Hot-formed plated steel sheet with excellent impact resistance, hot-formed member, and manufacturing method thereof

Abstract

The present invention relates to a plated steel sheet for hot forming, a hot-formed member, and a method for manufacturing the same, and more particularly to a plated steel sheet for hot forming, a hot-formed member, and a method for manufacturing the same, which have excellent impact resistance.

Description

[Technical Field]

[0001] The present invention relates to a plated steel sheet for hot forming, a hot-formed member, and a method for manufacturing the same, and more particularly to a plated steel sheet for hot forming, a hot-formed member, and a method for manufacturing the same, which have excellent impact resistance. [Background technology]

[0002] Recently, hot-formed members have been widely applied to structural members of automobiles for the purposes of improving fuel efficiency and passenger protection through weight reduction of automobiles, and can be particularly used in bumpers, doors, or filler reinforcements that require ultra-high strength or high energy absorption capacity.

[0003] Patent Document 1 has been proposed as a representative technology relating to such hot forming technology. In this patent, an Al-Si plated steel sheet is heated to 850°C or higher, and then hot formed using a press and rapidly cooled to form the structure of the part into martensite, thereby ensuring ultra-high strength with high tensile strength. When such ultra-high strength steel for hot forming is used, forming at high temperatures makes it easy to form complex shapes, and the increased strength due to rapid cooling in the mold can be expected to lead to weight reduction effects.

[0004] However, at the same time, automobile manufacturers are increasingly demanding improved crashworthiness for passenger safety. However, ordinary hot-formed steel has a martensitic structure, which, while having high strength in a crash, has poor crashworthiness, so improvements are needed. [Prior art documents] [Patent documents]

[0005] [Patent Document 1] U.S. Patent No. 6,296,805 (published October 2, 2001) Summary of the Invention [Problem to be solved by the invention]

[0006] According to one aspect of the present invention, there is provided a plated steel sheet for hot forming, a hot-formed member, and methods for manufacturing the same, which have excellent impact resistance.

[0007] The object of the present invention is not limited to the above-mentioned content, and a person skilled in the art will have no difficulty in understanding further object of the present invention from the entire content of this specification. [Means for solving the problem]

[0008] One aspect of the present invention is a coating material comprising a base steel sheet containing, by weight, 0.06 to 0.5% carbon (C) and 0.01 to 0.1% antimony (Sb), and a coating layer formed on a surface of the base steel sheet, The base steel sheet contains an antimony (Sb)-enriched layer therein, When analyzing the element content in the thickness direction of the base steel sheet using a glow discharge spectrometer, the antimony (Sb) content in the antimony (Sb)-enriched layer is found to be at its maximum value (Sb max In this case, it is possible to provide a plated steel sheet in which the carbon (C) content at the depth showing the Cr content (C0) is 10 to 70% of the nominal carbon content (C0) of the base steel sheet.

[0009] The carbon (C) decarburization rate (α) in the region from the interface between the base steel sheet and the coating layer to a depth of 30 μm in the thickness direction may be 14 to 35%.

[0010] The plated steel sheet may have a point where the carbon (C) content is 50% of the nominal carbon content (C0) at a depth of more than 1.5 μm and less than 6 μm from the interface between the base steel sheet and the plated layer in the thickness direction.

[0011] The plated steel sheet may have a point where the carbon (C) content is 80% of the nominal carbon content (C0) at a depth of more than 6 μm and less than 15 μm from the interface between the base steel sheet and the plated layer in the thickness direction.

[0012] The plated steel sheet has an R value defined by the following relational expression 1 of 1.2 or more, The B value defined by the following relational expression 2 may be 0.008 or more. [Equation 1]

number

number

[0013] The region from the interface between the base steel sheet and the coating layer to a depth of 10 μm in the thickness direction can have a microstructure in which ferrite is the main phase and which contains 1 area % or more of pearlite.

[0014] The base steel sheet can contain carbon (C): 0.06 to 0.5%, antimony (Sb): 0.01 to 0.1%, silicon (Si): 0.001 to 2%, manganese (Mn): 0.1 to 4%, molybdenum (Mo): 1% or less, phosphorus (P): 0.05% or less, sulfur (S): 0.02% or less, aluminum (Al): 0.001 to 1%, chromium (Cr): 1% or less, nitrogen (N): 0.02% or less, titanium (Ti): 0.1% or less, boron (B): 0.01% or less, the balance being iron (Fe) and impurities.

[0015] The plating layer may be made of aluminum or an aluminum alloy.

[0016] Another aspect of the present invention is a coating material comprising: a base iron containing, by weight percent, 0.06 to 0.5% carbon (C) and 0.01 to 0.1% antimony (Sb); and a coating layer formed on a surface of the base iron, The base iron contains an antimony (Sb)-enriched layer therein, When analyzing the element content in the thickness direction of the base steel using a glow discharge spectrometer, the antimony (Sb) content in the antimony (Sb)-enriched layer was found to be the maximum value (Sb max In this way, it is possible to provide a component in which the carbon (C) content at the depth showing the tensile strength (T) is 80% or less of the nominal carbon content (C0) of the base steel.

[0017] The above-mentioned material has an antimony (Sb) content of up to a maximum value (Sb max The carbon (C) content at the depth showing the pore size distribution (C0) can be 15 to 80% of the nominal carbon content (C0) of the base iron.

[0018] The member has an R value defined by the following relational expression 1 of 1.5 or more, The B value defined by the following relational expression 2 may be 0.01 or more. [Equation 1]

number

number

[0019] The region from the interface between the base steel and the coating layer to a depth of 45 to 100 μm in the thickness direction may have a softening rate (β) of 2 to 7%.

[0020] The region extending from the interface between the base steel and the coating layer to a depth of 50 μm in the thickness direction can contain less than 5 area % of ferrite in the microstructure.

[0021] The region from the interface between the base steel and the coating layer to a depth of 50 μm in the thickness direction may have a microstructure with martensite as the main phase, less than 5 area % of ferrite, and the remainder upper and lower bainite.

[0022] The base iron may contain 0.06 to 0.5% carbon (C), 0.01 to 0.1% antimony (Sb), 0.001 to 2% silicon (Si), 0.1 to 4% manganese (Mn), 1% or less molybdenum (Mo), 0.05% or less phosphorus (P), 0.02% or less sulfur (S), 0.001 to 1% aluminum (Al), 1% or less chromium (Cr), 0.02% or less nitrogen (N), 0.1% or less titanium (Ti), 0.01% or less boron (B), and the remainder being iron (Fe) and impurities.

[0023] The plating layer may be made of aluminum or an aluminum alloy.

[0024] The member may have a product of tensile strength and bending angle of 80,000 MPa·° or more.

[0025] The member may have a diffusible hydrogen content of 0.2 ppm or less.

[0026] One aspect of the present invention is a method for manufacturing a steel sheet, the method comprising the steps of: preparing a cold-rolled steel sheet containing, by weight, carbon (C): 0.06 to 0.5% and antimony (Sb): 0.01 to 0.1%; Annealing the cold-rolled steel sheet in a temperature range of Ac1 to Ac3; and plating the annealed cold-rolled steel sheet, During the above annealing, the product of the annealing time and absolute humidity must be between 10,000 and 80,000 s·g / m 3 and It is possible to provide a method for producing a plated steel sheet in which, during the annealing, the average heating rate from room temperature to 500°C is 2.7 to 10.0°C / s, the average heating rate in the 500 to 700°C range is 0.5 to 2.5°C / s, and the average heating rate from 700°C to the annealing temperature is 0.01 to 0.4°C / s, based on the surface temperature of the steel sheet.

[0027] During the annealing, the annealing time is 100 to 200 seconds, and the absolute humidity is 100 to 400 g / m 3 It can be.

[0028] The cold-rolled steel sheet is Reheating the steel slab to a temperature range of 1050-1300°C; finish rolling the reheated steel slab at a temperature in the range of 800 to 950°C; Coiling and cooling the rolled steel at a temperature in the range of 500 to 700°C; and The method may include cold rolling the cooled steel at a reduction rate of 30 to 80%.

[0029] The steel slab may contain carbon (C): 0.06 to 0.5%, antimony (Sb): 0.01 to 0.1%, silicon (Si): 0.001 to 2%, manganese (Mn): 0.1 to 4%, molybdenum (Mo): 1% or less, phosphorus (P): 0.05% or less, sulfur (S): 0.02% or less, aluminum (Al): 0.001 to 1%, chromium (Cr): 1% or less, nitrogen (N): 0.02% or less, titanium (Ti): 0.1% or less, boron (B): 0.01% or less, the balance being iron (Fe) and impurities.

[0030] The plating may be performed with aluminum or an aluminum alloy.

[0031] Another aspect of the present invention is a method for manufacturing a plated steel sheet from a blank; heating the blank to a temperature range of Ac3 to 975°C and maintaining the temperature for 10 to 1000 seconds; A method for manufacturing a component may be provided, which includes forming and cooling the heated blank.

[0032] During the cooling, the cooling rate can be 20° C. / s or more. [Effects of the Invention]

[0033] According to one aspect of the present invention, it is possible to provide a plated steel sheet for hot forming, a hot-formed member, and methods for manufacturing the same, which have excellent impact resistance.

[0034] According to one aspect of the present invention, it is possible to provide a plated steel sheet for hot forming, a hot-formed member, and methods for manufacturing the same, which are excellent in fatigue resistance and crash resistance. [Brief explanation of the drawings]

[0035] [Figure 1] 1 is a schematic diagram showing an exemplary change in Sb content according to the present invention to represent an Sb-enriched layer. [Figure 2] 1 is a schematic diagram showing the profiles of Sb and C contents in the thickness direction from the interface in a plated steel sheet according to one embodiment of the present invention. [Figure 3] 1 is a schematic diagram showing a profile of the decarburization rate (α) in the thickness direction from the interface in a plated steel sheet according to one embodiment of the present invention. [Figure 4] 1 is a schematic diagram showing the profile of an Sb-enriched layer of a plated steel sheet according to an embodiment of the present invention. [Figure 5] 1 is a schematic diagram showing the profiles of Sb and C contents in the thickness direction from the interface in a member according to an embodiment of the present invention. [Figure 6] 1 is a schematic diagram showing a hardness softening rate (β) profile in the thickness direction from the interface in a member according to an embodiment of the present invention. [Figure 7] 1 shows a profile of the C content in a plated steel sheet according to an embodiment of the present invention. [Figure 8] 1 is a photograph of the microstructure of a plated steel sheet according to an embodiment of the present invention, observed with a scanning electron microscope (SEM). [Figure 9] 1 is a photograph showing the microstructure of a member according to an embodiment of the present invention. DETAILED DESCRIPTION OF THE INVENTION

[0036] Hereinafter, preferred embodiments of the present invention will be described. The embodiments of the present invention can be modified in various forms, and the scope of the present invention should not be construed as being limited to the embodiments described below. The present embodiments are provided to further explain the present invention in detail to those skilled in the art to which the invention pertains.

[0037] To solve the problems of the prior art, it has been considered to apply decarburization technology to improve bendability, but this can lead to deterioration of fatigue properties due to a local decrease in hardness in the surface layer, limiting the application of this technology to automotive components. The present inventors conducted extensive research to solve these problems and discovered that forming an antimony (Sb)-enriched layer in the base steel sheet and maintaining the decarburization rate of the steel sheet at an appropriate level can solve the problem of deterioration of fatigue properties and improve internal crash properties, thereby achieving the present invention.

[0038] The present invention will now be described in more detail.

[0039] One aspect of the present invention is directed to a plated steel sheet comprising a base steel sheet and a coating layer formed on the surface of the base steel sheet. Impact resistance and fatigue resistance can be significantly affected by the degree of decarburization of the steel sheet, but the effects of the present invention can be advantageously achieved by appropriately controlling the decarburization rate using a concentrated layer formed within the base steel sheet. That is, a plated steel sheet according to one aspect of the present invention comprises a base steel sheet and a coating layer formed on the surface of the base steel sheet (which may refer to the interface between the base steel sheet and the coating layer), and the base steel sheet comprises an antimony (Sb) concentrated layer formed therein.

[0040] According to an embodiment of the present invention, when an antimony (Sb)-enriched layer is formed in a steel sheet, the decarburization rate according to the depth in the thickness direction of the steel sheet can be appropriately controlled according to the formation of the Sb-enriched layer.

[0041] The Sb-enriched layer according to one embodiment of the present invention and its role will be described in detail below with reference to the graph in FIG. 1. FIG. 1 is a schematic diagram illustrating an exemplary change in Sb content according to the present invention to represent an Sb-enriched layer. The x-axis in FIG. 1 represents the linear distance from the coating layer at an arbitrary position in the coated steel sheet in the lateral direction of the base steel sheet, i.e., the thickness direction, and the y-axis represents the Sb content measured using a glow discharge spectrometer (GDS). FIG. 1 illustrates the change in Sb content in the coating layer 1, the Sb-enriched layer 2, and the base steel sheet 3 excluding the Sb-enriched layer 2. Here, the Sb-enriched layer 2 may have an Sb content that is 1.05 times or more the nominal Sb content (Sb0) of the base steel sheet, and the Sb-enriched layer 2 has a maximum Sb content at point 200 (Sb max ) can exist. In addition, the Sb-enriched layer 2 can be formed by advancing in the x-axis direction until it reaches a point 200 (Sb max ) and the point 200 where the Sb content is at its maximum (Sb max ) in the x-axis direction, the Sb content decreases. In FIG. 1, the average Sb content line 10 of the coating layer 1 and the point 200 (Sb max ), the last contact point 11 in the x-axis (+) direction of the Sb content line 100 can be set as the starting point of the rising section 21 of the Sb content.

[0042] In one embodiment, the average Sb content line 10 of the coating layer 1 is a point 200 (Sb max ) can refer to a horizontal extension of the average Sb content in the section from point A, which is 15 μm away from the coating layer 1 side, to point B, which is 20 μm away from the coating layer 1 side.

[0043] Similarly, the point 200 (Sb maxIn the descending section 22 of the Sb content in the (+) direction of the x-axis from the line 30, the first contact point 31 in the (+) direction of the x-axis between the average Sb content line 30 of the base steel sheet and the Sb content line 100 is regarded as the end point of the Sb-enriched layer 2.

[0044] In one embodiment, the Sb average content line 30 of the base steel sheet 3 excluding the Sb-enriched layer is a point 200 (Sb max ) to point C, which is 15 μm away from the base steel sheet 3, to point D, which is 20 μm away from the base steel sheet 3.

[0045] In one embodiment of the present invention, the Sb-enriched layer may be formed directly below the interface between the base steel sheet and the coating layer. For example, when the profile of Al content is analyzed in the depth (thickness) direction from the surface of the coated steel sheet using a glow discharge spectrometer (GDS), the Sb-enriched layer may be defined as the point where the Al content is 15%. Furthermore, the thickness of the Sb-enriched layer may be, for example, 1 to 30 μm.

[0046] In addition, according to an embodiment of the present invention, when analyzing the antimony (Sb) content in the thickness direction of the base steel sheet using a glow discharge spectrometer (GDS), the antimony (Sb) content in the antimony (Sb)-enriched layer is found to have a maximum value (Sb max ) may be 10 to 70% of the nominal carbon content (C0) of the base steel sheet. In one embodiment of the present invention, the nominal carbon content (C0) may refer to the average carbon content in a 1 / 4 to 3 / 4 thickness region based on the cross section of the base steel sheet, and more specifically, may be the average carbon content obtained by analyzing a carbon profile at a distance of 50 μm or more from an arbitrary point in a 1 / 4 to 3 / 4 thickness region of the base steel sheet using a glow discharge spectrometer (GDS).

[0047] FIG. 2 is a schematic diagram showing the profiles of Sb and C content in the thickness direction from the interface in a coated steel sheet according to an embodiment of the present invention. The x-axis of FIG. 2 represents the depth (μm) from the interface between the base steel sheet and the coating layer, and the y-axis represents the element content (wt%). As shown in FIG. 2, 70% of the nominal carbon content (C0) is 0.154%. Here, the nominal carbon content (C0) is 0.22%, which is obtained by analyzing a certain thickness (depth) in the region of 1 / 4 to 3 / 4 of the thickness of the base steel sheet using a glow discharge spectrometer (GDS), as described above. At this time, the Sb content reaches its maximum value (Sb max It can be seen that the carbon content at the depth showing ) is 70% or less of the nominal carbon content (C0).

[0048] As shown in FIG. 2, in one embodiment of the present invention, the Sb content in the Sb-enriched layer is at a maximum value (Sb max At the depth where the carbon content is 10 to 70%, the ratio of the carbon content to the nominal carbon content (C0) is controlled to 10 to 70%, and the carbon content at this time affects the softening rate of the surface hardness and bendability of the component.

[0049] On the other hand, the Sb content reaches its maximum value (Sb max If the carbon content exceeds 70% of the nominal carbon content (C0) at the depth where the nominal carbon content (C0) is reached, the hardness of the surface layer may increase, resulting in deterioration of bendability. Also, according to one embodiment of the present invention, if the carbon content is less than 10% of the nominal carbon content (C0), the hardness may decrease excessively, resulting in deterioration of fatigue resistance.

[0050] In one embodiment of the present invention, the Sb content in the Sb-enriched layer is set to a maximum value (Sb max At the depth at which the carbon content is less than the nominal carbon content (C0), the ratio of the carbon content to the nominal carbon content (C0) may be 10.0 to 70.0%.

[0051] According to one embodiment of the present invention, the carbon (C) decarburization rate (α) in the region from the interface between the base steel sheet and the coating layer to a depth of 30 μm in the thickness direction may be 14 to 35%.

[0052] FIG. 3 is a schematic diagram illustrating a decarburization rate (α) profile in the thickness direction from the interface of a coated steel sheet according to an embodiment of the present invention. In FIG. 3, the decarburization rate (α) can be obtained from the results of measuring carbon in the coated steel sheet using a glow discharge spectroscopy (GDS). The y-axis in the drawing represents the ratio (%) of the carbon content at that location relative to the nominal carbon content (C0), and the x-axis represents the distance (μm) in the thickness (depth) direction from the interface between the base steel sheet and the coating layer. As shown in the drawing, a rectangle can be drawn with a horizontal side corresponding to a depth of 0 to 30 μm from the interface in the thickness direction of the base steel sheet along the x-axis, and a vertical side corresponding to a length of 0 to 100% along the y-axis. A carbon profile curve showing the ratio of the carbon content at that depth relative to the nominal carbon content (C0) is plotted within the rectangle, and the decarburization rate (α) can be defined as the ratio (%) of the area above the carbon profile curve within the rectangle to the total area of ​​the rectangle.

[0053] That is, the decarburization rate (α) in the present invention means the ratio (%) of the area of ​​the region above the carbon profile curve to the total area of ​​a rectangle, with the horizontal axis representing the distance (μm) in the thickness (depth) direction from the interface between the base steel sheet and the coating layer and the vertical axis representing the ratio (%) of the carbon content at that position to the nominal carbon content (C0).

[0054] If the carbon (C) decarburization rate (α) in the region from the interface to a depth of 30 μm in the thickness direction is less than 14%, the carbon concentration in the base steel sheet will excessively increase the hardness of the part after hot forming, which may significantly reduce the effect of improving bendability.On the other hand, if the decarburization rate exceeds 35%, the carbon content in the surface layer of the base steel sheet will decrease significantly, resulting in a problem of poor fatigue resistance of the part.

[0055] According to one embodiment of the present invention, the carbon (C) decarburization rate (α) in a region from the interface between the base steel sheet and the coating layer to a depth of 30.0 μm in the thickness direction may be 14.0 to 35.0%.

[0056] In the plated steel sheet according to another embodiment of the present invention, a point where the carbon (C) content is 50% of the nominal carbon content (C0) may exist at a depth of more than 1.5 μm and less than 6 μm from the interface between the base steel sheet and the coating layer in the thickness direction.

[0057] The reason for controlling the carbon (C) content ratio according to the nominal carbon content (C0) at a depth of more than 1.5 μm but less than 6 μm from the interface in the thickness direction is to simultaneously ensure fatigue resistance and impact resistance. If there is a point within this depth range where the carbon (C) content is 50% of the nominal carbon content (C0), it is advantageous for simultaneously ensuring impact resistance and fatigue resistance. However, if the 50% point is located at a depth of 6 μm or more, excessive decarburization can cause deterioration of fatigue resistance. On the other hand, if the 50% point is located at a depth of 1.5 μm or less, decarburization is insufficient, making it difficult to ensure the desired bendability.

[0058] In the plated steel sheet according to another embodiment of the present invention, a point where the carbon (C) content is 50.0% of the nominal carbon content (C0) may exist at a depth of more than 1.50 μm and less than 6.0 μm from the interface between the base steel sheet and the coating layer in the thickness direction.

[0059] In addition, in one embodiment of the present invention, a point where the carbon (C) content is 80% of the nominal carbon content (C0) may exist at a depth of more than 6 μm and less than 15 μm in the thickness direction from the interface between the base steel sheet and the coating layer.

[0060] If a point where the carbon (C) content ratio according to the nominal carbon content (C0) is 80% exists at a depth of more than 6 μm but less than 15 μm from the interface in the thickness direction, this can be advantageous for ensuring appropriate bendability and preventing excessive deterioration of fatigue resistance. On the other hand, if the 80% point exists at a depth of 15 μm or more, fatigue resistance may be reduced due to excessive decarburization, and if the 80% point exists at a depth of 6 μm or less, decarburization may be insufficient, making it difficult to ensure the desired bendability.

[0061] In one embodiment of the present invention, a point where the carbon (C) content is 80.0% of the nominal carbon content (C0) may exist at a depth of more than 6.0 μm and less than 15.0 μm in the thickness direction from the interface between the base steel sheet and the coating layer.

[0062] In one embodiment of the present invention, the R value defined by the following relational expression 1 may be 1.2 or more, and the B value defined by the following relational expression 2 may be 0.008 or more.

[0063] In one embodiment of the present invention, the R value defined by the following relational expression 1 may be 1.20 or more, and the B value defined by the following relational expression 2 may be 0.0080 or more.

[0064] When an Sb-enriched layer is formed in the base steel sheet, it becomes difficult for oxygen dissociated in the annealing furnace to penetrate into the base steel sheet, and the layer can act as a protective film that makes decarburization difficult. In the present invention, it has been confirmed that the decarburization rate can be appropriately controlled by controlling the Sb content according to the depth in the thickness direction, and the present invention proposes the following Relational Formulas 1 and 2. [Equation 1]

number

number

[0065] Fig. 4 is a schematic diagram showing the profile of the Sb-enriched layer of the plated steel sheet according to one embodiment of the present invention. In Fig. 4, the area corresponding to the B value of the above-mentioned relational expression 2 is shown by a shaded area. coat The measurement points and Sb maxThe degree of Sb enrichment can be shown by Δt, which indicates the distance between the measurement points.

[0066] If the R value defined by the above relational expression 1 is less than 1.2 or the B value defined by the above relational expression 2 is less than 0.008, excessive decarburization occurs, the decarburization rate in the plated steel sheet becomes too high, and the surface hardness of the member after hot forming is significantly reduced, which may deteriorate the fatigue resistance of the member.

[0067] In one embodiment of the present invention, the R value defined by Relational Formula 1 may be limited to 1.5 or more. In another embodiment of the present invention, the B value defined by Relational Formula 2 may be limited to 0.02 or more. However, if the R value or the B value is excessively high, the carbon on the surface may not be removed, resulting in excessively high surface hardness of the part after hot forming, which may reduce the bendability of the surface layer. Therefore, in one embodiment of the present invention, the upper limit of the R value may be limited to 6.5. In another embodiment of the present invention, the upper limit of the B value may be limited to 0.15.

[0068] In one embodiment of the present invention, the R value defined by Relational Formula 1 may be limited to 1.50 or more. In another embodiment of the present invention, the B value defined by Relational Formula 2 may be limited to 0.020 or more. However, if the R value or the B value is excessively high, the carbon on the surface may not be removed, resulting in excessively high surface hardness of the part after hot forming, which may reduce the bendability of the surface layer. Therefore, in one embodiment of the present invention, the upper limit of the R value may be limited to 6.50. In another embodiment of the present invention, the upper limit of the B value may be limited to 0.150.

[0069] As described above, by controlling the R-value and B-value of the plated steel sheet within the suggested range, the R-value and B-value of the component can be controlled within an appropriate range, thereby effectively suppressing hydrogen penetration.

[0070] According to a further embodiment of the present invention, the plated steel sheet may have a microstructure in which ferrite is the main phase and pearlite accounts for 1 area % or more in a region from the interface between the base steel sheet and the coating layer to a depth of 10 μm in the thickness direction. In the present invention, a phase that accounts for 50 area % or more of the total area of ​​the microstructure can be considered as the main phase.

[0071] In the plated steel sheet of the present invention, if the ferrite fraction in the region from the interface between the base steel sheet and the plated layer to a depth of 10 μm in the thickness direction is insufficient, the fatigue resistance of the member may deteriorate.

[0072] The pearlite in the region up to a distance of 10 μm in the thickness (depth) direction from the interface between the base steel sheet and the coating layer provides carbon to the structure directly below the coating layer during heat treatment for hot forming, thereby preventing deterioration of the hardness of the surface layer. Therefore, the present invention may contain pearlite at an area percentage of 1% or more.

[0073] On the other hand, if the pearlite content is less than 1% by area, the hardness of the surface layer after hot forming may decrease excessively, resulting in a high hardness softening rate, which may cause a problem of deterioration in the fatigue resistance of the member.

[0074] According to another embodiment of the present invention, the region from the interface between the base steel sheet and the coating layer to a depth of 10.0 μm in the thickness direction can contain ferrite as the main phase and 1.0 area % or more of pearlite.

[0075] The composition of the base steel sheet of the present invention will be described in detail below.

[0076] The base steel sheet according to one embodiment of the present invention may contain, by weight, carbon (C): 0.06 to 0.5% and antimony (Sb): 0.01 to 0.1%.

[0077] The base steel sheet according to one embodiment of the present invention may contain, by weight, carbon (C): 0.060 to 0.50%, and antimony (Sb): 0.010 to 0.10%.

[0078] In the present invention, unless otherwise specified, the percentage representing the content of each element is based on weight.

[0079] Carbon (C): 0.06~0.5% Carbon (C) is an element that improves the strength and hardenability of hot-formed parts and must be added appropriately as an essential element for strength adjustment. If the carbon (C) content is less than 0.06%, the hardenability is low, and as the cooling rate decreases, sufficient martensite cannot be obtained, making it difficult to obtain the desired strength due to the formation of ferrite. In one embodiment of the present invention, the carbon (C) content may be 0.1% or more. On the other hand, if the carbon (C) content exceeds 0.5%, the strength may increase excessively, which may induce brittleness and poor weldability. In one embodiment of the present invention, the upper limit of the carbon (C) content may be 0.45%. According to another embodiment of the present invention, carbon (C) may be contained in an amount of 0.060 to 0.50%. According to another embodiment of the present invention, carbon (C) may be 0.10% or more. According to another embodiment of the present invention, the upper limit may be 0.450%.

[0080] Antimony (Sb): 0.01-0.1% Antimony (Sb) enriches in the base steel sheet, thereby controlling the amount of carbon released during internal oxidation annealing and preventing excessive hardness loss in the component. If the antimony (Sb) content is less than 0.01%, a sufficient enriched layer is not formed at the interface between the coating layer and the base steel sheet, causing excessive decarburization and resulting in an excessive decrease in surface hardness, which can degrade fatigue resistance. According to one embodiment of the present invention, the lower limit of the antimony (Sb) content may be 0.02%. On the other hand, if the antimony (Sb) content exceeds 0.1%, excessive antimony (Sb) precipitates at grain boundaries, which can induce grain boundary fracture under stress, resulting in material degradation. According to one embodiment, the upper limit of the antimony (Sb) content may be 0.08%. According to another embodiment of the present invention, antimony (Sb) can be contained in an amount of 0.010 to 0.10%. According to another embodiment of the present invention, antimony (Sb) can be 0.020% or more. According to another embodiment of the present invention, the upper limit may be 0.080%.

[0081] The type and content of the additive elements of the base steel sheet used in the hot-forming coated steel sheet of the present invention are not particularly limited as long as they are commonly added. However, non-limiting examples of elements that can be added to the base steel sheet according to one embodiment of the present invention include silicon (Si), manganese (Mn), molybdenum (Mo), phosphorus (P), sulfur (S), aluminum (Al), chromium (Cr), nitrogen (N), titanium (Ti), boron (B), copper (Cu), nickel (Ni), vanadium (V), calcium (Ca), niobium (Nb), tin (Sn), tungsten (W), magnesium (Mg), cobalt (Co), arsenic (As), zirconium (Zr), bismuth (Bi), and rare earth elements (REM), and the steel sheet may further contain one or more of these.

[0082] According to one embodiment of the present invention, the base steel sheet can contain, in weight percent, silicon (Si): 0.001 to 2%, manganese (Mn): 0.1 to 4%, molybdenum (Mo): 1.0% or less, phosphorus (P): 0.05% or less, sulfur (S): 0.02% or less, aluminum (Al): 0.001 to 1%, chromium (Cr): 1.00% or less, nitrogen (N): 0.02% or less, titanium (Ti): 0.1% or less, boron (B): 0.01% or less, the balance being iron (Fe) and impurities.

[0083] Silicon (Si): 0.001 to 2% Silicon (Si) can be added as a deoxidizer during steelmaking. It is also a solution strengthening element and an inhibitor of carbide formation, effectively homogenizing the internal structure. It also contributes to increasing the strength of hot-formed components and is added as an effective element for homogenizing material properties. However, if the Si content is less than 0.001%, the above effects cannot be expected. If the Si content exceeds 2%, excessive Si oxides formed on the steel sheet surface during annealing can significantly reduce galvanic properties. According to one embodiment of the present invention, the lower limit of the silicon (Si) content can be 0.005%, and in some cases, 0.01%. According to another embodiment of the present invention, the upper limit of the silicon (Si) content can be 0.7%, and in some cases, 0.65%. According to another embodiment of the present invention, silicon (Si) can be contained in an amount of 0.001 to 2.0%. According to another embodiment of the present invention, silicon (Si) may be 0.0050% or more. According to another embodiment of the present invention, the upper limit can be 0.70%. According to another embodiment of the present invention, silicon (Si) may be 0.010% or more. According to another embodiment of the present invention, the upper limit may be 0.650%.

[0084] Manganese (Mn): 0.1-4% Manganese (Mn) is necessary not only to ensure the desired strength through solid solution strengthening, but also to suppress ferrite formation during hot forming by improving hardenability. If the manganese (Mn) content is less than 0.1%, it is difficult to obtain sufficient hardenability, and the insufficient hardenability requires excessive addition of other expensive alloying elements, which can significantly increase manufacturing costs. According to one embodiment of the present invention, manganese (Mn) may be contained in an amount of 0.5% or more, and in another embodiment, 0.8% or more. However, if the manganese (Mn) content exceeds 4%, the band-like structure aligned in the rolling direction of the microstructure deepens, causing inhomogeneity in the internal structure and resulting in deterioration of impact resistance. In one embodiment of the present invention, the upper limit of the manganese (Mn) content may be 3.5%. According to another embodiment of the present invention, manganese (Mn) may be contained in an amount of 0.010 to 4.0%. According to another embodiment of the present invention, manganese (Mn) may be contained in an amount of 0.050 to 4.0%. According to another embodiment of the present invention, manganese (Mn) may be contained in an amount of 0.080 to 4.0%. According to another embodiment of the present invention, manganese (Mn) may be contained in an amount of 0.050 to 3.50%. According to another embodiment of the present invention, manganese (Mn) may be contained in an amount of 0.080 to 3.50%.

[0085] Molybdenum (Mo): 1.0% or less Molybdenum (Mo) can be included as an element that strengthens grains and improves bendability. However, if the Mo content exceeds 1.0%, manufacturing costs may increase significantly. According to one embodiment of the present invention, the upper limit of the Mo content may be 0.5%, or even 0.45%. According to another embodiment of the present invention, molybdenum (Mo) may be contained in an amount of 1.0% or less. According to one embodiment of the present invention, the upper limit of the molybdenum (Mo) content may be 0.50%, and in some cases, may be 0.450%.

[0086] Phosphorus (P): 0.05% or less Phosphorus (P) exists as an impurity in steel, and if its content exceeds 0.05%, it can deteriorate the weldability of hot-formed parts and the material properties due to high-temperature grain boundary segregation. In one embodiment, the upper limit can be limited to 0.015%. Meanwhile, in one embodiment of the present invention, the lower limit can be limited to 0.001% because controlling the P content to a very low level incurs significant manufacturing costs. According to another embodiment of the present invention, phosphorus (P) may be contained in an amount of 0.050% or less. According to one embodiment, the upper limit may be limited to 0.0150%, while in some cases the lower limit may be limited to 0.0010%.

[0087] Sulfur (S): 0.02% or less Sulfur (S) is an impurity in steel that impairs the ductility, impact properties, and weldability of components, so its upper limit can be limited to 0.02%. In one embodiment of the present invention, controlling its content to a very low level can significantly increase manufacturing costs, so its lower limit can be limited to 0.0001%. According to another embodiment of the present invention, sulfur (S) can be contained in an amount of 0.020% or less. According to one embodiment, the lower limit can be limited to 0.00010%.

[0088] Aluminum (Al): 0.001 to 1% Aluminum (Al) is an element that, together with Si, acts as a deoxidizer during steelmaking, increasing the cleanliness of the steel. If the aluminum (Al) content is less than 0.001%, this effect may be difficult to achieve. According to one embodiment of the present invention, the lower limit of aluminum (Al) may be 0.01%, and in some cases, 0.02%. On the other hand, if the aluminum (Al) content exceeds 1%, excessive AlN precipitates formed during the continuous casting process may reduce high-temperature ductility and cause slab cracking, leading to manufacturing problems. In one embodiment, the upper limit may be limited to 0.1%, and in some cases, 0.09%. According to another embodiment of the present invention, aluminum (Al) can be contained in an amount of 0.0010 to 1.0%. According to another embodiment of the present invention, aluminum (Al) can be contained in an amount of 0.010 to 1.0%. According to another embodiment of the present invention, aluminum (Al) can be contained in an amount of 0.020 to 1.0%. According to another embodiment of the present invention, aluminum (Al) may be contained in an amount of 0.010 to 0.10%. According to another embodiment of the present invention, aluminum (Al) may be contained in an amount of 0.010 to 0.090%. According to another embodiment of the present invention, aluminum (Al) may be contained in an amount of 0.020 to 0.10%. According to another embodiment of the present invention, aluminum (Al) may be contained in an amount of 0.020 to 0.090%.

[0089] Chromium (Cr): 1% or less Chromium (Cr), like Mn, can be added as an element to ensure the hardenability of steel and suppress the formation of ferrite after hot forming. If the chromium (Cr) content exceeds 1%, not only is the effect of improving hardenability relative to the amount added small, but excessive formation of coarse iron carbides can induce cracks under stress, resulting in material degradation. In one embodiment of the present invention, the upper limit can be 0.8%. Meanwhile, in one embodiment of the present invention, the lower limit can be limited to 0.01%, or in some cases, 0.05%, to effectively ensure the above-mentioned effects. According to another embodiment of the present invention, chromium (Cr) may be contained in an amount of 1.0% or less. According to another embodiment of the present invention, chromium (Cr) may be contained in an amount of 0.80% or less. According to another embodiment of the present invention, chromium (Cr) may be contained in an amount of 0.01 to 1.0%. According to another embodiment of the present invention, chromium (Cr) may be contained in an amount of 0.01 to 0.8%. According to another embodiment of the present invention, chromium (Cr) may be contained in an amount of 0.05 to 1.0%. According to another embodiment of the present invention, chromium (Cr) may be contained in an amount of 0.05 to 0.8%.

[0090] Nitrogen (N): 0.02% or less Nitrogen (N) can be contained in steel as an impurity. If the nitrogen (N) content exceeds 0.02%, it may form AlN like added Al, which may cause slab cracking. However, controlling the nitrogen (N) content to a minimum may incur excessive manufacturing costs, so in one embodiment of the present invention, the lower limit of nitrogen (N) may be limited to 0.001%. According to another embodiment of the present invention, nitrogen (N) may be contained at 0.020% or less. According to another embodiment of the present invention, nitrogen (N) can be contained in an amount of 0.0010 to 0.02%. According to another embodiment of the present invention, nitrogen (N) can be contained in an amount of 0.0010 to 0.020%.

[0091] Titanium (Ti): 0.1% or less Titanium (Ti) combines with N remaining as an impurity in steel to form TiN, thereby protecting B, which ensures hardening, from forming compounds. It also contributes to precipitation strengthening and grain refinement through the formation of TiC precipitates. However, if the Ti content exceeds 0.1%, a large amount of coarse TiN is formed, degrading the quality of the steel. In one embodiment of the present invention, the upper limit of the Ti content can be set to 0.09%. According to another embodiment of the present invention, titanium (Ti) may be contained in an amount of 0.10% or less. According to another embodiment of the present invention, titanium (Ti) may be contained in an amount of 0.090% or less.

[0092] Boron (B): 0.01% or less Boron (B) is an element that can effectively improve hardening ability and segregates at the prior austenite grain boundaries to suppress the brittleness of hot-formed parts caused by the grain boundary segregation of impurities such as P or S. However, when its content exceeds 0.01%, Fe 23 The formation of CB6 complex compounds can cause embrittlement during hot rolling. In one embodiment of the present invention, the upper limit of the boron (B) content can be limited to 0.008%. According to another embodiment of the present invention, boron (B) may be contained in an amount of 0.010% or less. According to another embodiment of the present invention, boron (B) may be contained in an amount of 0.0080% or less.

[0093] Furthermore, as one embodiment of the present invention, the composition may further contain one or more of copper (Cu): 1% or less, nickel (Ni): 1% or less, vanadium (V): 1.0% or less, calcium (Ca): 0.01% or less, niobium (Nb): 0.1% or less, tin (Sn): 1% or less, tungsten (W): 1% or less, magnesium (Mg): 0.1% or less, cobalt (Co): 1% or less, arsenic (As): 1% or less, zirconium (Zr): 1% or less, bismuth (Bi): 1% or less, and rare earth elements (REM): 0.3% or less.

[0094] Furthermore, as one embodiment of the present invention, the composition may further contain one or more of copper (Cu): 1.0% or less, nickel (Ni): 1.0% or less, vanadium (V): 1.0% or less, calcium (Ca): 0.010% or less, niobium (Nb): 0.10% or less, tin (Sn): 1.0% or less, tungsten (W): 1.0% or less, magnesium (Mg): 0.10% or less, cobalt (Co): 1.0% or less, arsenic (As): 1.0% or less, zirconium (Zr): 1.0% or less, bismuth (Bi): 1.0% or less, and rare earth elements (REM): 0.30% or less.

[0095] In addition to the above-described composition, the base steel sheet of the present invention may contain the remaining iron (Fe) and inevitable impurities. Since inevitable impurities may be unintentionally mixed in during normal manufacturing processes, they cannot be excluded. Since such impurities are known to anyone skilled in the field of normal steel manufacturing, the full details of such impurities will not be specifically mentioned in this specification.

[0096] According to one embodiment of the present invention, the plating layer of the plated steel sheet may be an aluminum or aluminum-based alloy plating layer. Further, according to one embodiment, the plating layer may be an alloyed aluminum-based plating layer.

[0097] In addition, as an embodiment of the present invention, the plating layer may contain Si, Mg, Fe in addition to Al, and may optionally contain Mn, Cr, Cu, Mo, Ni, Sb, Sn, Ti, Ca, Sr, Zn, etc. In the present invention, the thickness of the plating layer is not particularly limited, and the plating layer may have a thickness within a general range.

[0098] In one embodiment of the present invention, the plating layer contains, by weight, one or more elements selected from the group consisting of 5 to 11% Si, 5% or less Fe, and 5% or less Mg, with the remainder being Al and other impurities. If necessary, the above composition may further contain elements such as Mn, Cr, Cu, Mo, Ni, Sb, Sn, Ti, Ca, Sr, and Zn in a total amount of 30% or less.

[0099] In one embodiment of the present invention, the plating layer contains, by weight percent, one or more selected from the group consisting of 5.0 to 11.0% Si, 5.0% or less Fe, and 5.0% or less Mg, with the remainder being Al and other impurities. If necessary, the above composition may further contain elements such as Mn, Cr, Cu, Mo, Ni, Sb, Sn, Ti, Ca, Sr, and Zn in a total amount of 30.0% or less.

[0100] The components of the present invention will be described in detail below.

[0101] A member according to one embodiment of the present invention may include a base steel and a plating layer formed on a surface of the base steel.

[0102] The base steel according to one aspect of the present invention may have the same alloy composition as the base steel sheet of the coated steel sheet proposed in the present invention.

[0103] According to one embodiment of the present invention, the plating layer may be formed on at least one surface of the base steel. The plating layer of the member may have a composition in which the plating layer of the above-mentioned plated steel sheet and components of the base steel sheet, including Fe, are diffused and alloyed.

[0104] A component according to one embodiment of the present invention may include an antimony (Sb) enriched layer formed within the base iron.

[0105] The Sb-enriched layer of the present invention can be identified by analyzing the change in Sb content from any point in the coating layer toward the base steel in the thickness direction using a glow discharge spectrometer (GDS). This can be applied in the same manner as the method for identifying the Sb-enriched layer in a coated steel sheet proposed in the present invention. In one embodiment of the present invention, the antimony (Sb)-enriched layer may be formed directly below the interface where the base steel and the coating layer contact. According to one embodiment of the present invention, the interface between the base steel and the coating layer may refer to the point where the Al content is 15%.

[0106] In the embodiment of the present invention, when the content of elements in the thickness direction of the base steel is analyzed using a glow discharge spectrometer (GDS), the antimony (Sb) content in the antimony (Sb)-enriched layer is a maximum value (Sb max The carbon (C) content at the depth showing the pore size distribution (C0) can be 80% or less of the nominal carbon content (C0) of the base iron.

[0107] In the embodiment of the present invention, when the content of elements in the thickness direction of the base steel is analyzed using a glow discharge spectrometer (GDS), the antimony (Sb) content in the antimony (Sb)-enriched layer is a maximum value (Sb max The carbon (C) content at the depth showing the pore size distribution (C0) can be 80.0% or less of the nominal carbon content (C0) of the base iron.

[0108] The maximum Sb content in the Sb-enriched layer (Sb max ) influences the hardness of the surface structure and affects the bendability. On the other hand, when the Sb content reaches its maximum value (Sb max If the carbon content at the depth where the Sb concentration is greater than 80% of the nominal carbon content (C0), the hardness of the surface layer increases and the bendability may deteriorate. However, if the Sb concentration in the Sb-enriched layer exceeds the maximum value (Sb maxIf the carbon content at the depth where the carbon content is too low, the hardness of the surface layer will be insufficient, making it difficult to ensure fatigue resistance. Therefore, in one embodiment of the present invention, the lower limit may be set to 15%. In one embodiment of the present invention, the lower limit may be set to 15.0%.

[0109] FIG. 5 is a schematic diagram showing the Sb and C content profiles in the thickness direction from the interface in a member according to one embodiment of the present invention. The x-axis of FIG. 5 represents the depth (μm) from the interface between the base steel and the coating layer, and the y-axis represents the element content (wt%). As shown in FIG. 5, 80% of the nominal carbon content (C0) is 0.176%. Here, the nominal carbon content (C0) is 0.22%, which can be obtained by analyzing a certain thickness (depth) in the region of 1 / 4 to 3 / 4 of the thickness of the base steel using a glow discharge spectrometer (GDS), as described above. It can be seen that the carbon content at the depth where the Sb content is at its maximum is less than 80% of the nominal carbon content (C0).

[0110] The member according to one embodiment of the present invention may have an R value defined by the following relational expression 1 of 1.5 or more, and a B value defined by the following relational expression 2 of 0.01 or more.

[0111] A member according to an embodiment of the present invention may have an R value defined by the following relational expression 1 of 1.50 or more, and a B value defined by the following relational expression 2 of 0.010 or more. [Equation 1]

number

number

[0112] When a plated steel sheet is heated for hot forming, the Sb concentration in the Sb-enriched layer may become even more pronounced. During heat treatment for hot forming, the Sb-enriched layer effectively blocks permeating diffusible hydrogen. Diffusible hydrogen promotes grain boundary cracking when stress is generated. Therefore, reducing this diffusible hydrogen can improve bendability. That is, if the above-mentioned relationship is not satisfied, specifically, if the R value defined by Relationship 1 is less than 1.5 or the B value defined by Relationship 2 is less than 0.01, the diffusible hydrogen that permeates during hot forming may not be sufficiently blocked, resulting in poor crashworthiness. In one embodiment of the present invention, the R value defined by Relationship 1 may be 1.7 or greater. In another embodiment of the present invention, the B value defined by Relationship 2 may be 0.014 or greater. However, if the R value or B value is excessively high, the hardness of the surface layer of the component may be excessively high, resulting in poor bendability. Therefore, in one embodiment of the present invention, the upper limit of the R value may be limited to 6.4. In another embodiment of the present invention, the upper limit of the B value may be limited to 0.5.

[0113] In one embodiment of the present invention, the R value defined by Relational Formula 1 may be 1.70 or more. Also, in one embodiment of the present invention, the B value defined by Relational Formula 2 may be 0.0140 or more. In one embodiment of the present invention, the upper limit of the R value may be limited to 6.40. Also, in one embodiment of the present invention, the upper limit of the B value may be limited to 0.50.

[0114] In the member according to one aspect of the present invention, a region from the interface between the base steel and the coating layer to a depth of 45 to 100 μm in the thickness direction can have a softening rate (β) of 2 to 7%.

[0115] The region from the interface between the base steel and the coating layer to a depth of 45 to 100 μm in the thickness direction influences the hardness of the surface layer of the component, which may affect bendability.

[0116] If the softening rate in the region from 45 to 100 μm deep in the thickness direction is less than 2%, the hardness of the surface layer will be too high, which may reduce the effect of improving bendability.On the other hand, if the softening rate exceeds 7%, the hardness of the surface layer will be too low, which may cause problems such as deterioration of fatigue resistance.

[0117] The hardness softening rate in the present invention can be measured as shown in FIG. 6. FIG. 6 is a schematic diagram showing the hardness softening rate (β) profile at a depth of 45 to 100 μm from the interface in the thickness direction in a member according to one embodiment of the present invention. Specifically, Vickers hardness is measured using a 1 kg weight. The hardness inside the base steel is defined as the reference hardness (H0), which can be measured at a point 1 / 5 of the way through the thickness of the base steel. The y-axis in FIG. 6 represents the ratio (%) of the hardness value (H) at the corresponding position to the reference hardness value (H0), and the x-axis represents the distance (μm) in the thickness direction from the interface. As shown in FIG. 6, a rectangle was drawn with the y-axis ranging from 0 to 100% and the x-axis ranging from 45 to 100 μm deep from the interface. The square shows a hardness profile curve indicating the ratio of hardness values ​​depending on the depth from the interface, and the ratio of the area of ​​the upper region within the square of the hardness profile to the total area of ​​the square can be defined as the hardness softening rate (β, %). Within the range of 45 μm from the interface, the indentation or affected area of ​​the Vickers hardness test using 1 kg may be exposed to the plating layer or the outside, making it difficult to obtain an accurate hardness value. Therefore, in the present invention, a hardness profile was created at a depth of 45 to 100 μm and used to determine the hardness softening rate (β).

[0118] That is, the hardness softening rate (β) in the present invention means the percentage (%) of the area of ​​the upper region of the hardness profile curve to the total area of ​​a rectangle, where the horizontal axis represents the distance (μm) in the thickness (depth) direction from the interface between the base steel sheet and the coating layer, and the vertical axis represents the percentage (%) of the hardness value (H) at that position relative to the reference hardness value (H0).

[0119] In the member according to one aspect of the present invention, a region from the interface between the base steel and the coating layer to a depth of 45.0 to 100.0 μm in the thickness direction can have a softening rate (β) of 2.0 to 7.0%.

[0120] According to one embodiment of the present invention, the region of the member from the interface between the base steel and the coating layer to a depth of 50 μm in the thickness direction can contain less than 5 area % of ferrite in the microstructure.

[0121] The ferrite in the region extending from the interface between the base steel and the coating layer to a depth of 50 μm in the thickness direction can promote crack propagation. That is, if the ferrite content in this region is 5% or more, when stress is generated in the surface layer, local stress is concentrated in the relatively soft ferrite, which can promote crack propagation and deteriorate bendability and fatigue resistance.

[0122] In a member according to one embodiment of the present invention, a region extending from the interface between the base steel and the coating layer to a depth of 50 μm in the thickness direction may have a microstructure containing martensite as the main phase, less than 5 area % of ferrite, and the remainder upper and lower bainite. In the present invention, a phase having an area fraction of 50% or more of the total fraction of the microstructure may be considered to be the main phase.

[0123] If the martensite fraction is insufficient, the physical properties aimed at in the present invention may be insufficient.

[0124] According to one embodiment of the present invention, the region of the member from the interface between the base steel and the coating layer to a depth of 50.0 μm in the thickness direction can include less than 5.0 area % of ferrite in the microstructure.

[0125] In a member according to one embodiment of the present invention, a region from the interface between the base steel and the coating layer to a depth of 50.0 μm in the thickness direction may have a microstructure containing martensite as the main phase, less than 5.0 area% of ferrite, and the remainder upper and lower bainite. In the present invention, a phase having an area fraction of 50.0% or more of the total fraction of the microstructure may be considered to be the main phase.

[0126] The method for producing a plated steel sheet of the present invention will be described in detail below.

[0127] According to one aspect of the present invention, a plated steel sheet may be manufactured by annealing and plating a cold-rolled steel sheet having the above-described alloy composition. Here, the cold-rolled steel sheet may be manufactured by reheating, hot-rolling, coiling, cooling, and cold-rolling a steel slab having the above-described alloy composition.

[0128] reheating A steel slab satisfying the alloy composition according to one embodiment of the present invention can be reheated to a temperature range of 1050 to 1300°C.

[0129] If the reheating temperature is less than 1050°C, the slab structure is not sufficiently homogenized, making it difficult to redissolve precipitated elements when using them. On the other hand, if the temperature exceeds 1300°C, an excessive oxide layer is formed, which increases the manufacturing cost for removing the oxide layer and can cause surface defects after hot rolling.

[0130] hot rolling The reheated steel slab can be finish rolled at a temperature range of 800 to 950°C.

[0131] If the finish rolling temperature is less than 800°C, the rolling process will proceed in the two-phase region, resulting in the introduction of ferrite into the surface layer of the steel sheet, making it difficult to control the sheet shape. On the other hand, if the temperature exceeds 950°C, grain coarsening may occur.

[0132] Winding and cooling The rolled steel can be coiled and cooled in the temperature range of 500 to 700°C.

[0133] If the coiling temperature is less than 500°C, the tension during coiling is too high, which can lead to poor width shape of the hot rolled coil and equipment problems.On the other hand, if the temperature exceeds 700°C, excessive coarse carbides are formed, which promotes cracking when stress occurs in the hot-formed part, resulting in a problem of reduced crash resistance.

[0134] cold rolling The cooled steel can be cold rolled at a reduction ratio of 30 to 80% to produce a cold rolled steel sheet.

[0135] In the present invention, the cold rolling reduction is not particularly limited, but can be carried out within the range of 30 to 80% to obtain a predetermined target thickness.

[0136] annealing The cold-rolled steel sheet can be annealed in the temperature range of Ac1 to Ac3.

[0137] If the annealing temperature is below Ac1, the recrystallization of the cold-rolled structure may not be fully completed, resulting in poor sheet shape, and antimony may not be sufficiently concentrated, making it difficult to fully achieve the effects of the present invention in the final product. On the other hand, if the temperature exceeds Ac3, it may cause equipment problems in the annealing furnace and promote the formation of surface oxides, resulting in surface defects. According to one embodiment of the present invention, the lower limit of the annealing temperature may be 750°C. In another embodiment of the present invention, the upper limit of the annealing temperature may be limited to 860°C.

[0138] During the above annealing, the product of the annealing time and absolute humidity must be between 10,000 and 80,000 s·g / m 3 It can be.

[0139] During the annealing, the atmosphere and humidity can be adjusted using hydrogen gas, hydrogen-nitrogen mixed gas, etc. to form an oxidizing atmosphere, and it is important to control the annealing time and absolute humidity in the Ac1 to Ac3 temperature range to obtain an appropriate decarburization rate of the steel sheet.

[0140] Therefore, during the above annealing, the product of the annealing time and absolute humidity should be 10,000 to 80,000 s·g / m 3 It can be.

[0141] The product of annealing time and absolute humidity is 10,000 s·g / m 3 If the value is less than 80,000 s·g / m, the decarburization reaction due to internal oxidation will not occur sufficiently, making it difficult to achieve the desired decarburization rate, and the excessive carbon concentration in the component will prevent the improvement of crashworthiness from being expected. 3 If the temperature exceeds 100°C, excessive oxidation of the steel sheet surface may cause surface oxides to be generated, which may lead to surface defects during plating. According to one embodiment of the present invention, the annealing time may be 100 to 200 seconds. Also, according to one embodiment of the present invention, the absolute humidity is 100 to 400 g / m 3 It can be.

[0142] Furthermore, based on the surface temperature of the steel sheet, the average heating rate from room temperature to 500°C can be controlled to 2.7 to 10.0°C / s, the average heating rate in the 500 to 700°C range to 0.5 to 2.5°C / s, and the average heating rate from 700°C to the annealing temperature to 0.01 to 0.4°C / s.

[0143] The reason for limiting the average heating rate from room temperature to 500°C to 2.7 to 10.0°C / s in the steel sheet surface temperature standard is to ensure the formation of an Sb-enriched layer. If the average heating rate from room temperature to 500°C deviates from 2.7 to 10.0°C / s, specifically, if it is less than 2.7°C / s, the concentrated layer may not be sufficiently formed. If it exceeds 10°C / s, rapid heating increases the temperature non-uniformity across the width of the steel sheet, resulting in structural variations and line trouble. The surface temperature of the steel sheet in the 500 to 700°C range may affect the Sb concentration in the base steel. In other words, if the average heating rate in this range deviates from 0.5 to 2.5°C / s, the Sb-enriched layer may not be sufficiently formed. From the surface temperature of the steel sheet at 700°C to the target annealing temperature, which is a temperature at which an Sb-enriched layer is sufficiently formed on the base steel, the average heating rate is preferably 0.01 to 0.4°C / s to prevent the Sb-enriched layer and defects in the steel sheet surface that satisfy Relational Formulas 1 and 2.

[0144] plating The annealed cold-rolled steel sheet can be plated.

[0145] The plating bath according to one aspect of the present invention can be aluminum or an aluminum-based alloy.

[0146] According to one embodiment of the present invention, the plating bath composition may contain Si, Mg, Fe in addition to Al, and may also contain Mn, Cr, Cu, Mo, Ni, Sb, Sn, Ti, Ca, Sr, Zn, etc. The coating weight during plating is not particularly limited, and may be within a general range.

[0147] In one embodiment of the present invention, the composition of the plating bath may include, in weight percent, one or more selected from Si: 5 to 11%, Fe: 5% or less, and Mg: 5% or less, with the remainder being Al and other impurities.

[0148] According to one embodiment of the present invention, an alloying step may be included after plating. The alloying step is not particularly limited and may be carried out under normal conditions.

[0149] The method for manufacturing a member of the present invention will be described in detail below.

[0150] A component according to one aspect of the present invention can be manufactured by forming a blank from the plated steel sheet manufactured by the above-described method, heating, holding, forming, and cooling the blank.

[0151] Blank manufacturing The coated steel sheet proposed in the present invention can be produced in blanks for hot forming.

[0152] Heating and Maintenance The blank produced as described above can be heated in a temperature range of Ac3 to 975°C and maintained at that temperature for 10 to 1000 seconds.

[0153] If the blank heating temperature is below Ac3, the presence of untransformed ferrite in the two-phase region may make it difficult to ensure strength and impact resistance. On the other hand, if the heating temperature exceeds 975°C, excessive oxides may form on the surface of the part, making it difficult to ensure spot weldability and increasing manufacturing costs due to maintaining the high temperature. The heated blank then preferably has a heat treatment dwell time of 10 to 1,000 seconds within the above temperature range. If the dwell time is less than 10 seconds, it may be difficult to achieve uniform temperature distribution throughout the blank, resulting in material variations at different locations. On the other hand, if the dwell time exceeds 1,000 seconds, excessive oxides may form on the surface of the part and an interdiffusion layer may grow excessively, making it difficult to ensure spot weldability and increasing manufacturing costs.

[0154] Forming and Cooling The heated blank can be formed and cooled.

[0155] The heated blank can be transferred to a press and subjected to hot forming and die quenching at a cooling rate of 20°C / s or more to produce a final part. At a cooling rate of less than 20°C / s, ferrite phases can be introduced and formed at grain boundaries during cooling, which can degrade strength and impact resistance. According to one embodiment of the present invention, the blank can be cooled at 25°C / s or more after being formed.

[0156] The plated steel sheet of the present invention manufactured in this manner can maintain the surface hardness below a certain level and can maintain the blade life at a certain level or more when shearing for manufacturing blanks for hot forming, thereby reducing costs.

[0157] The member of the present invention manufactured in this manner has a product of tensile strength and bending angle of 80,000 MPa·° or more, a diffusible hydrogen content of 0.2 ppm or less, and excellent fatigue resistance and bendability.

[0158] In this invention, tensile strength (TS) * bending angle (BA) was used as an index for measuring crashworthiness. The bending angle, which is an index of crashworthiness, is affected by the tensile strength, and this tends to be inversely proportional. Therefore, as the product of tensile strength and bending angle (TS * BA) increases, crashworthiness increases. The BA value can be measured by bendability evaluation according to the VDA238-100 standard and is expressed as the bending angle converted to the maximum bending strength. [Example]

[0159] The present invention will be described in more detail with reference to the following examples. However, it should be noted that the following examples are intended to illustrate and explain the present invention in more detail, and are not intended to limit the scope of the present invention.

[0160] (Example) A 40 mm thick slab having a composition of 0.22C-0.25Si-1.25Mn-0.2Cr-0.03Al-0.03Ti-0.0025B and containing the Sb content disclosed in Table 1 below was produced by vacuum melting. The slab was heated to 1200°C and maintained for 1 hour, then hot rolled at a hot rolling finish temperature of 900°C, and coiled at a temperature of 600°C. After this, a pickling process was carried out, and cold rolling was carried out at a reduction ratio of 30 to 80% to produce a cold-rolled steel sheet. The cold-rolled steel sheet was annealed at temperatures Ac1 to Ac3, and the annealing time (s) and absolute humidity (g / m 3 The resulting product of the annealing time and absolute humidity is shown in Table 1. After annealing, the steel sheet was immersed in a plating bath containing Al-9%Si-2%Fe and trace amounts of impurities to perform plating.

[0161] [Table 1]

[0162] Table 2 below shows the Sb-enriched layer, microstructure, and decarburization rate of the manufactured coated steel sheets. First, a scanning electron microscope (SEM) was used to observe the structure directly below the interface and measure the pearlite area fraction. All specimens were observed to have ferrite as the remaining fraction, excluding the pearlite area fraction. Furthermore, a GDS850A (model name, manufactured by LECO) DC and RF equipment was used to measure the carbon decarburization rate in the region from the interface to a depth of 30 μm in the thickness direction of the base steel sheet. The decarburization rate (α) and the depth depending on the carbon content ratio were obtained from the carbon profile obtained using this equipment. A glow discharge spectrometer (GDS) was used to compare the values ​​of Relations 1 and 2 with the maximum antimony content (Sb max The carbon content at the depth where the unplated area was measured was shown. Furthermore, the plated steel sheet was visually inspected to determine whether or not there were unplated areas, and the number of unplated areas with an average diameter of 1 mm or more was 2 / m. 2 If it exceeds this, it is marked with a circle, and if it is less than this, it is marked with an ×.

[0163] [Table 2] [Equation 1]

number

number

[0164] The components were manufactured by hot forming using coated steel sheets with no uncoated areas. The heat treatment temperature and time for hot forming were 900°C and 360 seconds, and the transfer time from the heat treatment furnace to the forming press was 10 seconds.

[0165] Table 3 below shows the structure and properties of the parts manufactured by the hot forming process, measured using the same methods as described above. Vickers hardness was measured by applying a load of 1.0 kg to a region 45–100 μm deep from the interface between the coating layer and the base steel of the part, and the hardness softening rate (β) was calculated using Figure 6 and the method described above. The amount of diffusible hydrogen was also measured using a thermal desorption analysis (TDA) device (Bruker G8; model name). Specifically, the temperature was raised to 400°C at a rate of 20°C / min, and a diffusible hydrogen curve was measured for a certain period of time until a diffusible hydrogen peak was obtained. The diffusible hydrogen content in the steel was calculated by integrating this curve.

[0166] In addition, the ferrite area fraction within a 30 μm depth range from the interface between the base steel and the coating layer was measured using an optical microscope and is shown in Table 3. All specimens were observed to have martensite as the remaining area fraction, excluding the ferrite area fraction. Furthermore, fatigue limit strength was measured by performing fatigue tests over 10,000,000 cycles according to JIS Z2275. Fatigue resistance was confirmed by dividing the fatigue limit strength by the tensile strength by a value of 0.25 or greater, indicated as ◯, and a value less than 0.25, indicated as ×. Impact resistance was expressed as the product of tensile strength and bending angle. Tensile strength was measured using JIS-5 specimens via room-temperature tensile testing according to ISO 6892. The bending angle was calculated as the bending angle converted from the maximum bending strength specified in the VDA238-100 bendability evaluation method.

[0167] [Table 3] [Equation 1]

number

number

[0168] As shown in Tables 2 and 3, in the case of Examples 1 to 6, which satisfy the alloy composition and manufacturing conditions of the present invention, the characteristics proposed by the present invention are satisfied and the physical properties targeted by the present invention are also secured.

[0169] In Comparative Examples 1 and 2, the product of the annealing time and absolute humidity during annealing did not reach the range proposed by the present invention, and the decarburization rate of the plated steel sheet was outside the range proposed by the present invention. As a result, the hardness softening rate of the member was low, and the crash resistance was deteriorated due to excessive carbon concentration in the surface layer.

[0170] In Comparative Example 3, the Sb content in the steel was outside the range of the present invention, and the Sb-enriched layer was not sufficiently formed, which caused excessive internal oxidation during annealing. As a result, the hardness and softening of the component after heat treatment were excessively reduced, and a large amount of ferrite was formed in the surface layer, resulting in deterioration of fatigue resistance.

[0171] Figure 7 shows the carbon profile of a plated steel sheet according to one embodiment of the present invention. It can be seen from Figure 7 that Examples 1 and 3 of the present invention demonstrated sufficient decarburization control as proposed in the present invention, resulting in a high product of tensile strength and bending angle and a certain level of fatigue resistance. On the other hand, it can be seen that Comparative Example 1 did not achieve sufficient decarburization according to depth, and Comparative Example 3 demonstrated poor physical properties due to excessive decarburization caused by insufficient formation of an Sb-enriched layer.

[0172] 8 is a photograph of the microstructure observed by SEM immediately below the interface in the plated steel sheets of Example 3 according to an embodiment of the present invention and Comparative Example 3. It can be seen that in Example 3, 2.9% of pearlite was observed, while in Comparative Example 3, less than 1% of pearlite was observed.

[0173] 9 is an optical photograph of the interface between the plating layer and the base steel in the members of Example 3 of the present invention and Comparative Example 3. In the case of Example 3, ferrite was observed at less than 1%, but in the case of Comparative Example 3, ferrite was 7.3%, and the fatigue resistance characteristics aimed at in the present invention could not be ensured.

[0174] In Comparative Example 4, the product of the Sb content in the steel, the annealing time, and the absolute humidity was outside the range proposed by the present invention, and the amount of diffusible hydrogen in the part was excessive, which deteriorated bendability, so that the value of the product of tensile strength and bending angle, which is an indicator of crash resistance, did not reach the desired level.

[0175] In Comparative Examples 5 and 6, the product of the annealing time and absolute humidity during annealing exceeded the range of the present invention, which resulted in severe oxidation of the surface layer during annealing, resulting in the formation of Fe oxides on the surface, which in turn deteriorated the coating adhesion and caused the occurrence of uncoated areas.

[0176] Although the present invention has been described in detail with reference to the above embodiments, other embodiments are possible, and the spirit and scope of the claims set forth below should not be limited to the embodiments. [Explanation of symbols]

[0177] 1 plating layer 2 Sb concentrated layer In the 21Sb-enriched layer, the Sb content increases in the x-axis (+) direction. In the 22Sb-enriched layer, the Sb content decreases in the x-axis (+) direction. 3 Base steel sheet excluding the Sb-enriched layer 10 Average Sb content of plating layer 11 The last point of contact in the x-axis (+) direction between the average Sb content line of the plating layer and the Sb content line by GDS 30 Average Sb content of base steel sheet 31 The first contact point in the x-axis (+) direction between the average Sb content line of the base steel sheet and the Sb content line by GDS 100 Sb content line by GDS 200 The point where the Sb content is maximum in the Sb-enriched layer

Claims

1. The invention comprises a base steel sheet containing, by weight, carbon (C): 0.06 to 0.5% and antimony (Sb): 0.01 to 0.1% and a coating layer formed on a surface of the base steel sheet, The base steel sheet contains an antimony (Sb) concentrated layer therein, When the content of elements in the thickness direction of the base steel sheet is analyzed using a glow discharge spectrometer, the antimony (Sb) content in the antimony (Sb)-enriched layer is found to have a maximum value (Sb ｍａｘ The carbon (C) content at the depth where the nominal carbon content (C) of the base steel sheet is ０ ) 10 to 70% of the plated steel sheet.

2. The plated steel sheet according to claim 1, wherein a decarburization rate (α) of carbon (C) in a region from the interface between the base steel sheet and the plated layer to a depth of 30 μm in the thickness direction is 14 to 35%.

3. The plated steel sheet has a carbon (C) content that is greater than or equal to the nominal carbon content (C) at a depth of more than 1.5 μm and less than 6 μm from the interface between the base steel sheet and the plated layer in the thickness direction. ０ 3. The plated steel sheet according to claim 1, wherein there are points where the surface roughness is 50% of the surface roughness.

4. The plated steel sheet has a carbon (C) content that is greater than or equal to the nominal carbon content (C) at a depth of more than 6 μm and less than 15 μm from the interface between the base steel sheet and the plated layer in the thickness direction. ０ The plated steel sheet according to claim 1 , wherein there are points where the surface roughness is 80% of the surface roughness.

5. The plated steel sheet has an R value defined by the following relational expression 1 of 1.2 or more, The plated steel sheet according to claim 1 , wherein the B value defined by the following relational expression 2 is 0.008 or more. [Relationship 1] [Equation 1] [Relationship 2] [Equation 2] (In the formula, Sb ｍａｘ represents the maximum value of the Sb content in the Sb-enriched layer, and Sb ｃｏａｔ represents the average Sb content in the coating layer, expressed in weight percent. Furthermore, Δt represents the amount of Sb from the interface between the coating layer and the base steel sheet. ｍａｘ (The unit is μm.)

6. 6. The plated steel sheet according to claim 1, wherein a region from an interface between the base steel sheet and the plated layer to a depth of 10 μm in the thickness direction has a microstructure in which ferrite is a main phase and which contains 1 area % or more of pearlite.

7. 7. The plated steel sheet according to claim 1, wherein the base steel sheet contains carbon (C): 0.06 to 0.5%, antimony (Sb): 0.01 to 0.1%, silicon (Si): 0.001 to 2%, manganese (Mn): 0.1 to 4%, molybdenum (Mo): 1% or less, phosphorus (P): 0.05% or less, sulfur (S): 0.02% or less, aluminum (Al): 0.001 to 1%, chromium (Cr): 1% or less, nitrogen (N): 0.02% or less, titanium (Ti): 0.1% or less, boron (B): 0.01% or less, the balance being iron (Fe) and impurities.

8. The plated steel sheet according to claim 1 , wherein the plating layer is made of aluminum or an aluminum alloy.

9. The invention comprises a base iron containing, by weight, 0.06 to 0.5% carbon (C) and 0.01 to 0.1% antimony (Sb), and a plating layer formed on a surface of the base iron, The base iron contains an antimony (Sb)-enriched layer therein, When the content of elements in the thickness direction of the base steel is analyzed using a glow discharge spectrometer, the antimony (Sb) content in the antimony (Sb)-enriched layer is found to have a maximum value (Sb ｍａｘ The carbon (C) content at the depth where the carbon (C) content is the nominal carbon content (C ０ ) or less.

10. The component has a maximum antimony (Sb) content (Sb ｍａｘ The carbon (C) content at the depth where the carbon (C) content is the nominal carbon content (C ０ 10. The member according to claim 9, wherein the thickness is 15 to 80% of the total thickness.

11. The member has an R value defined by the following relational expression 1 of 1.5 or more, The member according to claim 9 or 10, wherein the B value defined by the following relational expression 2 is 0.01 or more. [Relationship 1] [Equation 3] [Relationship 2] [Equation 4] (In the formula, Sb ｍａｘ represents the maximum value of the Sb content in the Sb-enriched layer, and Sb ｃｏａｔ represents the average Sb content in the coating layer, expressed in weight percent. Furthermore, Δt is the distance from the interface where the coating layer and the base steel contact ｍａｘ (The unit is μm.)

12. The member according to any one of claims 9 to 11, wherein a region from an interface between the base steel and the plating layer to a depth of 45 to 100 µm in the thickness direction has a softening rate (β) of 2 to 7%.

13. 13. The member according to claim 9, wherein a region from an interface between the base steel and the coating layer to a depth of 50 μm in the thickness direction has a microstructure containing less than 5 area % of ferrite.

14. 14. The member according to claim 9, wherein a region from an interface between the base steel and the coating layer to a depth of 50 μm in the thickness direction has a microstructure containing martensite as a main phase, less than 5 area % of ferrite, and the remainder upper and lower bainite.

15. 15. The member according to any one of claims 9 to 14, wherein the base iron contains 0.06 to 0.5% carbon (C), 0.01 to 0.1% antimony (Sb), 0.001 to 2% silicon (Si), 0.1 to 4% manganese (Mn), 1% or less molybdenum (Mo), 0.05% or less phosphorus (P), 0.02% or less sulfur (S), 0.001 to 1% aluminum (Al), 1% or less chromium (Cr), 0.02% or less nitrogen (N), 0.1% or less titanium (Ti), 0.01% or less boron (B), the balance being iron (Fe) and impurities.

16. The member according to claim 9 , wherein the plating layer is made of aluminum or an aluminum alloy.

17. 17. The member of any one of claims 9 to 16, wherein the member has a product of tensile strength and bend angle of 80,000 MPa·° or greater.

18. 18. The component according to claim 9, wherein the component has a diffusible hydrogen content of 0.2 ppm or less.

19. A step of preparing a cold-rolled steel sheet containing, by weight, carbon (C): 0.06 to 0.5% and antimony (Sb): 0.01 to 0.1%; The cold-rolled steel sheet is １ ~Ac ３ annealing at a temperature range of plating the annealed cold-rolled steel sheet; During the annealing, the product of the annealing time and the absolute humidity is 10,000 to 80,000 s g / m ３ and during the annealing, based on the surface temperature of the steel sheet, an average heating rate from room temperature to 500°C is 2.7 to 10.0°C / s, an average heating rate in the 500 to 700°C range is 0.5 to 2.5°C / s, and an average heating rate from 700°C to the annealing temperature is 0.01 to 0.4°C / s.

20. During the annealing, the annealing time is 100 to 200 seconds, and the absolute humidity is 100 to 400 g / m ３ The method for producing a plated steel sheet according to claim 19, wherein

21. The cold-rolled steel sheet is reheating the steel slab to a temperature range of 1050-1300°C; finish rolling the reheated steel slab at a temperature in the range of 800 to 950°C; Coiling and cooling the rolled steel at a temperature in the range of 500 to 700°C; and The method for producing a plated steel sheet according to claim 19 or 20, comprising a step of cold rolling the cooled steel at a reduction ratio of 30 to 80%.

22. 22. The method for producing a plated steel sheet according to claim 19, wherein the steel slab contains carbon (C): 0.06 to 0.5%, antimony (Sb): 0.01 to 0.1%, silicon (Si): 0.001 to 2%, manganese (Mn): 0.1 to 4%, molybdenum (Mo): 1% or less, phosphorus (P): 0.05% or less, sulfur (S): 0.02% or less, aluminum (Al): 0.001 to 1%, chromium (Cr): 1% or less, nitrogen (N): 0.02% or less, titanium (Ti): 0.1% or less, boron (B): 0.01% or less, the balance being iron (Fe) and impurities.

23. The method for producing a plated steel sheet according to any one of claims 19 to 22, wherein the plating is performed with aluminum or an aluminum alloy.

24. Producing the plated steel sheet according to any one of claims 19 to 23 from a blank; The blank was ３ heating to a temperature range of 10 to 1000°C and maintaining for 10 to 1000 seconds; and forming and cooling the heated blank.

25. The method for manufacturing a member according to claim 24, wherein the cooling is performed at a cooling rate of 20°C / s or more.

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