is inevitably added during the manufacturing process, 0% can be excluded.

[0064] Nitrogen (N): 0.01% or less (excluding 0%) Nitrogen (N) is a solid solution strengthening element, but if its content exceeds 0.01%, the risk of brittleness increases, and it can combine with Al in the steel to cause excessive AlN precipitation, which may impair continuous casting quality.

[0065] Therefore, the above N can be included in an amount of 0.01% or less, and 0% can be excluded considering the level that will inevitably be added.

[0066] The remaining component of this invention is iron (Fe). However, in the normal manufacturing process, unintended impurities may inevitably be introduced from the raw materials or the surrounding environment, and therefore cannot be eliminated. Since these impurities are easily recognizable to any technician in the normal manufacturing process, their details are not specifically mentioned in this specification.

[0067] The steel sheet of the present invention having the alloy composition described above can be composed of a recrystallized ferrite bridge structure formed by linking the hard phases together with ferrite, which is a soft phase, and bainite and martensite, which are hard phases, as its microstructure.

[0068] In the present invention, the most significant change in the microstructure phase due to the formation of recrystallized ferrite bridges is the disappearance of the rolling directionality of the existing ferrite and the increased degree of coupling around the hard phase. Furthermore, the formation of recrystallized ferrite bridges during heating reduces the formation location of reverse-transformed austenite, delaying the formation of austenite at high temperatures and enabling the generation of a smaller secondary phase after cooling. The unrecrystallized ferrite region is an elongated structure with a rolling directionality, remaining as an irregular, rough interface, while the recrystallized ferrite bridge grain boundaries are characterized by having a polygonal, smooth interface. The above-mentioned recrystallized ferrite bridges can be determined, for example, by classification using EBSD crystal orientation (electron backscatter diffraction orientation), or by optical classification by etching with an aqueous hydrogen peroxide solution (e.g., 140 ml distilled water, 100 ml hydrogen peroxide, 4 g oxalic acid, 2 ml sulfuric acid, 1.5 ml hydrofluoric acid).

[0069] Specifically, the steel sheet of the present invention contains a ferrite phase in an area fraction of 10-30%, a recrystallized ferrite bridge phase in an area fraction of 10-25%, and can also contain 20-30% bainite, which is a hard phase, and the remainder being martensite. In addition, it can contain a trace amount of retained austenite phase.

[0070] In the present invention, the above-mentioned recrystallized ferrite bridge phase is a structure that is advantageous in eliminating the unidirectionality of the hard phase and suppressing the propagation of voids that occur along the grain boundaries of the hard phase, and is a structure that is distinct from existing granular ferrite (polygonal ferrite).

[0071] Furthermore, the above-mentioned recrystallized ferrite bridge is a structure distinct from general recrystallized ferrite, being relatively coarse, and preferably having an average size of 1 to 6 μm based on the equivalent circular diameter. If the size of the recrystallized ferrite bridge phase is less than 1 μm, it is difficult to eliminate the directionality of the hard phase, and the desired effect cannot be obtained. On the other hand, if its size exceeds 6 μm, the structure becomes excessively coarse, and physical properties such as strength may be impaired.

[0072] As shown in Figure 2, when the hard phase is formed to have directionality (a), the voids formed along the grain boundaries become easier to connect, which is advantageous for crack propagation. On the other hand, in the structure intended by the present invention (b), in particular, when the recrystallized ferrite bridge phase is formed as a structure that connects the hard phases to each other while eliminating the unidirectionality of the hard phase, voids become less likely to coalesce along the grain boundaries, and as a result, the occurrence and propagation of cracks are greatly suppressed.

[0073] If the fraction of the recrystallized ferrite bridge phase is excessively high, the relative fraction of the hard phase becomes low, making it impossible to secure the target level of strength. Taking this into consideration, the recrystallized ferrite bridge phase can be included in a fraction of 25% or less. On the other hand, if its fraction is less than 10%, the above-mentioned effects (such as the elimination of unidirectionality in the hard phase and the suppression of void propagation) cannot be sufficiently obtained, and the hole-expanding properties will be inferior.

[0074] In other words, the technical significance of the present invention lies in further improving moldability by introducing a recrystallized ferrite bridge phase in addition to the soft ferrite phase and the hard bainite and martensite phases, and by controlling the shape and distribution of the recrystallized ferrite bridge phase.

[0075] If the ferrite phase fraction is less than 10%, it is unfavorable for ensuring the ductility of the steel. On the other hand, if the fraction exceeds 30%, the relative fraction of the hard phase becomes low, making it difficult to achieve the target level of strength.

[0076] If the bainite phase fraction is less than 20%, it becomes difficult to ensure sufficient strength, and there is also a problem of a large difference in hardness between the soft phase and the martensite phase. On the other hand, if the fraction exceeds 30%, the proportion of the soft phase decreases, making it difficult to ensure ductility.

[0077] Excluding the ferrite phase, recrystallized ferrite bridge phase, and bainite phase mentioned above, the martensite phase does not have a specific fractional limit, but it may contain an area fraction of 15% or more to ensure ultra-high strength of 980 MPa or more. However, if the fraction of the martensite phase exceeds 60%, the ductility decreases, making it difficult to ensure the target level of formability.

[0078] On the other hand, it is advantageous that the residual austenite phase does not exceed 3%, and it is clear that even at 0%, it is possible to secure the intended physical properties.

[0079] The high-strength thick steel sheet of the present invention, having the microstructure described above, has a tensile strength of 980 MPa or more, a yield strength of 550 to 700 MPa, and an elongation (total elongation) of 14% or more, thus possessing both high strength and high ductility.

[0080] Furthermore, the steel plate, having a hole expansion ratio (HER) of 30% or more, exhibits excellent resistance to cracks that may occur during processing and resistance to impact fracture.

[0081] Furthermore, the high-strength thick steel plate of the present invention may have a thickness of 1 to 3 mm, and more preferably 1.5 to 2.5 mm.

[0082] The following describes in detail a method for manufacturing a high-strength thick steel plate with excellent hole-expanding properties and ductility according to another aspect of the present invention.

[0083] Simply put, the present invention allows for the production of a target steel sheet through the processes of [steel slab heating - hot rolling - coiling - cold rolling - continuous annealing], and each process will be described in detail below.

[0084] [Heating steel slabs] First, a steel slab satisfying the alloy composition described above can be prepared and then heated.

[0085] This process is carried out to ensure a smooth subsequent hot rolling process and to obtain the desired physical properties of the steel sheet. In this invention, there are no particular restrictions on the conditions of this heating process; ordinary conditions are acceptable. For example, the heating process can be carried out in a temperature range of 1100 to 1300°C.

[0086] [Hot rolling] As described above, the heated steel slab can be hot-rolled to produce a hot-rolled steel sheet, and in this process, the finish hot-rolling can be performed at an exit temperature of Ar3 or higher and 1000°C or lower.

[0087] During the above-mentioned finish hot rolling, if the exit temperature is below Ar3, the hot deformation resistance increases sharply, causing the top, tail, and edge of the hot-rolled coil to become single-phase regions, increasing in-plane anisotropy and potentially degrading formability. On the other hand, if the temperature exceeds 1000°C, while the relative reduction in rolling load is advantageous for productivity, there is a risk of thick oxide scale formation.

[0088] More specifically, the above-mentioned finish hot rolling can be performed in a temperature range of 760 to 940°C.

[0089] [Rewind] The hot-rolled steel sheet manufactured as described above can be wound into a coil.

[0090] The above winding process can be carried out in a temperature range of 400 to 700°C. If the winding temperature is below 400°C, excessive martensite or recrystallized ferrite bridge phases may form, leading to an excessive increase in the strength of the hot-rolled steel sheet and potentially causing problems such as shape defects due to load during subsequent cold rolling. On the other hand, if the winding temperature exceeds 700°C, surface scale increases, resulting in a deterioration of pickling properties.

[0091] [cooling] It is preferable to cool the wound hot-rolled steel sheet to room temperature at an average cooling rate of 0.1°C / s or less (excluding 0°C / s). In this case, the wound hot-rolled steel sheet can be cooled after processes such as transport and stacking, but it is important to clarify that the processes before cooling are not limited to these.

[0092] In this way, by cooling the rolled hot-rolled steel sheet at a constant rate, it is possible to obtain a hot-rolled steel sheet in which carbides, which serve as austenite nucleation sites, are finely dispersed.

[0093] [Cold rolling] The hot-rolled steel sheet wound up as described above can be cold-rolled to produce a cold-rolled steel sheet, and in this invention, the cold rolling can be performed at a cold reduction ratio of 55 to 80%.

[0094] The present invention allows for the application of an appropriate cold reduction ratio during cold rolling, and further promotes the recrystallization of ferrite during the subsequent continuous annealing process. This induces the formation of fine ferrite, and the austenite formed at the ferrite grain boundaries can also be made small and uniform.

[0095] If the reduction ratio during cold rolling is less than 55%, ferrite recrystallization is delayed, making it difficult to obtain a fine and uniform austenite phase. On the other hand, if it exceeds 80%, excessive recrystallization reduces the yield strength too much, making it impossible to secure the target level of strength. More advantageously, it can be carried out at 78% or less.

[0096] In the present invention, the above cold reduction ratio can be achieved not only from a hot-rolled thick material having a constant thickness using the high rolling capacity of a ZRM (for example, at the 5000 KN / mm level), but also by a process of achieving the target reduction ratio through repeated rolling using a reversing mill. As a non-limiting example, the above hot-rolled thick material can have a thickness of 4 to 8 mm, and it is revealed that a cold rolling process using a reversing mill can be performed when the thickness of the hot-rolled thick material is 6 mm or more.

[0097] A reversing rolling mill is generally a type of rolling mill used for rolling thin materials, and refers to a rolling mill that rolls the material while moving it back and forth between a pair of rolls, and one of the reversing paths of the material can be set to one pass.

[0098] The present invention reveals that hot-rolled steel sheets can be pickled before cold rolling, and that the pickling process can be carried out by conventional methods.

[0099] [Continuous annealing] It is preferable to perform continuous annealing on the cold-rolled steel sheet produced as described above. This continuous annealing can be carried out, for example, in a continuous annealing furnace (CAL).

[0100] Typically, a continuous annealing furnace (CAL) can consist of a heating zone, a homogenizing zone, a cooling zone (including slow cooling and rapid cooling zones), and (if necessary, an over-aging zone). In such a continuous annealing furnace, cold-rolled steel sheets are loaded, heated to a specific temperature in the heating zone, and then maintained in the homogenizing zone for a certain period of time after reaching the target temperature.

[0101] In this invention, the temperatures of the heating zone and the soaking zone during continuous annealing can be controlled identically, which means that the end temperature of the heating zone and the start temperature of the soaking zone can be controlled identically (Figure 1).

[0102] Specifically, the temperatures of the heating zone and the uniform zone can be controlled to 790-850°C.

[0103] If the temperature of the heating zone is below 790°C, it becomes impossible to apply sufficient heat for recrystallization. On the other hand, if the temperature exceeds 850°C, productivity decreases, an excessive austenite phase is formed, and after subsequent cooling, the fraction of the hard phase increases significantly, potentially resulting in poor ductility of the steel.

[0104] Furthermore, if the temperature of the homogenization zone is below 790°C, excessive cooling is required at the end of the heating zone, which is economically disadvantageous and may result in insufficient heat for recrystallization. On the other hand, if the temperature exceeds 850°C, the austenite fraction becomes excessive, and the moldability may decrease due to an increase in the hard phase during cooling.

[0105] Increasing the temperature of the homogenized zone within the aforementioned temperature range increases the fraction of the hard phase in the final structure, thereby increasing the yield strength, while the introduction of bainite reduces the hardness difference between phases, improving the hole expansion rate.

[0106] On the other hand, the present invention is characterized by inducing the formation of recrystallized ferrite bridges by causing sufficient recrystallization during the annealing process.

[0107] Specifically, the present invention allows for the introduction of a recrystallization zone, which maintains an intermediate temperature for a certain period of time, when raising the temperature of the cold-rolled steel sheet to the temperature range of the heating zone. More preferably, the process involves maintaining the temperature in the range of 600 to 700°C for 1 to 3 minutes (Figure). 1 (Dashed line graph).

[0108] If the temperature of the recrystallization zone is below 600°C or the maintenance time is less than 1 minute, the ferrite will not recrystallize sufficiently, and the recrystallized ferrite bridge phase cannot be formed at the target fraction. On the other hand, if the temperature exceeds 700°C or the maintenance time exceeds 3 minutes, the recrystallization will be excessive, which may lead to a decrease in strength and a deterioration of physical properties due to grain coarsening.

[0109] The present invention improves workability by introducing a recrystallized ferrite bridge phase along with an appropriate fraction of hard and soft phases as the final microstructure through the above-described recrystallization zone process, thereby enhancing crack toughness, or in other words, crack resistance, while maintaining strength.

[0110] [Gradual cooling] As described above, the target microstructure can be formed by cooling the cold-rolled steel sheet that has undergone continuous annealing treatment as described above, and it is preferable to perform this cooling stepwise.

[0111] In the present invention, the above-mentioned stepwise cooling can consist of primary cooling and secondary cooling. Specifically, after the above-mentioned continuous annealing, primary cooling can be performed to a temperature range of 650-700°C at an average cooling rate of 1-10°C / s, followed by secondary cooling to a temperature range of 450-500°C at an average cooling rate of 5-50°C / s.

[0112] In this case, by performing the primary cooling more slowly than the secondary cooling, it is possible to suppress plate shape defects caused by a rapid temperature drop during the secondary cooling phase, which is a relatively rapid cooling phase.

[0113] If the end temperature during the primary cooling described above is below 650°C, the low temperature results in low carbon diffusion activity, leading to a high carbon concentration within the ferrite. Conversely, as the carbon concentration within the austenite decreases, the proportion of the hard phase becomes excessive, increasing the yield ratio and thus increasing the tendency for cracks to occur during processing. Additionally, the cooling rate between the homogenized zone and the cooling zone (slow cooling zone) becomes too high, resulting in a problem of uneven plate shape. If the end temperature exceeds 700°C, there is the disadvantage of requiring an excessively high cooling rate during subsequent cooling (secondary cooling).

[0114] Furthermore, if the average cooling rate during the primary cooling described above exceeds 10°C / s, sufficient carbon diffusion will not occur. On the other hand, considering productivity, the primary cooling described above can be performed at an average cooling rate of 1°C / s or higher.

[0115] As mentioned above, after primary cooling is complete, rapid cooling (secondary cooling) can be performed at a cooling rate above a certain level. In this case, if the secondary cooling completion temperature is below 450°C, cooling variations may occur in the width and length directions of the steel plate, potentially degrading the plate shape. On the other hand, if the temperature exceeds 500°C, it may become impossible to secure a sufficient hard phase, potentially reducing the strength.

[0116] Furthermore, if the average cooling rate during the secondary cooling described above is less than 5°C / s, the proportion of the soft phase may become excessive. On the other hand, if it exceeds 50°C / s, the hard phase may become insufficient.

[0117] Furthermore, if necessary, overaging treatment can be performed after the above-mentioned stepwise cooling is completed.

[0118] The above overaging treatment is a process of maintaining the temperature for a certain period after the secondary cooling is completed, and has the effect of improving shape quality by performing uniform heat treatment in the width and length directions of the coil. For this reason, the above overaging treatment can be performed for 200 to 800 seconds.

[0119] Since the above overaging treatment can be performed immediately after the completion of secondary cooling, its temperature can be the same as the secondary cooling completion temperature, or within the range of the secondary cooling completion temperature, and it can also be performed at a lower temperature. More advantageously, it can be performed in the temperature range of 300 to 450°C.

[0120] The high-strength thick steel sheet of the present invention, manufactured as described above, has a microstructure composed of hard and soft phases. By maximizing ferrite recrystallization through particularly optimized cold rolling and annealing processes, it is possible to obtain a structure in which the martensite phase, which is the hard phase, is uniformly distributed in the recrystallized ferrite matrix. Furthermore, by introducing relatively coarse recrystallized ferrite bridge phases so as to connect the hard phases, the crack resistance during processing is enhanced.

[0121] Therefore, despite having high strength of 980 MPa or more in tensile strength, the thick steel sheet of the present invention ensures excellent formability, such as hole-expanding properties, by securing a low yield ratio and high ductility. [Examples]

[0122] The present invention will be described in more detail below with reference to examples. However, such examples are provided to illustrate the implementation of the present invention and do not limit the present invention. The scope of the present invention is determined by the matters described in the claims and matters that can be reasonably inferred therefrom.

[0123] (Examples) After preparing steel slabs with the alloy compositions shown in Table 1 below, each slab was heated at 1200°C for 1 hour, and then hot-rolled at a finish rolling temperature of 880-920°C to produce hot-rolled steel sheets. Subsequently, each hot-rolled steel sheet was wound at 650°C and cooled to room temperature at a cooling rate of 0.1°C / s. After that, the wound hot-rolled steel sheets were cold-rolled and continuously annealed under the conditions shown in Table 2 below, followed by stepwise cooling (primary-secondary) and then over-aging treatment at 360°C for 520 seconds to produce a final steel sheet with a thickness of 1.8 mm.

[0124] During this process, the primary cooling was performed at an average cooling rate of 3°C / s, and the secondary cooling at an average cooling rate of 20°C / s.

[0125] The microstructure of each steel sheet manufactured as described above was observed, and the tensile and workability characteristics, as well as the utilization property index of the processing steps such as hole expansion rate, were evaluated. The results are shown in Table 3 below.

[0126] In this case, for each test specimen, a tensile test specimen of JIS No. 5 size was taken perpendicular to the rolling direction, and then the tensile test was performed at a strain rate of 0.01 / s.

[0127] On the other hand, the hole expansion rate (HER, %) measurement test was conducted according to the ISO 16630 standard. Specifically, after punching a circular hole in the test specimen, the hole was expanded using a conical punch, and the amount of hole expansion until a crack formed at the edge of the hole penetrated through in the thickness direction was expressed as a percentage of the initial hole. At this time, the dimensions of the test specimen were 120 mm × 120 mm, the clearance was 12%, the punching hole diameter was 10 mm, the punching holding load was 20 tons, and the test speed was set to 12 mm / min.

[0128] The bainite phase, which corresponds to the hard phase of the tissue, was observed via picral etching, and the martensite phase was observed via nital etching, followed by observation via SEM at 2000x and 5000x magnification.

[0129] Furthermore, the fractions of the ferrite phase and the recrystallized ferrite bridge phase were measured using SEM and an image analyzer program after nital etching.

[0130] [Table 1]

[0131] [Table 2]

[0132] [Table 3]

[0133] As shown in Tables 1-3 above, Examples 1-6 of the invention, in which the alloy composition of the steel and manufacturing conditions, particularly the cold rolling and continuous annealing processes, satisfy all the conditions proposed in this invention, were formed by sufficient recrystallization of ferrite during the annealing process after cold rolling, resulting in the hard phase being linked by recrystallized ferrite bridge phases. As a result, the material had high strength while possessing a yield strength suitable for plate processing and excellent elongation. Furthermore, it was confirmed that the homogeneous distribution of the hard phase resulted in excellent hole-expanding properties and the ability to secure the target level of formability.

[0134] On the other hand, in Comparative Examples 1-4 and 6-9, which did not go through the recrystallization zone during the heating process of continuous annealing in the steel sheet manufacturing process, recrystallization did not occur sufficiently, resulting in an insufficient recrystallized ferrite bridge phase. Of these, Comparative Examples 1-4, which had relatively low continuous annealing temperatures, were inferior in at least one of the physical properties of elongation and hole-expanding properties, while Comparative Examples 6-9, which had relatively high continuous annealing temperatures, had excessive bainite phase formation, resulting in excessively high yield strength and poor elongation.

[0135] Comparative Examples 5 and 10 had sufficient annealing temperature for recrystallization and sufficient austenite stability for strength, but insufficient reduction ratio prevented sufficient recrystallization, resulting in a failure to form a uniform structure. Consequently, they had inferior elongation and relatively high yield strength.

[0136] Furthermore, comparative examples 11-14, which had very low secondary cooling temperatures, had excessively high yield strength, increasing the risk of crack formation during processing, and also exhibited inferior elongation due to the absence of a recrystallized ferrite bridge phase.

[0137] Figure 3 shows microstructural images of Comparative Examples 1, 4-5, and 9, and Figure 4 shows microstructural images of Invention Examples 1, 3-4, and 6.

[0138] As shown in Figure 3, in Comparative Examples 1, 4, and 9, the recrystallization zone process was not introduced during the heating process of continuous annealing, so the connecting structure of the hard phase by recrystallized ferrite bridges could hardly be observed. In addition, in Comparative Example 5, the fraction of the recrystallized zone was low due to insufficient reduction ratio, and recrystallization was insufficient due to low driving force, resulting in the formation of a structure in which the hard phases clustered together, thus forming a microstructure with low crack propagation resistance.

[0139] On the other hand, as shown in Figure 4, in the example of the invention, a relatively coarse recrystallized ferrite bridge phase is observed through the recrystallization process, and it can be seen that the unidirectionality of the hard phase is eliminated by this phase.

Claims

1. By weight percent, it contains carbon (C): 0.05-0.12%, manganese (Mn): 2.0-3.0%, silicon (Si): 0.5% or less (excluding 0%), chromium (Cr): 1.0% or less (excluding 0%), niobium (Nb): 0.1% or less (excluding 0%), titanium (Ti): 0.1% or less (excluding 0%), boron (B): 0.003% or less (excluding 0%), aluminum (sol. Al): 0.02-0.05%, phosphorus (P): 0.05% or less (excluding 0%), sulfur (S): 0.01% or less (excluding 0%), and nitrogen (N): 0.01% or less (excluding 0%), and consists of iron (Fe) and other unavoidable impurities. The microstructure includes ferrite with an area fraction of 10-30%, recrystallized ferrite bridges with an area fraction of 10-25%, and a hard phase. The aforementioned hard phase consists of 20-30% bainite and 39.4% or less martensite. The aforementioned recrystallized ferrite bridges exist to connect the hard phase, and the average circle equivalent diameter is 1 to 6 μm, resulting in a high-strength, thick steel plate with excellent hole-expanding properties and ductility.

2. The steel sheet is a high-strength, thick steel sheet with excellent hole-expanding properties and ductility, as described in claim 1, wherein the steel sheet contains a martensite phase in an area fraction of 15% or more.

3. The steel sheet further comprises a retained austenite phase in an area fraction of 3% or less (including 0%), wherein the steel sheet is a high-strength thick steel sheet with excellent hole-expanding properties and ductility as described in claim 1.

4. The steel plate is a high-strength, thick steel plate with excellent hole-expanding properties and ductility, as described in claim 1, having a tensile strength of 980 MPa or more, a yield strength of 550 to 700 MPa, and a total elongation of 14% or more.

5. The steel plate is a high-strength, thick steel plate with excellent hole-expanding properties and ductility, as described in claim 1, wherein the hole-expanding ratio (HER) is 30% or more.

6. The steel plate is a high-strength thick steel plate with excellent hole-expanding properties and ductility, as described in claim 1, having a thickness of 1 to 3 mm.

7. The process involves preparing a steel slab containing, by weight percent, carbon (C): 0.05-0.12%, manganese (Mn): 2.0-3.0%, silicon (Si): 0.5% or less (excluding 0%), chromium (Cr): 1.0% or less (excluding 0%), niobium (Nb): 0.1% or less (excluding 0%), titanium (Ti): 0.1% or less (excluding 0%), boron (B): 0.003% or less (excluding 0%), aluminum (sol. Al): 0.02-0.05%, phosphorus (P): 0.05% or less (excluding 0%), sulfur (S): 0.01% or less (excluding 0%), and nitrogen (N): 0.01% or less (excluding 0%), and consisting of iron (Fe) and other unavoidable impurities. The steps include heating the steel slab to a temperature range of 1100 to 1300°C, The steps include: hot rolling the heated steel slab to produce a hot-rolled steel sheet; The steps include winding the hot-rolled steel sheet at a temperature range of 400 to 700°C, After the aforementioned winding, the hot-rolled steel sheet is cooled to room temperature. The steps include: cold rolling the cooled hot-rolled steel sheet at a cold reduction ratio of 55-80% to produce a cold-rolled steel sheet; The steps include: continuously annealing the cold-rolled steel sheet, The first step is to perform primary cooling after continuous annealing to a temperature range of 650 to 700°C at an average cooling rate of 1 to 10°C / s, The step includes a secondary cooling step after the primary cooling, in which the temperature is reduced to a range of 450 to 500°C at an average cooling rate of 5 to 50°C / s. The method for manufacturing a high-strength thick steel sheet with excellent hole-expanding properties and ductility, as described in claim 1, wherein the continuous annealing is performed in a facility equipped with a heating zone, a soaking zone and a cooling zone, and when the cold-rolled steel sheet is heated to the heating zone, it passes through a recrystallization zone where it is maintained at 600 to 700°C for 1 to 3 minutes.

8. The method for manufacturing a high-strength thick steel sheet with excellent hole-expanding properties and ductility, as described in claim 7, wherein the hot rolling is performed as a finish hot rolling at an exit temperature of Ar3 or higher and 1000°C or lower.

9. The method for manufacturing a high-strength thick steel sheet with excellent hole-expanding properties and ductility, according to claim 7, wherein the cooling after winding is performed at a cooling rate of 0.1°C / s or less (excluding 0°C / s).

10. The method for manufacturing a high-strength thick steel sheet with excellent hole-expanding properties and ductility, according to claim 7, wherein the heating zone and the uniform zone are controlled to a temperature range of 790 to 850°C.

11. The process further includes a step of performing overaging treatment after the aforementioned secondary cooling, The method for manufacturing a high-strength thick steel plate with excellent hole-expanding properties and ductility, according to claim 7, wherein the over-aging treatment is performed for 200 to 800 seconds.

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