Coated steel sheet and method for manufacturing the same

Abstract

The present invention relates to a steel sheet comprising, in weight percent, carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, and boron (B): greater than 0 and less than or equal to 0.003%, and the remainder comprising iron (Fe) and other unavoidable impurities, wherein the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0; A plated steel plate comprising a plating layer on the base steel plate and a plated layer thereon; wherein the deformation energy from the point of highest stress to the point of fracture in a tensile test on the plated steel plate is 30% or more of the deformation energy from the point of highest stress to the point of fracture in a tensile test on the base steel plate.

Description

Technology Field

[0001] The present invention relates to a steel plate and a method for manufacturing the same, and more specifically, to a plated steel plate and a method for manufacturing the same. Background Technology

[0002] Automotive steel sheets have been developed with a focus on increasing strength to ensure user safety and reduce vehicle body weight, as well as securing elongation to facilitate processing. Currently, common ultra-high strength steels include dual-phase steel (DP steel), which secures elongation through two phases of ferrite and martensite, and transformation-induced plasticity steel (TRIP steel), which secures strength and elongation through the phase transformation of residual austenite in the final structure during plastic deformation by controlling the phase transformation of austenite during heat treatment. However, development based on DP steel, which cannot overcome the limitations of the rule of mixture (ROM), and TRIP steel, which has relatively lower strength as it consists of a ferrite-based matrix with some bainite and tempered martensite, has reached a limit. Therefore, there is a need to develop next-generation ultra-high-strength automotive steel sheets that secure the ultra-high strength and high formability required by customers by improving the microstructure of the transformation-induced plastic steel. Prior art literature

[0003] Korean Patent Publication No. 20200075949A The problem to be solved

[0004] The problem that the present invention aims to solve is to provide an ultra-high-strength galvanized steel sheet with excellent formability and weldability while overcoming the limitations of the mechanical properties of existing DP steel or TRIP steel, and a method for manufacturing the same.

[0005] However, these tasks are exemplary, and the technical concept of the present invention is not limited thereto. means of solving the problem

[0006] A plated steel sheet according to one aspect of the present invention is provided. The plated steel sheet comprises, in weight percent, carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, and boron (B): greater than 0 and less than or equal to 0.003%, and the remainder being iron (Fe) and other unavoidable impurities, wherein the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0. and a plating layer on the base steel plate; wherein the strain energy from the point of highest stress to the point of fracture in a tensile test on the plated steel plate is 30% or more of the strain energy from the point of highest stress to the point of fracture in a tensile test on the base steel plate.

[0007] Immediately after spot welding with respect to the plated steel plate, in an area within 0.1 μm in the direction of the base steel plate at the interface between the base steel plate and the plating layer, the volume fraction of the alloy image gamma (Γ) phase may be 50% or more.

[0008] In the above-mentioned plated steel sheet, the plating layer is a zinc plating layer, and immediately after spot welding with respect to the above-mentioned plated steel sheet, in an area within 0.1 μm in the direction of the base steel sheet at the interface between the base steel sheet and the plating layer, the ratio of the volume fraction of the alloyed gamma (Γ) phase to the volume fraction of the liquid pure zinc (Zn) phase may be 1.5 or more.

[0009] The above galvanized steel sheet may have a yield strength of 850 to 1070 MPa, a tensile strength of 1180 MPa or more, and a total elongation of 14% or more.

[0010] In the above-mentioned plated steel sheet, the final microstructure of the base steel sheet may be, in terms of area fraction, ferrite: 5 to 25%, retained austenite: 5 to 20%, total of tempered martensite and bainite: 30 to 90%, and the remainder may be at least one of pearlite and fresh martensite.

[0011] The final microstructure of the base steel sheet in the above-mentioned plated steel sheet can satisfy the following Equation 1.

[0012] (Formula 1)

[0013] (X Pearlite + X Fresh Martensite ) / (100 - X Ferrite - X Tempered Martensite - X Bainite ) ≤ 0.30

[0014] (Above X Pearlite is the value of the area fraction of pearlite (unit: %), and the above X Fresh Martensite is the value of the area fraction of fresh martensite (unit: %), and the above X Ferrite is the value of the area fraction of ferrite (unit: %), and X Tempered Martensite is the value of the area fraction of tempered martensite (unit: %), and the above X Bainite is the value of the area fraction of bainite (unit: %)

[0015] A method for manufacturing a plated steel sheet according to one aspect of the present invention is provided. The method for manufacturing the plated steel sheet comprises the step of providing a hot-rolled steel sheet by hot-rolling a steel material characterized in that, in weight percent, it contains carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, and boron (B): greater than 0 and less than or equal to 0.003%, and the remainder contains iron (Fe) and other unavoidable impurities, wherein the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0. The method comprises the steps of: cold rolling the hot-rolled steel sheet to provide a cold-rolled steel sheet; annealing the cold-rolled steel sheet in an abnormal temperature range; slowly cooling the annealed steel sheet at a first cooling rate; rapidly cooling the slowly cooled steel sheet to a temperature below the martensite transformation start temperature (Ms) at a second cooling rate greater than the first cooling rate; reheating the rapidly cooled steel sheet in a temperature range above the martensite transformation start temperature (Ms) and below the bainite transformation start temperature (Bs); performing a plating treatment on the reheated steel sheet as a base steel sheet to form a plating layer on the base steel sheet; and an alloying heat treatment step of maintaining the base steel sheet and the plating layer in a temperature range of 480 to 560°C and then cooling to room temperature.

[0016] In the above method for manufacturing the galvanized steel sheet, the hot rolling can be performed under conditions of reheating temperature: 1150 ~ 1250℃, finishing rolling temperature: 800 ~ 950℃, and coiling temperature: 400 ~ 650℃.

[0017] In the method for manufacturing the above-described plated steel sheet, the annealing heat treatment step includes a step of maintaining the cold-rolled steel sheet in a temperature range of 800 to 850°C for 60 to 130 seconds, the first cooling rate is lower than 20°C / s, the second cooling rate is higher than 20°C / s, the slow cooling end temperature of the slow cooling step is 650 to 750°C, and the rapid cooling end temperature of the rapid cooling step may be 200 to 300°C.

[0018] In the method for manufacturing the above-mentioned plated steel sheet, the alloying heat treatment step includes a step of maintaining at a temperature range of 480 to 560°C for 15 to 45 seconds and then cooling to room temperature at a third cooling rate, wherein the third cooling rate may be greater than the first cooling rate and smaller than the second cooling rate. Effects of the invention

[0019] According to the present invention, it is possible to realize an ultra-high-strength plated steel sheet with excellent formability and weldability while overcoming the limitations of the mechanical properties of conventional DP steel or TRIP steel, and a method for manufacturing the same.

[0020] The effects of the present invention described above are illustrative and the scope of the present invention is not limited by these effects. Brief explanation of the drawing

[0021] FIG. 1 is a flowchart illustrating a method for manufacturing a plated steel sheet according to one embodiment of the present invention. Figure 2 is a graph showing the results of the Glibl tensile test for the plated steel sheet and base steel sheet of the present invention. FIG. 3 is a diagram illustrating the simulation results of the interface movement between the plating layer and the substrate steel sheet and the aluminum diffusion pattern immediately after spot welding in a plated steel sheet according to one embodiment of the present invention. FIG. 4 is a diagram illustrating the change in stability of the interface phase due to aluminum diffusion in a plated steel sheet according to one embodiment of the present invention. FIGS. 5 to 7 are photographs of the microstructure of the substrate steel sheet among the plated steel sheets according to Comparative Examples 1 to 3, respectively. FIGS. 8 to 10 are photographs of the microstructure of the substrate steel sheet among the plated steel sheets according to Invention Examples 7 to 9. Specific details for implementing the invention

[0022] Hereinafter, preferred embodiments of the present invention will be described in detail with reference to the attached drawings. The embodiments of the present invention are provided to more completely explain the technical concept of the present invention to those skilled in the art, and the following embodiments may be modified in various different forms, and the scope of the technical concept of the present invention is not limited to the following embodiments. Rather, these embodiments are provided to make the present disclosure more faithful and complete and to fully convey the technical concept of the present invention to those skilled in the art. In this specification, the same reference numerals denote the same elements throughout. Furthermore, various elements and areas in the drawings are depicted schematically. Accordingly, the technical concept of the present invention is not limited by the relative sizes or spacing depicted in the attached drawings.

[0023] Recently, the automotive industry has seen increasing interest in lightweight vehicle bodies utilizing ultra-high-strength steel to simultaneously satisfy crash safety and fuel efficiency regulations. In response to these demands from automotive companies, the steel industry is actively conducting research on the development of ultra-high-strength steel. Q&P (Quenching and Partitioning) heat treatment technology, developed to simultaneously secure high strength and ductility in automotive steel materials, is a technique that suppresses the formation of carbon carbide precipitates released from the martensite structure during quenching and facilitates the diffusion of carbon into the retained austenite structure through partitioning. By re-diffusion of carbon, the retained austenite structure is stabilized even at room temperature, ultimately enabling the achievement of high ductility from the retained austenite structure and high strength from the martensite structure. These Q&P steel sheets contain a large amount of silicon (Si) compared to ordinary steel to inhibit the movement of iron (Fe) atoms in order to suppress the formation of carbide precipitates within the structure, and also contain a large amount of austenite-stabilizing alloying elements such as carbon (C) and manganese (Mn) to increase the volume fraction of stabilized residual austenite structure and improve TRIP (Transformation-Induced Plasticity) behavior.

[0024] Meanwhile, although it is obvious that the technical concept of the present invention can be applied to the Q&P steel plate described above, it is evident that the technical concept of the present invention can be widely applied to various steel plates and is not limited to the application of the Q&P steel plate alone.

[0025] In order to overcome the limitations of the mechanical properties of existing DP steel and TRIP steel, the present invention aims to develop a high-formability, ultra-high-strength galvanized steel sheet for automobiles that can secure high strength and appropriate elongation by replacing the main matrix of TRIP steel with tempered martensite, bainite, and inevitably formed fresh martensite instead of ferrite.

[0026] Specifically, the present invention reduces the ratio of austenite phase transformation, which may inevitably occur during the heat treatment of QP steel to TRIP steel plated materials, through composition control and heat treatment control. More specifically, the composition system is configured by controlling the total amount and ratio of silicon (Si) and aluminum (Al) alloys, and carbon diffusion into austenite and auto-tempering at low temperatures are additionally induced compared to conventional heat treatment processes by maintaining the material in the section prior to reheating after rapid cooling in the manufacturing process. This controls the ratio of unnecessary austenite phase transformation during reheating and alloying heat treatment, thereby describing a plated steel sheet having a target material and a method for manufacturing the same.

[0027] FIG. 1 is a flowchart illustrating a method for manufacturing a plated steel sheet according to one embodiment of the present invention.

[0028] Referring to FIG. 1, a method for manufacturing a plated steel sheet according to an embodiment of the present invention sequentially comprises the steps of: providing a hot-rolled steel sheet (S10); cold-rolling the hot-rolled steel sheet to provide a cold-rolled steel sheet (S20); annealing the cold-rolled steel sheet (S30); slowly cooling the annealed steel sheet at a first cooling rate (S40); rapidly cooling the slowly cooled steel sheet at a second cooling rate greater than the first cooling rate (S50); reheating the rapidly cooled steel sheet at a predetermined temperature range (S60); performing a plating treatment on the reheated steel sheet as a base steel sheet to form a plating layer on the base steel sheet (S70); and an alloying heat treatment step (S80) of maintaining the base steel sheet and the plating layer at a predetermined temperature range and then cooling to room temperature. Below, a method for manufacturing an ultra-high-strength plated steel sheet with excellent formability and weldability according to an embodiment of the present invention will be examined in detail.

[0029] hot-rolled steel sheet supply step (S10)

[0030] A plated steel sheet according to one embodiment of the present invention consists of a base steel sheet (substrate) and a plating layer formed on the base steel sheet. Hereinafter, the role and content of exemplary components included in the base steel sheet are described as follows. In this case, the content of all component elements refers to weight percent.

[0031] Carbon (C): 0.1~0.3%

[0032] Carbon is the most important alloying element in steelmaking, and its primary purpose is to perform basic strengthening and austenite stabilization. A high carbon concentration in the austenite improves austenite stability, making it easier to secure an appropriate amount of austenite for material improvement. If the carbon content is less than 0.1%, it is difficult to secure the desired yield strength and elongation. If the carbon content exceeds 0.3%, or strictly speaking, 0.25%, it may lead to a decrease in weldability due to the increase in carbon equivalent. Therefore, it is desirable for the carbon content to be 0.1% to 0.3% of the total weight of the base steel sheet, and strictly, 0.1% to 0.25%.

[0033] Manganese (Mn): 2.0~4.0%

[0034] Manganese is an austenite stabilizing element, and as manganese is added, the martensite transformation start temperature Ms gradually decreases, which can lead to an increase in the residual austenite fraction during the continuous annealing heat treatment process. If the manganese content is less than 2.0%, or strictly less than 2.5%, the effect of manganese addition is insufficient. If the manganese content exceeds 4.0%, the carbon equivalent increases, significantly reducing weldability, and oxides (MnO) may form on the surface of the steel sheet during the process, leading to a decrease in plating performance due to poor wettability in that area. Therefore, it is desirable that the manganese content be 2.0% or more and 4.0% or less of the total weight of the base steel sheet, and strictly, 2.5% or more and 4.0% or less.

[0035] Silicon (Si): Greater than 0% and less than or equal to 2.0%

[0036] Silicon is an element that inhibits the formation of carbides (e.g., Fe3C) within ferrite and increases the activity of carbon, thereby increasing the diffusion rate of austenite. Silicon is also well known as a ferrite-stabilizing element, and is known to increase ductility by increasing the ferrite fraction during cooling. If the silicon content exceeds 2.0%, oxides (SiO2) may form on the surface of the base steel sheet during processing, which may lead to a decrease in plating performance due to poor wettability in that area. Additionally, Liquid Metal Embrittlement (LME) may occur, a phenomenon in which the plating layer melts during Resistance Spot Welding of the plated steel sheet, causing molten zinc to penetrate toward the interface of residual austenite present on the surface of the base steel sheet and induce brittleness. According to one embodiment of the present invention, since the base steel sheet contains silicon, it is preferable that the silicon content be greater than 0 and less than or equal to 2.0% of the total weight of the base steel sheet. However, according to another embodiment of the present invention, the base steel sheet may not contain silicon. Accordingly, in the present invention, the silicon content may be 0% or more and 2.0% or less of the total weight of the base steel sheet. However, even if actual silicon (Si) is not intentionally added, a very small amount may be detected.

[0037] Aluminum (Al): Greater than 0 and less than or equal to 2.0%

[0038] Aluminum plays a role in reducing the austenite transformation fraction due to temperature changes during the heating process in annealing. The addition of aluminum can reduce material variation by decreasing changes in the fraction of abnormal structures during the heating process.

[0039] Meanwhile, next-generation steel sheets such as Q&P steel sheets may contain a higher amount of silicon (Si) compared to general steel to inhibit the movement of iron (Fe) atoms and suppress the formation of carbide precipitates within the microstructure; however, since a high silicon content can make the material sensitive to liquid metal embrittlement (LME), the silicon content may be reduced. In this case, aluminum may be added to suppress phase transformation during reheating and alloying heat treatment. That is, suppressing the phase transformation of residual austenite during alloying heat treatment after the formation of the plating layer in the manufacturing process of plated steel sheets is effective in terms of securing elongation. Therefore, this invention proposes an alloy design that achieves a high aluminum content and a low silicon content.

[0040] However, if the aluminum content exceeds 2.0%, the annealing temperature required to secure an ideal structure may increase excessively, which may reduce mass production efficiency, and surface defects such as dents may increase due to the formation of oxide foreign matter during annealing. Additionally, it may cause an increase in steelmaking inclusions and surface oxidation during annealing. Therefore, it is desirable that the aluminum content be greater than 0% and less than or equal to 2.0% of the total weight of the base steel sheet.

[0041] Furthermore, the aluminum content may strictly be 0.5% or more and 2.0% or less of the total weight of the base steel sheet, and more strictly, may exceed 0.5% and be 2.0% or less of the total weight of the base steel sheet. In this case, the carbides formed during the alloying process after plating are small, and the chemical stability within the austenite is improved. Therefore, during the final cooling process after the alloying heat treatment, the phase transformation from austenite to pearlite or fresh martensite is relatively small, allowing the reduction rate of residual austenite to be managed at a low level. In addition, immediately after spot welding with respect to the plated steel sheet, in the region within 0.1 μm in the direction of the base steel sheet at the interface between the base steel sheet and the plating layer, the volume fraction of the alloying phase gamma (Γ) phase is 50% or more, and the ratio of the volume fraction of the alloying phase gamma (Γ) phase to the volume fraction of the liquid phase pure zinc (Zn) phase can be realized to be 1.5 or more. In addition, when conducting a Glib test, the deformation energy from the point where maximum tensile strength (UTS) occurs to the point where fracture occurs can be reduced by an energy reduction rate of 70% or less for the plated steel sheet compared to the unplated base steel sheet. That is, it can be confirmed that the deformation energy from the point of highest stress to the point of fracture in a tensile test on the plated steel sheet is at least 30% of the deformation energy from the point of highest stress to the point of fracture in a tensile test on the base steel sheet. The liquid metal embrittlement (LME) characteristics of QP steel or TRIP steel plating materials can be improved through the strict control of the aluminum (Al) alloy content described above. Depending on the strict control of the aluminum (Al) alloy content described above, the amount of aluminum (Al) diffusion to the interface between the plating layer and the steel sheet changes, and consequently, resistance to liquid metal embrittlement (LME) can be secured through changes in phase stability in the interface region.That is, for galvanized steel sheets, the stability of the Γ phase increases due to changes in the interface composition and phase stability caused by Al interface diffusion, and the fraction of liquid phase Zn at the interface decreases due to the decrease in the stability of liquid phase Zn. Consequently, the amount of penetration into the substrate steel sheet is reduced due to the reduction of liquid phase Zn at the interface, thereby ensuring resistance to liquid metal embrittlement (LME).

[0042] Sum of silicon (Si) and aluminum (Al) content: 2.0% or less, Ratio of silicon (Si) content to aluminum (Al) content: less than 4.0

[0043] In the present invention, the ratio of austenite phase transformation, which may inevitably occur during heat treatment of QP steel or TRIP steel plated materials, is reduced through compositional control. For example, by configuring the compositional system through the total amount and ratio control of silicon (Si) and aluminum (Al) alloying amounts, carbon diffusion into austenite is additionally induced during auto-tempering at low temperatures and austenite, thereby controlling the ratio of unnecessary austenite phase transformation during reheating and alloying heat treatment, and thus enabling the realization of a plated steel sheet having a target material.

[0044] Under conditions where the sum of the silicon (Si) content and the aluminum (Al) content in the base steel sheet is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0, the ratio of unnecessary austenite phase transformation (austenite → pearlite, fresh martensite) during reheating and alloying heat treatment can be controlled to 30% or less.

[0045] In contrast, when the sum of the silicon (Si) and aluminum (Al) content exceeds 2.0%, liquid metal embrittlement (LME), a phenomenon in which the plating layer melts during resistance spot welding of the plated steel sheet and molten zinc penetrates toward the interface of residual austenite present on the surface of the base steel sheet to cause brittleness, may occur significantly, and the annealing temperature required to secure an abnormal structure may increase excessively, which may reduce mass production efficiency, and surface defects such as dents may increase as oxidative foreign matter is formed during annealing, and may cause an increase in steelmaking inclusions and surface oxidation during annealing.

[0046] Meanwhile, even under conditions where the sum of the silicon (Si) content and aluminum (Al) content in the base steel sheet is 2.0% or less, if the ratio of the silicon (Si) content to the aluminum (Al) content is 4.0 or more, the ratio of unnecessary austenite phase transformation (austenite → pearlite, fresh martensite) during reheating and alloying heat treatment cannot be suppressed to 30% or less, so a problem may occur in which the total elongation is achieved at less than 14%.

[0047] Phosphorus (P): Greater than 0% ~ 0.02% or less

[0048] Phosphorus can play a role similar to silicon in steel. However, if phosphorus is added in an amount exceeding 0.02% of the total weight of the steel sheet, it can reduce the weldability of the steel sheet and increase brittleness, thereby causing material degradation. Therefore, it is desirable to limit the phosphorus content to 0.02% or less of the total weight of the base steel sheet.

[0049] Sulfur (S): Greater than 0% ~ 0.005% or less

[0050] Sulfur is an element that is inevitably contained during the manufacture of steel, and it impairs the toughness and weldability of steel and reduces the corrosion resistance and impact properties of steel by combining with manganese (Mn) to form MnS. Therefore, it is desirable to limit the sulfur content to 0.005% or less of the total weight of the base steel sheet.

[0051] Nitrogen (N): Greater than 0, 0.006% or less

[0052] Nitrogen is an element that is inevitably contained during the manufacture of steel and is an element that degrades aging resistance, so it is desirable to reduce it as much as possible. Therefore, it is desirable to limit the nitrogen content to 0.006% or less of the total weight of the base steel sheet.

[0053] Boron (B): Greater than 0 and less than or equal to 0.003%

[0054] Boron is a hardenable element in steel that can segregate at grain boundaries during cooling to suppress the formation of ferrite and function as a grain boundary strengthening element. However, if the boron content exceeds 0.003%, the strength may increase excessively, and high-temperature ductility may decrease by forming nitrides such as BN, and in the case of plated materials, plating peeling may occur. Therefore, it is desirable to add boron in an amount greater than 0% and less than or equal to 0.003% of the total weight of the base steel sheet.

[0055] The remaining component of the base steel sheet, which is the base material of the above-mentioned galvanized steel sheet, is iron (Fe). However, since unintended impurities from raw materials or the surrounding environment may inevitably be incorporated during the conventional steelmaking process, they cannot be excluded. As these impurities are known to any skilled technician in the conventional manufacturing process, all details thereof are not specifically mentioned in this specification.

[0056] In the manufacturing method according to the present invention, the semi-finished product subject to the hot rolling and cold rolling processes may be, for example, a slab. The slab in the semi-finished product state can be obtained through a continuous casting process after obtaining molten steel of a predetermined composition through a steelmaking process.

[0057] A step is performed to form a hot-rolled steel sheet by applying a hot rolling process to the above steel material. Since the above steel material is a high-alloy steel, it is necessary to minimize edge cracking and rolling load to ensure mass production, so the rolling finish temperature and coiling temperature can be set to a high temperature range.

[0058] The above steel is reheated, for example, at a slab reheating temperature (SRT) in the range of 1150°C to 1250°C. Through this reheating, the resolution of segregated components and precipitates may occur. If the reheating temperature is below 1150°C, a problem may arise where the hot rolling load increases rapidly. If the reheating temperature exceeds 1250°C, charging and discharging from the furnace may be difficult due to slab bending, and it may be difficult to secure the strength of the final product steel sheet due to the coarsening of the initial austenite grains. The reheating temperature may vary depending on the steel.

[0059] Subsequently, the reheated steel is hot-rolled, and hot-rolling can be performed at a Finish Delivery Temperature (FDT) of, for example, 800°C to 950°C. If the Finish Delivery Temperature exceeds 950°C, there is a risk that the quality of the steel sheet will deteriorate due to the formation of surface scale on the steel sheet. In addition, if the Finish Delivery Temperature is less than 800°C, it may cause an increase in rolling load and a decrease in productivity. The Finish Delivery Temperature may vary depending on the steel.

[0060] Next, the hot-rolled steel is cooled at a cooling rate of 10 to 30°C / s and then coiled at a coiling temperature (CT) in the range of, for example, 400°C to 650°C. The coiling temperature may vary depending on the steel. If the coiling temperature exceeds 650°C, an undesirable internal oxide layer may form on the hot-rolled steel sheet or the coiled hot-rolled coil. Since the internal oxidation of the coiled hot-rolled coil varies, it may be difficult to uniformly control the thickness of the internal oxide layer. If the coiling temperature is below 400°C, an undesirable low-temperature structure may be formed.

[0061] Softening heat treatment and pickling step (S15)

[0062] Meanwhile, in the method for manufacturing a plated steel sheet according to one embodiment of the present invention, a softening heat treatment step and a pickling step may be performed sequentially after performing a hot rolling process and before performing a cold rolling process. In the softening heat treatment step, the hot-rolled steel sheet is softened to achieve material softening, thereby mitigating problems with reduction rate load and shape defects during subsequent cold rolling. That is, to improve cold rolling efficiency, the hot-rolled steel sheet can be softened through softening heat treatment to ensure cold rolling performance. If the strength of the hot-rolled steel sheet is high, problems such as thickness hunting and shape defects may occur during cold rolling. However, since the softening heat treatment process for ultra-high strength steel is applied to hot-rolled coils with remaining scale rather than cold-rolled coils, countermeasures are required to address changes in surface characteristics due to the high-temperature reaction of scale during softening heat treatment. Generally, it is known that in the case of ultra-high strength steel materials containing large amounts of Si and Mn, internal oxides are generated along the grain boundaries of the steel sheet along with scale at high temperatures. The oxide layer generated by internal oxidation exhibits inferior pickling performance because the main component of the matrix is ​​Fe. Consequently, the internal oxide layer cannot be completely removed within the pickling time equivalent to that of general hot-rolled steel sheets, and the long pickling time required leads to a problem of reduced productivity. Such internal oxidation occurs when the activity of easily oxidizable elements, such as Si and Mn, is high and they exist under specific oxygen partial pressure conditions. Therefore, when hot-rolled coils containing remaining scale are heat-treated in a high-temperature reducing gas atmosphere, additional internal oxidation occurs due to oxygen generated during the scale reduction reaction. Specifically, internal oxides, which were not observed in the hot-rolled material, developed unevenly across the entire length of the coil after softening heat treatment, and the behavior of scale reduction and internal oxidation development may vary depending on the location within the coil.In the case of the outer coil of the softening heat treatment coil, internal oxidation growth due to the scale hydrogen reduction reaction was mainly observed, while in the case of the inner coil, internal oxidation growth due to the scale eutectoid reaction (4FeO → 4Fe + 2O2) was observed. It is determined that this difference in internal oxidation growth behavior is due to the difference in the ease of penetration of reaction gases depending on the position of the coil being wound, and it was understood that oxygen generated during the scale hydrogen reduction and eutectoid reactions during the softening heat treatment diffuses into the base material and acts as an internal oxidation reaction material. Since it is desirable for the internal oxide layer to be formed as uniformly as possible across the entire steel sheet, it is desirable to suppress the internal oxide layer during the winding stage as much as possible and to form the internal oxide layer during the softening heat treatment.

[0063] Considering these aspects, in the method for manufacturing a plated steel sheet according to the present invention, the coiling temperature is controlled to 400°C to 650°C and the softening heat treatment temperature is controlled to 450°C to 650°C. The softening heat treatment can be performed in a batch annealing furnace (BAF) while the hot-rolled steel sheet is coiled, and can be performed in a hydrogen atmosphere. The hot-rolled steel sheet that has undergone softening heat treatment under the above process conditions has its material softened, thereby ensuring cold rolling properties. In addition, the internal oxide layer formed by the softening heat treatment has a thickness of less than a predetermined thickness (e.g., 10 μm or less), thereby simultaneously ensuring subsequent pickling properties.

[0064] When a softening heat treatment is applied at a temperature below 450°C, the martensite formed after hot rolling does not undergo recrystallization, and only tempering proceeds, causing supersaturated carbon within the structure to form cementite (θ) and spheroidization to occur. In this case, the brittleness of the martensite may be exhibited, which may lead to safety accidents such as plate breakage during cold rolling. In other words, when the softening heat treatment is performed, excessive austenite is formed, and martensite is formed during cooling, which may prevent the effective expression of strength reduction.

[0065] In addition, when the softening heat treatment is applied at a temperature exceeding 650°C, the internal oxide layer formed by the softening heat treatment exceeds a predetermined thickness (e.g., exceeding a thickness of 10 μm), making it difficult to ensure subsequent pickling performance. Furthermore, when the softening heat treatment is applied at a temperature exceeding 650°C, austenite is excessively formed, and as martensite is formed during cooling, the reduction in strength is not effectively manifested.

[0066] Meanwhile, in the method for manufacturing a plated steel sheet with excellent weldability according to one embodiment of the present invention, whether or not to apply the softening heat treatment step may be selective depending on the steel grade or target strength. For example, if the target tensile strength after cold rolling / annealing is 1180 MPa or higher, the softening heat treatment step may be performed, and if the target tensile strength after cold rolling / annealing is 980 MPa, the softening heat treatment step may not be performed. The time for performing the softening heat treatment may be 1 to 12 hours.

[0067] In the pickling step, after performing the softening heat treatment, a pickling treatment can be performed to clean the hot-rolled steel sheet with acid. By pickling the hot-rolled steel sheet, at least a portion of the internal oxide layer can be removed. The pickling treatment can be performed, for example, at a temperature of 70°C to 90°C, for example, with a hydrochloric acid concentration of 5% to 15%, for example, 20 seconds to 40 seconds. Additionally, it can have an inhibitor concentration of 0.1% to 0.5%.

[0068] Cold rolling step (S20)

[0069] A step (S20) is performed to form a cold-rolled steel sheet by applying a cold rolling process to the above hot-rolled steel sheet. In the case of cold rolling, it is performed using hot-rolled material to match the thickness of the final produced steel sheet.

[0070] In the cold rolling step (S20), the pickled hot-rolled steel sheet can be cold-rolled with, for example, an average reduction rate of 40% to 60%, and a cold-rolled steel sheet can be manufactured accordingly. The microstructure of the cold-rolled steel sheet has an elongated shape compared to the microstructure of the hot-rolled steel sheet, and the microstructure of the steel sheet produced in the subsequent heat treatment is determined.

[0071] Annealing heat treatment step (S30)

[0072] A cold-rolled steel sheet can be heat-treated by annealing in a continuous annealing furnace with a slow cooling section. The step of heat-treating by annealing (S30) may include a step of maintaining the cold-rolled steel sheet at a temperature range of 800 to 850°C for 60 to 130 seconds. That is, the annealing heat treatment temperature may be set in an abnormal temperature range where austenite and ferrite coexist.

[0073] The above annealing heat treatment is performed to form an austenite structure. The annealing heat treatment temperature and time affect the austenite grain size and, therefore, can have a significant impact on the strength of the cold-rolled steel sheet.

[0074] The above annealing heat treatment is performed by heating at a heating rate of, for example, 1°C / s or more, or at a heating rate in the range of, for example, 1°C / s to 10°C / s. If the heating rate is less than 1°C / s, it takes a long time to reach the target annealing heat treatment temperature, which reduces production efficiency and may increase the size of the crystal grains.

[0075] The above annealing heat treatment can be performed, for example, by maintaining the temperature in the austenite and ferrite phase region of Ae1 to Ae3, for example, in the range of 800°C to 850°C, for a time in the range of 60 seconds to 130 seconds. During this heating and annealing heat treatment step, the cold-rolled structure undergoes a reverse transformation into austenite. If the annealing heat treatment temperature is below 800°C, austenite cannot be formed to create the final structure, tempered martensite. If the annealing heat treatment temperature exceeds 850°C, the austenite grains become coarse, which may result in a decrease in strength.

[0076] As the annealing heat treatment time increases, it affects coarsening due to austenite grain growth, similar to the annealing heat treatment temperature. If the annealing heat treatment time exceeds 130 seconds, the heat treatment efficiency may decrease. If the annealing heat treatment time is less than 60 seconds, the annealing heat treatment effect may be insufficient. Therefore, the annealing heat treatment time ranges from 60 seconds to 130 seconds.

[0077] In the annealing heat treatment step, the dew point temperature of the atmosphere gas inside the annealing furnace is controlled to be between -10°C and +20°C. The atmosphere gas inside the annealing furnace may be, for example, a mixed gas of nitrogen: 90% and oxygen: 10%. By using a humidifier attached to the outer wall of the annealing furnace to spray pure H2O into the furnace, the dew point of the heat treatment atmosphere gas can be managed to be between -10°C and +20°C. Through annealing heat treatment having the above-described dew point temperature range, a decarburization reaction may occur on the surface of the base steel sheet constituting the plated steel sheet of the present invention.

[0078] Multi-stage cooling steps (S40, S50)

[0079] The annealed cold-rolled steel sheet is cooled in multiple stages. The cooling step may be performed in the following two steps. That is, the multi-stage cooling includes a step of slow cooling the annealed steel sheet at a first cooling rate (S40); and a step of rapidly cooling the slow-cooled steel sheet to a temperature below the martensite transformation start temperature (Ms) at a second cooling rate greater than the first cooling rate (S50). The first cooling rate may be lower than 20℃ / s, and the second cooling rate may be higher than 20℃ / s. The slow cooling end temperature of the slow cooling step (S40) may be 650 to 750℃, and the rapid cooling end temperature of the rapid cooling step (S50) may be 200 to 300℃. The rapid cooling step (S50) is characterized by maintaining the temperature at the rapid cooling end temperature for 20 to 60 seconds after reaching the rapid cooling end temperature. The above multi-stage cooling steps are explained in detail below.

[0080] First, the annealed cold-rolled steel sheet is first cooled by slow cooling at a first cooling rate of, for example, 3°C / s or more and less than 20°C / s, to a temperature range that suppresses ferrite transformation, for example, to a first cooling end temperature in the range of 650°C to 750°C. By controlling the cooling rate from the annealing temperature, the formation of ferrite with a high manganese (Mn) concentration can be suppressed, and the amount of ferrite formed can be suppressed. If the average cooling rate is slower than 3°C / s, ferrite with a high manganese (Mn) concentration may be formed during cooling, which may degrade bendability or reduce strength.

[0081] Next, a second cooling step is performed in which the first cooled cold-rolled steel sheet is rapidly cooled to a second cooling end temperature below the martensite transformation start temperature (Ms) at a second cooling rate of, for example, greater than 20℃ / s and less than or equal to 100℃ / s. The second cooling end temperature is a temperature below the Ms temperature, and may, for example, have a range of 200℃ to 300℃.

[0082] In the above second cooling step, cooling is performed at a rapid cooling rate, so that the transformation of ferrite, pearlite, or bainite is suppressed, and some of the austenite can transform into martensite.

[0083] In the above second cooling step, if the cooling rate is less than 20℃ / s, the fraction of austenite after cooling is too high, making it difficult to ensure the stability of the retained austenite, and even if the bainite transformation structure is increased, the martensite fraction is low, which may result in a decrease in strength. In addition, some martensite structures increase internal stress, thereby increasing the bainite nucleation rate and causing the bainite transformation to proceed rapidly even at low temperatures below the martensite transformation initiation temperature (Ms).

[0084] The above second-cooled cold-rolled steel sheet can be maintained at the second-cooling end temperature for a predetermined time, for example, from 20 seconds to 60 seconds. During the second-cooling end temperature maintenance period, a portion of the austenite transforms into martensite, and precipitates such as metal carbides may be formed within the martensite generated during the second cooling stage. If precipitates are excessively formed within the martensite during the second-cooling end temperature maintenance period, the elongation of the final product may exhibit a relatively low value. Therefore, the second-cooling end temperature maintenance time must be performed within a range that suppresses the formation of excessive precipitates; accordingly, in the present invention, the maintenance time is from 20 seconds to 60 seconds.

[0085] In the present invention, the ratio of austenite phase transformation, which may inevitably occur during heat treatment of QP steel or TRIP steel plated material, is reduced by configuring the composition system through the total amount and ratio control of silicon (Si) and aluminum (Al) alloy amounts, and then maintaining the temperature in the section prior to reheating after the rapid cooling step (S50), which is the second cooling described above. For example, by securing a section of the second cooling end temperature maintenance for 20 to 60 seconds after the rapid cooling step (S50), which is the second cooling described above, auto-tempering at low temperatures and carbon diffusion into austenite are additionally induced, thereby controlling the ratio of unnecessary austenite phase transformation during reheating and alloying heat treatment, and thus enabling the realization of a plated steel sheet having the target material.

[0086] If the second cooling end temperature holding period after the aforementioned second cooling rapid cooling step (S50) is less than 20 seconds, the ratio of unnecessary austenite phase transformation (austenite → pearlite, fresh martensite) during reheating and alloying heat treatment cannot be suppressed to 30% or less, so the total elongation is achieved at less than 14%. If the second cooling end temperature holding period after the aforementioned second cooling rapid cooling step (S50) exceeds 60 seconds, or if metal carbides are excessively precipitated within the martensite, it becomes difficult to secure an elongation of 14% or more of the final product.

[0087] Reheating heat treatment step (S60)

[0088] A reheating step (S60) is performed in which the above-mentioned rapidly cooled steel sheet is maintained in a temperature range above the martensite transformation start temperature (Ms) and below the bainite transformation start temperature (Bs). That is, the above-mentioned multi-stage cooled cold-rolled steel sheet is reheated at a heating rate of, for example, 40℃ / s or more, and a reheating heat treatment is performed by maintaining it at a temperature in the range of, for example, 350℃ to 470℃ for a time in the range of, for example, 20 seconds to 50 seconds.

[0089] In the reheating heat treatment (partitioning heat treatment) step, carbon diffuses into the retained austenite and becomes concentrated, acting to stabilize the retained austenite. Additionally, some austenite may undergo martensitic transformation. This martensitic transformation can refine the shape of the retained austenite after rapid cooling, thereby contributing to the stabilization of the retained austenite. If the partitioning heat treatment temperature is less than 350°C, the partitioning effect described above may be insufficient. If the reheating heat treatment temperature exceeds 470°C, the size of the carbides may coarsen, resulting in a decrease in strength.

[0090] If the holding time for the reheating heat treatment is less than 20 seconds, it may be difficult to obtain a stable partitioning effect. If the holding time for the reheating heat treatment exceeds 50 seconds, the heat treatment efficiency decreases, and the size of the carbides increases, which may result in a decrease in strength.

[0091] Plating treatment step (S70)

[0092] Next, a step (S70) is performed to form a plating layer on the substrate steel plate by performing a plating treatment on the reheated steel plate as a substrate steel plate. The plating layer may be formed by immersing the reheated steel plate in a molten plating bath and performing molten plating. The plating bath may be a molten zinc plating bath, and the entry temperature of the plating bath may be, for example, 460℃. The plating adhesion amount is 40 to 200 g / m² on both sides. 2 The plating layer thickness can be 10 to 30㎛.

[0093] Alloying heat treatment step (S80)

[0094] An alloying heat treatment step (S80) can be performed in which the above-described base steel plate and the above-described plating layer are maintained in a temperature range of 480 to 560°C. In the alloying heat treatment step (S80), the time for maintaining in the temperature range of 480 to 560°C may be 15 to 45 seconds. The plated steel plate realized through the alloying heat treatment after the above-described zinc hot-dip galvanizing is an alloyed hot-dip galvanized steel plate. An Fe-Zn alloy phase can be formed within the plating layer through the alloying heat treatment.

[0095] If the above-mentioned alloying heat treatment temperature is less than 480°C, the above-mentioned alloying heat treatment effect is insufficient, making it difficult to form an Fe-Zn alloy phase within the plating layer. If the above-mentioned alloying heat treatment temperature exceeds 560°C, transformation-induced plasticity is not expressed due to the decomposition of residual austenite along with the tempering effect of the matrix structure as alloying progresses, which may result in a decrease in strength.

[0096] Meanwhile, even if the time for maintaining the alloying heat treatment in the temperature range of 480 to 560°C is less than 15 seconds, the alloying heat treatment effect described above may be insufficient, making it difficult to form an Fe-Zn alloy phase within the plating layer, and if the time for maintaining the alloying heat treatment in the temperature range of 480 to 560°C exceeds 45 seconds, the size of the carbides may coarsen, and a decrease in strength may occur.

[0097] Meanwhile, the alloying heat treatment step (S80) may include a step of maintaining the temperature in the range of 480 to 560°C for 15 to 45 seconds and then cooling to room temperature at a third cooling rate. The third cooling rate may be greater than the first cooling rate of the slow cooling step (S40) described above and smaller than the second cooling rate of the rapid cooling step (S50) described above.

[0098] In the present invention, the material properties (especially elongation) of the steel sheet are secured by suppressing, to the maximum extent, the phenomenon in which austenite undergoes phase transformation into at least one of pearlite and fresh martensite during the alloying heat treatment step (S80). To achieve this, the compositional system is configured by controlling the total amount and ratio of silicon (Si) and aluminum (Al) alloys, and it is described that this can be achieved by maintaining the composition in the section prior to reheating after the rapid cooling step (S50), which is the secondary cooling described above.

[0099] The galvanized steel sheet implemented by the above-described manufacturing method is a base steel sheet comprising, in weight percent, carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, boron (B): greater than 0 and less than or equal to 0.003%, and the remainder being iron (Fe) and other unavoidable impurities; and a galvanized steel sheet comprising a galvanized layer on the base steel sheet; wherein the sum of the silicon (Si) content and the aluminum (Al) content in the base steel sheet is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0.

[0100] The above-mentioned galvanized steel sheet achieves physical properties of yield strength: 850 to 1070 MPa, tensile strength: 1180 MPa or higher, and total elongation: 14% or higher. For example, the above-mentioned galvanized steel sheet achieves physical properties of tensile strength: 1180 MPa or higher and 1300 MPa or lower, and total elongation: 14% or higher and 20% or lower.

[0101] The final microstructure of the base steel sheet among the plated steel sheets is characterized by, in terms of area fraction, ferrite: 5 to 25%, retained austenite: 5 to 20%, total of tempered martensite and bainite: 30 to 90%, and the remainder being at least one of pearlite and fresh martensite.

[0102] If the area fraction of ferrite is less than 5%, the workability of the base steel sheet is inferior, and if it exceeds 25%, it may not be easy to achieve a tensile strength of 1180 MPa or higher of the base steel sheet. Meanwhile, if the area fraction of retained austenite is less than 5%, it is difficult to expect a transformation-induced plasticity effect, and if it exceeds 20%, the proportion of low-temperature structures (tempered martensite and bainite) becomes relatively low, making it difficult to secure the strength of the steel sheet.

[0103] Meanwhile, the pearlite and fresh martensite are the result of the decomposition of residual austenite during the heat treatment of QP steel to TRIP steel plated materials in the present invention. For example, at least one of the pearlite and fresh martensite may be 30% or less in terms of area fraction. If at least one of the pearlite and fresh martensite exceeds 30% in terms of area fraction, it implies that the stability of the residual austenite is low, and since the residual austenite decomposes, the degree of transformation-induced plasticity is weak, which may result in inferior elongation characteristics. That is, if there is a large amount of pearlite and fresh martensite generated during the final cooling stage, the fraction of the final residual austenite decreases, which can adversely affect formability; therefore, the fraction of residual austenite transformed into pearlite and fresh martensite during the final cooling stage must be controlled to be 30% or less in terms of area fraction. Therefore, it is necessary to secure the stability of the austenite in advance prior to the alloying heat treatment step of the plating material, and the stability of the austenite can be evaluated in the final microstructure using the parameters disclosed in Formula 1 below.

[0104] (Formula 1)

[0105] (X Pearlite + X Fresh Martensite ) / (100 - X Ferrite - X Tempered Martensite - XBainite ) ≤ 0.30

[0106] In the above Formula 1, the above X Pearlite is the value of the area fraction of pearlite (unit: %), and the above X Fresh Martensite is the value of the area fraction of fresh martensite (unit: %), and the above X Ferrite is the value of the area fraction of ferrite (unit: %), and X Tempered Martensite is the value of the area fraction of tempered martensite (unit: %), and the above X Bainite is the value of the area fraction of bainite (unit: %).

[0107] In order to overcome the limitations of the mechanical properties of conventional DP steel or TRIP steel, the main matrix of TRIP steel is replaced with tempered martensite, bainite, and inevitably formed fresh martensite or pearlite instead of ferrite, thereby realizing a high-formability ultra-high-strength plated steel sheet for automobiles that can secure high strength and appropriate elongation.

[0108] Furthermore, the plated steel sheet of the present invention described above can improve liquid metal embrittlement (LME) characteristics.

[0109] A plated steel sheet according to one embodiment of the present invention comprises, in weight percent, carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, and boron (B): greater than 0 and less than or equal to 0.003%, and the remainder comprises iron (Fe) and other unavoidable impurities, wherein the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0; and a plating layer on the base steel plate; wherein the strain energy from the point of highest stress to the point of fracture in a tensile test on the plated steel plate is 30% or more of the strain energy from the point of highest stress to the point of fracture in a tensile test on the base steel plate.

[0110] Figure 2 is a graph showing the results of a Gleeble tensile test for the plated steel sheet and base steel sheet of the present invention. A Gleeble device, a high-temperature deformation simulator, was used to evaluate the tensile properties of the steel sheet under high-temperature conditions such as electric resistance welding. In Figure 2, the horizontal axis represents strain, which is the deformed length relative to a reference length, and the vertical axis represents stress received by the specimen, which is stress per unit area.

[0111] Referring to Fig. 2, when a tensile test is performed on a steel plate, it passes through the elastic deformation region (OY) and then through the strain hardening region (YT) until the point (T) where the maximum tensile strength (UTS), the point with the highest stress, appears. The point (Y) at one end of the elastic deformation region (OY) corresponds to the yield point. After passing the point (T) where the maximum tensile strength (UTS), the point with the highest stress, appears, a region (TB) occurs in which deformation continues but stress decreases; the material behavior phenomenon in this region is called necking. If stress is continuously applied to the specimen even after passing through the necking region, the specimen will eventually break at the point of fracture. Meanwhile, in the stress-strain graph, the area of ​​the region bounded by the horizontal axis and the stress curve corresponds to the strain energy.

[0112] A plated steel sheet according to one embodiment of the present invention is composed of a base steel sheet and a plating layer, wherein the deformation energy from the point of highest stress to the point of fracture in a tensile test on the plated steel sheet corresponds to the area of ​​the region (ZTC) bounded by the stress curve and the horizontal axis from the point (T) where the maximum tensile strength (UTS) appears in FIG. 2 to the point (C) where fracture occurs, and the deformation energy from the point of highest stress to the point of fracture in a tensile test on the base steel sheet corresponds to the area of ​​the region (ZTB) bounded by the stress curve and the horizontal axis from the point (T) where the maximum tensile strength (UTS) appears in FIG. 2 to the point (B) where fracture occurs.

[0113] A plated steel sheet according to one embodiment of the present invention is a plated steel sheet implemented with the alloy composition and manufacturing process described above, characterized in that the deformation energy from the point of highest stress to the point of fracture in a tensile test on the plated steel sheet is 30% or more of the deformation energy from the point of highest stress to the point of fracture in a tensile test on the base steel sheet. That is, when conducting a Glib test, the deformation energy from the point where maximum tensile strength (UTS) appears to the point where fracture occurs is characterized in that the energy reduction rate for the plated steel sheet is 70% or less compared to the base steel sheet which is unplated.

[0114] FIG. 3 is a diagram illustrating the simulation results of the interface movement between the plating layer and the substrate steel sheet immediately after spot welding in a plated steel sheet according to one embodiment of the present invention and the aluminum diffusion pattern, and FIG. 4 is a diagram illustrating the change in stability on the interface due to aluminum diffusion in a plated steel sheet according to one embodiment of the present invention.

[0115] Referring to FIGS. 3 and 4, it can be observed that when a plated steel sheet according to one embodiment of the present invention is subjected to electric resistance spot welding, liquid Zn and an alloy phase (Γ phase) are formed due to instantaneous heating and cooling, while the interface between the plating layer and the base steel sheet moves and aluminum (Al), a component contained in the base steel sheet, diffuses toward the interface. The point at which the concentration distribution stabilizes due to the diffusion of aluminum (Al) can be understood as the point in time 1.0 second after spot welding.

[0116] Specifically, referring to FIG. 3, the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0. As a result of simulation, the aluminum content is evaluated to be 2 to 3 weight% in the region within 0.1 μm in the direction of the substrate steel plate at the interface between the substrate steel plate and the plating layer.

[0117] Referring to Fig. 4, it can be seen that the interface composition and phase stability change due to the diffusion of the interface between the substrate steel plate and the plating layer, and that in the range where the aluminum content (concentration) is C1 (2 wt%) to C2 (3 wt%), the stability of the alloyed gamma (Γ) phase increases and the stability of the liquid pure zinc (Zn) phase decreases.

[0118] For a plated steel plate according to one embodiment of the present invention, immediately after spot welding, in a region within 0.1 μm in the direction of the base steel plate at the interface between the base steel plate and the plating layer, the volume fraction of the alloyed gamma (Γ) phase is 50% or more, and the ratio of the volume fraction of the alloyed gamma (Γ) phase to the volume fraction of the liquid pure zinc (Zn) phase is 1.5 or more. The period immediately after spot welding can be understood as a period within 1.0 second after spot welding.

[0119] So far, the plated steel sheet and the method for manufacturing the same according to the embodiments of the present invention have been described. In the present invention, the liquid metal embrittlement (LME) characteristics of QP steel to TRIP steel plated materials are improved through component control and process condition control. More specifically, by configuring the component system through the control of the total amount and ratio of silicon (Si) and aluminum (Al) alloys and manufacturing a QP / TRIP plated steel sheet of the target material through an appropriate heat treatment process, and then proceeding with spot welding, the amount of aluminum (Al) diffusion to the interface between the plating layer and the steel sheet varies depending on the total amount and ratio of silicon (Si) and aluminum (Al) alloys. Consequently, resistance to liquid metal embrittlement (LME) is secured through changes in phase stability in the interface region.

[0120] Experimental Example

[0121] Preferred experimental examples are presented below to aid in understanding the present invention. However, the following experimental examples are intended only to aid in understanding the present invention, and the present invention is not limited by the following experimental examples.

[0122] Table 1 shows the composition (unit: weight%) of the main components of the base steel sheet constituting the plated steel sheet according to the experimental example of the present invention. In addition, the composition of phosphorus (P): 0.01%, sulfur (S): 0.003%, nitrogen (N): 0.003%, and boron (B): 0.002% was applied identically.

[0123] C Si Mn Al [Si] / [Al] < 4 satisfaction Comparative Example 1 0.2 1.5 4 - X Comparative Example 2 0.2 1.5 4 - X Comparative Example 3 0.2 1.5 4 - X Comparative Example 4 0.2 1.25 4 0.25 X Comparative Example 5 0.2 1.25 4 0.25 X Comparative Example 6 0.2 1.25 4 0.25 X Invention Example 1 0.2 1 4 0.51 O Invention Example 2 0.2 1 4 0.51 O Invention Example 4 0.2 0.5 4 1 O Invention Example 5 0.2 0.5 4 1 O Invention Example 6 0.2 0.5 4 1 O Comparative Example 7 0.2 0.5 4 1 O Invention Example 7 0.2 - 4 1.5 O Invention Example 8 0.2 - 4 1.5 O Invention Example 9 0.2 - 4 1.5 O

[0124] Referring to Table 1, Invention Example 1-9 contains, in weight percent, carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, boron (B): greater than 0 and less than or equal to 0.003%, and satisfies all compositional ranges, with the remainder being iron (Fe), and satisfies all conditions that the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content is less than or equal to 4.0. Meanwhile, Invention Example 7-9 can be understood as having silicon (Si) added at 0.01% or less. This is because even if actual silicon (Si) is not intentionally added, minute amounts can be detected.

[0125] In contrast, Comparative Example 1-6 does not satisfy the condition that the ratio of the silicon (Si) content to the aluminum (Al) content is less than 4.0.

[0126] Table 2 shows the process conditions for manufacturing a galvanized steel sheet according to the experimental example of the present invention. The unit of the temperature item is °C, the unit of the speed item is °C / s, and the units of the time and holding items are seconds (s). In the experimental example of the present invention, other process conditions were applied as conditions having the same single value, satisfying the process conditions described with reference to FIG. 1. For example, the hot rolling process applied process conditions of reheating temperature: 1200 °C, finishing rolling temperature: 900 °C, and coiling temperature: 600 °C.

[0127] Annealing temperature Annealing hour Slow cooling speed Slow cooling temperature Rapid freezing speed Rapid freezing temperature Rapid freezing maintain reheat speed reheat temperature reheat maintain Alloying temperature Alloying maintain Comparative Example 1 800 80 3.4 750 46 230 34 88 460 30 480 24 Comparative Example 2 800 80 3.4 750 46 230 34 88 460 30 520 24 Comparative Example 3 800 80 3.4 750 46 230 34 88 460 30 560 24 Comparative Example 4 805 80 3.8 750 45 235 34 86 460 30 480 24 Comparative Example 5 805 80 3.8 750 45 235 34 86 460 30 520 24 Comparative Example 6 805 80 3.8 750 45 235 34 86 460 30 560 24 Invention Example 1 815 80 4.7 750 44 245 34 82 460 30 480 24 Invention Example 2 815 80 4.7 750 44 245 34 82 460 30 520 24 Invention Example 4 830 80 5.9 750 43 260 34 76 460 30 480 24 Invention Example 5 830 80 5.9 750 43 260 34 76 460 30 520 24 Invention Example 6 830 80 5.9 750 43 260 34 76 460 30 560 24 Comparative Example 7 830 80 5.9 750 43 260 1 76 460 30 560 24 Invention Example 7 850 80 7.5 750 42 270 34 72 460 30 480 24 Invention Example 8 850 80 7.5 750 42 270 34 72 460 30 520 24 Invention Example 9 850 80 7.5 750 42 270 34 72 460 30 560 24

[0128] Referring to Table 2, Invention 1-9 satisfies all conditions, including an annealing heat treatment temperature: 800 to 850°C, annealing heat treatment holding time: 60 to 130 seconds, slow cooling rate: less than 20°C, slow cooling end temperature: 650 to 750°C, rapid cooling rate: greater than 20°C, rapid cooling end temperature: 200 to 300°C, holding time after rapid cooling: 20 to 60 seconds, heating rate before reheating heat treatment: 40°C / s or more, reheating heat treatment temperature: 350 to 470°C, reheating heat treatment holding time: 20 to 50 seconds, alloying heat treatment temperature: 480 to 560°C, and alloying heat treatment holding time: 15 to 45 seconds. In contrast, Comparative Example 7 does not satisfy the range of holding time after rapid cooling: 20 to 60 seconds and falls below that range.

[0129] Table 3 shows the microstructural phase fractions (unit: %) of the substrate steel sheet constituting the plated steel sheet according to the experimental example of the present invention and the austenite decomposition rate during the alloying heat treatment step. In Table 3, the microstructure F represents the area fraction of ferrite, TM represents tempered martensite, B represents bainite, RA represents retained austenite, FM represents fresh martensite, and P represents pearlite; TM+B represents the sum of the area fractions of tempered martensite and bainite, and FM+P represents the sum of the area fractions of pearlite and fresh martensite. Meanwhile, the γ decomposition rate @ GA represents the decomposition rate of austenite during the alloying heat treatment step and can be calculated using the following Formula 1.

[0130] (Formula 1)

[0131] (X Pearlite + X Fresh Martensite ) / (100 - X Ferrite - X Tempered Martensite - X Bainite )

[0132] (Above X Pearlite is the value of the area fraction of pearlite (unit: %), and the above X Fresh Martensiteis the value of the area fraction of fresh martensite (unit: %), and the above X Ferrite is the value of the area fraction of ferrite (unit: %), and X Tempered Martensite is the value of the area fraction of tempered martensite (unit: %), and the above X Bainite is the value of the area fraction of bainite (unit: %)

[0133] F TM B RA FM P TM+B FM+P γ decomposition rate @ GA Comparative Example 1 10.1 51.2 10.9 18.3 6.8 2.7 62.1 9.5 0.342 Comparative Example 2 11.3 48.5 15.6 14.4 4.8 5.4 64.1 10.2 0.415 Comparative Example 3 10.6 50.3 12.2 12.8 7 7.1 62.5 14.1 0.524 Comparative Example 4 13.2 48.4 13.4 17.4 4.8 2.8 61.8 7.6 0.304 Comparative Example 5 12.6 50.2 11.8 15.7 4.8 4.9 62 9.7 0.382 Comparative Example 6 14.1 49.5 13.1 11.1 6.3 5.9 62.6 12.2 0.524 Invention Example 1 11.3 52.3 11.3 19.3 4.6 1.2 63.6 5.8 0.231 Invention Example 2 11.5 50.7 12.6 17.7 4.8 2.7 63.3 7.5 0.298 Invention Example 4 12.9 50.6 13.4 18.7 3.5 0.9 64 4.4 0.190 Invention Example 5 12.3 51.3 10.6 18.6 4.1 3.1 61.9 7.2 0.279 Invention Example 6 12.1 52.5 11.4 17.8 2.5 3.7 63.9 6.2 0.258 Comparative Example 7 11.3 51.4 10.2 16.4 4.9 5.8 61.6 10.7 0.395 Invention Example 7 12.7 51.6 14.8 17.3 3.6 0 66.4 3.6 0.172 Invention Example 8 11.6 51 14.3 18.4 3.3 1.4 65.3 4.7 0.203 Invention Example 9 12.8 51.5 13.6 17.2 2 2.9 65.1 4.9 0.222

[0134] Referring to Table 3, Invention Example 1-9 satisfies all of the ranges in which, as the final microstructure of the base steel sheet of the invention example, the area fraction is ferrite: 5 to 25%, retained austenite: 5 to 20%, the sum of tempered martensite and bainite: 30 to 90%, and the remainder is at least one of pearlite and fresh martensite, and all of the ranges in which the value of Formula 1 is 0.3 or less. In contrast, Comparative Example 1-7 can be confirmed to be a comparative example in which the value of Formula 1 does not satisfy the range of 0.3 or less and exceeds it. That is, in Comparative Example 1-7, it can be confirmed that the degree to which austenite decomposes into at least one of pearlite and fresh martensite during the alloying heat treatment step is higher than in the invention example.

[0135] Table 4 shows the physical properties of a plated steel sheet according to an experimental example of the present invention. In Table 4, YP represents yield strength (unit: MPa), TS represents tensile strength (unit: MPa), U.El represents uniform elongation (unit: %) until local elongation occurs, and T.El represents total elongation (unit: %) as the total elongation until fracture.

[0136] YP TS U.El T.El Material satisfaction Comparative Example 1 971 1344 9.2 11.8 X Comparative Example 2 966 1309 7.5 9.6 X Comparative Example 3 998 1258 6.9 9.1 X Comparative Example 4 983 1321 9.2 12.9 X Comparative Example 5 951 1316 7.5 10.8 X Comparative Example 6 955 1288 6.9 9.5 X Invention Example 1 967 1223 7.4 14.5 O Invention Example 2 934 1225 8.3 14.2 O Invention Example 4 935 1216 7.7 14.5 O Invention Example 5 899 1214 8.2 15.1 O Invention Example 6 874 1197 8.7 14.2 O Comparative Example 7 932 1307 7.6 11.2 X Invention Example 7 925 1211 8.3 15.9 O Invention Example 8 895 1201 8.8 15.3 O Invention Example 9 851 1184 9.1 14.7 O

[0137] Referring to Table 4, Inventive Example 1-9 satisfies all of the following criteria: yield strength: 850 to 1070 MPa, tensile strength: 1180 MPa or higher, and total elongation: 14% or higher. In contrast, Comparative Example 1-7 fails to satisfy the total elongation: 14% or higher and falls below all of these criteria. This is analyzed to be because in Comparative Example 1-7, the degree to which austenite decomposes into at least one of pearlite and fresh martensite during the alloying heat treatment step is higher than in the Inventive Example.

[0138] FIGS. 5 to 7 are photographs of the microstructure of the base steel sheet among the plated steel sheets according to Comparative Examples 1 to 3, respectively, and FIGS. 8 to 10 are photographs of the microstructure of the base steel sheet among the plated steel sheets according to Invention Examples 7 to 9, respectively.

[0139] Referring to FIGS. 5 to 7, the effect of silicon (Si) as an alloying component can be confirmed. Specifically, it can be seen that while it is effective to suppress carbide formation at a low alloying temperature (480°C), at a high alloying temperature (520°C / 560°C), the effect of suppressing carbide formation is less, and a large amount of cementite (θ) phase or pearlite (P) is generated. That is, since the chemical stability within the austenite is reduced due to the carbides formed during the alloying process, a relatively large phase transformation from austenite to pearlite or fresh martensite occurs during the final cooling process after the alloying heat treatment, and it can be seen that the reduction rate of residual austenite according to the alloying temperature is large.

[0140] Referring to FIGS. 8 to 10, the effect of the aluminum (Al) alloying component can be confirmed. Specifically, while some carbide formation is observed at a low alloying temperature (480°C), it can be confirmed that the effect of suppressing carbide formation at a high alloying temperature (520°C / 560°C) is higher than in FIGS. 5 to 7. That is, since the carbides formed during the alloying process are small, the chemical stability within the austenite is improved compared to FIGS. 5 to 7. Consequently, during the final cooling process after the alloying heat treatment, the phase transformation from austenite to pearlite or fresh martensite is relatively small, and the reduction rate of the residual austenite is small.

[0141] Meanwhile, referring to Comparative Examples 1 to 3 and Invention Examples 7 to 9, a decrease in tensile strength with increasing alloying temperature is commonly observed, and this is analyzed to be because transformation-induced plasticity is not exhibited due to the tempering effect of the matrix structure and the decomposition of residual austenite during the alloying heat treatment process.

[0142] Meanwhile, compared to Invention Examples 7 to 9, which are aluminum additives, Comparative Examples 1 to 3, which are silicon additives, show higher strength under the same conditions. This is understood to be because there is a difference in the solid solution strengthening effect of aluminum / silicon, and transformation-induced plasticity is not exhibited due to the decomposition of residual austenite.

[0143] Meanwhile, compared to Comparative Examples 1 to 3, which are silicon additives, Invention Examples 7 to 9, which are aluminum additives, show a higher elongation under the same conditions. This is analyzed to be because transformation-induced plasticity was not expressed due to the decomposition of residual austenite.

[0144] Ultimately, it can be understood that suppressing the phase transformation of residual austenite during alloying heat treatment is effective for securing material properties (especially elongation).

[0145] Table 5 shows the interface-related properties of a plated steel sheet according to an experimental example of the present invention. In Table 5, the Γ phase ratio (A) and the liquid phase Zn ratio (B) represent the volume fraction of the alloyed gamma (Γ) phase and the volume fraction of the liquid pure zinc (Zn) phase, respectively, in the region within 0.1 μm in the direction of the substrate steel sheet at the interface between the substrate steel sheet and the plating layer immediately after spot welding for the plated steel sheet. The A / B item represents the ratio of the volume fraction of the alloyed gamma (Γ) phase to the volume fraction of the liquid pure zinc (Zn) phase. The POST UTS energy reduction rate represents the energy reduction rate for the plated steel sheet, which is the plated material, based on the unplated substrate steel sheet, for the deformation energy from the point where the maximum tensile strength (UTS) appears to the point where fracture occurs during the Glib evaluation.

[0146] Γ phase ratio (A) (vol.%) Liquid Zn Ratio (B) (vol.%) A / B POST UTS Energy Reduction Rate (%) Satisfaction with plating material Comparative Example 1 44 56 0.79 81 X Comparative Example 2 46 54 0.85 83 X Comparative Example 3 43 57 0.75 77 X Comparative Example 4 45 55 0.82 82 X Comparative Example 5 47 53 0.89 80 X Comparative Example 6 41 59 0.69 84 X Invention Example 1 66 34 1.94 66 O Invention Example 2 63 37 1.70 61 O Invention Example 4 73 27 2.70 57 O Invention Example 5 76 24 3.17 59 O Invention Example 6 75 25 3.00 53 O Invention Example 7 83 17 4.88 51 O Invention Example 8 82 18 4.56 49 O Invention Example 9 85 15 5.67 55 O

[0147] Referring to Table 5, it can be confirmed that in Invention Example 1-9, immediately after spot welding of the plated steel plate, the volume fraction of the alloying gamma (Γ) phase in the region within 0.1 μm in the direction of the base steel plate at the interface between the base steel plate and the plating layer is 50% or more, and the ratio of the volume fraction of the alloying gamma (Γ) phase to the volume fraction of the liquid pure zinc (Zn) phase is 1.5 or more. In addition, during the Glib test, it can be confirmed that the deformation energy from the point where the maximum tensile strength (UTS) appears to the point where fracture occurs has an energy reduction rate of 70% or less for the plated steel plate compared to the unplated base steel plate. That is, it can be confirmed that the deformation energy from the point of highest stress to the point of fracture in the tensile test of the plated steel plate is 30% or more of the deformation energy from the point of highest stress to the point of fracture in the tensile test of the base steel plate. In the invention example, it was confirmed that the liquid metal embrittlement (LME) characteristics of QP steel or TRIP steel plated materials can be improved through component control and process condition control. More specifically, it was confirmed that by configuring the component system by controlling the total amount and ratio of silicon (Si) and aluminum (Al) alloys, manufacturing a QP / TRIP plated steel sheet of the target material through an appropriate heat treatment process, and then proceeding with spot welding, the amount of aluminum (Al) diffusion to the interface between the plating layer and the steel sheet varies depending on the total amount and ratio of silicon (Si) and aluminum (Al) alloys. Consequently, it was confirmed that resistance to liquid metal embrittlement (LME) can be secured through changes in phase stability in the interface region.

[0148] That is, the stability of the Γ phase increases due to changes in interface composition and phase stability caused by Al interface diffusion, and the fraction of liquid phase Zn at the interface decreases due to the decrease in the stability of liquid phase Zn. Consequently, the amount of penetration into the substrate steel plate is reduced due to the reduction of liquid phase Zn at the interface, thereby ensuring resistance to liquid metal embrittlement (LME).

[0149] In contrast, Comparative Examples 1-6, as comparative examples, show that immediately after spot welding of the plated steel plate, in the region within 0.1 μm in the direction of the base steel plate at the interface between the base steel plate and the plating layer, the volume fraction of the alloying gamma (Γ) phase does not satisfy the condition of being 50% or more and is all lower; the ratio of the volume fraction of the alloying gamma (Γ) phase to the volume fraction of the liquid pure zinc (Zn) phase does not satisfy the condition of being 1.5 or more and is lower; and during the Glib test, the deformation energy from the point where the maximum tensile strength (UTS) appears to the point where fracture occurs does not satisfy the condition of being 70% or less for the plated steel plate compared to the unplated base steel plate and is all higher.

[0150] According to the comparative example, the stability of the Γ phase decreases due to changes in interface composition and phase stability caused by Al interface diffusion, and the fraction of liquid phase Zn at the interface increases due to the increase in the stability of liquid phase Zn. Therefore, it can be analyzed that the resistance to liquid metal embrittlement (LME) deteriorates as the amount of penetration into the substrate increases due to the increase in liquid phase Zn at the interface.

[0151] Although the present invention has been described above with reference to embodiments, various changes and modifications may be made by those skilled in the art. Such changes and modifications are considered to be within the scope of the present invention as long as they do not depart from the scope of the present invention. Accordingly, the scope of rights of the present invention should be determined by the claims set forth below.

Claims

Claim 1 A base steel sheet characterized by containing, in weight%, carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, and boron (B): greater than 0 and less than or equal to 0.003%, and the remainder being iron (Fe) and other unavoidable impurities, wherein the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content ([Si] / [Al]; where [Si] and [Al] represent the silicon (Si) and aluminum (Al) content, respectively) is less than 4.0; A plated steel sheet comprising: a plating layer on the substrate steel sheet; wherein, immediately after spot welding with respect to the plated steel sheet, in an area within 0.1 μm in the direction of the substrate steel sheet from the interface between the substrate steel sheet and the plating layer, the volume fraction of the alloying gamma (Γ) phase is 50% or more, the plating layer is a zinc plating layer, and immediately after spot welding with respect to the plated steel sheet, in an area within 0.1 μm in the direction of the substrate steel sheet from the interface between the substrate steel sheet and the plating layer, the ratio of the volume fraction of the alloying gamma (Γ) phase to the volume fraction of the liquid pure zinc (Zn) phase is 1.5 or more, and the final microstructure of the substrate steel sheet is, in terms of area fraction, ferrite: 5 to 25%, retained austenite: 5 to 20%, total of tempered martensite and bainite: 30 to 90%, and the remainder is at least one of pearlite and fresh martensite, and the final microstructure of the substrate steel sheet is as follows A plated steel sheet characterized by satisfying Formula 1. (Formula 1)(X Pearlite + X Fresh Martensite ) / (100 - X Ferrite - X Tempered Martensite - X Bainite ) ≤ 0.30(the above X Pearlite is the value of the area fraction of pearlite (unit: %), and the above X Fresh Martensite is the value of the area fraction of fresh martensite (unit: %), and the above X Ferrite is the value of the area fraction of ferrite (unit: %), and X Tempered Martensite is the value of the area fraction of tempered martensite (unit: %), and the above X Bainite is the value of the area fraction of bainite (unit: %) Claim 2 delete Claim 3 delete Claim 4 A galvanized steel sheet according to claim 1, characterized in that the deformation energy from the point of highest stress to the point of fracture in a tensile test of the galvanized steel sheet is 30% or more of the deformation energy from the point of highest stress to the point of fracture in a tensile test of the base steel sheet, and has a yield strength of 850 to 1070 MPa, a tensile strength of 1180 MPa or more, and a total elongation of 14% or more. Claim 5 delete Claim 6 delete Claim 7 In weight%, it contains carbon (C): 0.1 to 0.3%, silicon (Si): greater than 0% and less than or equal to 2.0%, manganese (Mn): 2.0 to 4.0%, aluminum (Al): greater than 0 and less than or equal to 2.0%, phosphorus (P): greater than 0 and less than or equal to 0.02%, sulfur (S): greater than 0 and less than or equal to 0.005%, nitrogen (N): greater than 0 and less than or equal to 0.006%, boron (B): greater than 0 and less than or equal to 0.003%, and the remainder includes iron (Fe) and other unavoidable impurities, wherein the sum of the silicon (Si) content and the aluminum (Al) content is 2.0% or less, and the ratio of the silicon (Si) content to the aluminum (Al) content ([Si] / [Al]; where [Si] and [Al] represent the silicon (Si) and aluminum (Al) content, respectively) is 4.A step of providing a hot-rolled steel sheet by hot-rolling a steel material characterized by being less than 0 under conditions of a reheating temperature: 1150 ~ 1250℃, a finishing rolling temperature: 800 ~ 950℃, and a coiling temperature: 400 ~ 650℃; a step of providing a cold-rolled steel sheet by cold-rolling the hot-rolled steel sheet; a step of performing an annealing heat treatment by maintaining the cold-rolled steel sheet in an abnormal temperature range of 800 ~ 850℃ for 60 ~ 130 seconds; a step of slow cooling the annealed steel sheet to a slow cooling end temperature of 650 ~ 750℃ at a first cooling rate less than 20℃ / s; and a step of rapid cooling the slow-cooled steel sheet to a rapid cooling end temperature of 200 ~ 300℃, which is a temperature below the martensite transformation start temperature (Ms), at a second cooling rate greater than 20℃ / s. A method for manufacturing a plated steel sheet, comprising: a reheating step of maintaining the rapidly cooled steel sheet for a time ranging from 20 seconds to 50 seconds in a temperature range above the martensite transformation start temperature (Ms) and below the bainite transformation start temperature (Bs); a step of performing a plating treatment on the reheated steel sheet as a base steel sheet to form a plating layer on the base steel sheet; and an alloying heat treatment step of maintaining the base steel sheet and the plating layer for 15 to 45 seconds in a temperature range of 480 to 560°C, and then cooling to room temperature at a third cooling rate which is greater than the first cooling rate and smaller than the second cooling rate. Claim 8 delete Claim 9 delete Claim 10 delete

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