KR102416968B1 - Member for automobile structure - Google Patents

Abstract

An embodiment of the present invention is a member for an automobile structure comprising a base steel plate and a plating layer covering at least one surface of the base steel plate, and has a tensile strength of 1350 MPa or more and a yield strength of 900 MPa or more, and the base steel plate is 80% or more martensite phase having an area fraction; an iron-based carbide located inside the martensite phase and having an area fraction of less than 5% based on the martensite phase; and precipitates distributed inside the base steel sheet, wherein the precipitates are present in 7.2×10 ³ to 1.65×10 ⁴ per unit area (100 μm ² ). Disclosed is a member for an automobile structure.

Description

Member for automobile structure

Embodiments of the present invention relate to a member for an automobile structure.

As environmental and fuel economy regulations are tightened around the world, the need for lighter vehicle materials is increasing. Accordingly, research and development of ultra-high-strength steel and hot stamping steel are being actively carried out. The double hot stamping process consists of heating/forming/cooling/trimming, and uses the phase transformation of the material and the change of the microstructure during the process.

Recently, studies to improve delayed fracture, corrosion resistance, and weldability occurring in a hot stamping member manufactured by a hot stamping process are being actively conducted. As a technology related thereto, there is Korean Patent Laid-Open Publication No. 10-2018-0095757 (Title of the Invention: Method of Manufacturing Hot Stamping Member) and the like.

No. 10-2018-0095757

SUMMARY Embodiments of the present invention provide a member for an automobile structure that prevents or minimizes delayed fracture due to residual hydrogen.

An embodiment of the present invention is a member for an automobile structure comprising a base steel plate and a plating layer covering at least one surface of the base steel plate, and has a tensile strength of 1350 MPa or more and a yield strength of 900 MPa or more, and the base steel plate is 80% or more martensite phase having an area fraction; an iron-based carbide located inside the martensite phase and having an area fraction of less than 5% based on the martensite phase; and precipitates distributed inside the base steel sheet, wherein the precipitates are present in 7.2×10 ³ to 1.65×10 ⁴ per unit area (100 μm ² ). Disclosed is a member for an automobile structure.

In this embodiment, 60% or more of the precipitates may have an equivalent circle diameter of 10 nm or less.

In the present embodiment, 4.5×10 ³ to 1.6×10 ⁴ of the precipitates having an equivalent circle diameter of 10 nm or less may be present per unit area (100 μm ² ).

In this embodiment, 25% or more of the precipitates may have an equivalent circle diameter of 5 nm or less.

In this embodiment, the average distance between two adjacent precipitates among the precipitates may be 0.4 μm or more and 0.8 μm or less.

In this embodiment, each of the precipitates may include a carbide of at least one of titanium (Ti), niobium (Nb), and vanadium (V), and may trap hydrogen.

In this embodiment, the iron-based carbide may have a diameter of less than 0.2 μm and a length of less than 10 μm.

In this embodiment, the martesite phase includes a Lath phase, and among the iron-based carbides, the area fraction of the iron-based carbide horizontal to the longitudinal direction of the Lath phase is the area of the iron-based carbide perpendicular to the longitudinal direction of the Lath phase may be greater than the fraction.

In this embodiment, among the iron-based carbides, the area fraction of the iron-based carbides forming an angle of 20° or less with the longitudinal direction of the lath phase may be 50% or more.

In this embodiment, among the iron-based carbides, the area fraction of the iron-based carbides forming an angle of 70° or more and 90° or less with the longitudinal direction of the lath may be less than 50%.

In this embodiment, the base steel sheet, carbon (C) 0.19 wt% to 0.25 wt%, silicon (Si) 0.1 wt% to 0.6 wt%, manganese (Mn) 0.5 wt% to 0.5 wt% based on the total weight of the base steel sheet 2.0 wt%, phosphorus (P) 0.03 wt% or less, sulfur (S) 0.003 wt% or less, chromium (Cr) 0.05 wt% to 0.6 wt%, boron (B) 0.0005 wt% to 0.005 wt%, the balance iron and It may contain unavoidable impurities.

In this embodiment, the plating layer may include aluminum (Al).

According to the embodiments of the present invention, it is possible to prevent hydrogen-delayed destruction of a member for an automobile structure and secure excellent mechanical properties.

1 is a cross-sectional view schematically illustrating a part of a member for an automobile structure according to an embodiment of the present invention.
FIG. 2 is a cross-sectional view schematically illustrating a portion A of FIG. 1 .
3 is a plan view showing a portion of the martensite phase in the base steel sheet of FIG. 2 .
4 is a plan view schematically illustrating a part of the base steel plate of FIG. 2 .
5 and 6 are plan views schematically illustrating a state in which hydrogen is trapped by the precipitates.
7 is a flowchart schematically illustrating an example of a method of manufacturing a member for an automobile structure of FIG. 1 .
8 is a graph showing the comparison of tensile strength and bending stress according to the coiling temperature of Examples and Comparative Examples of the present invention.
9 and 10 are images showing the results of the 4-point bending test according to the winding temperature of Examples and Comparative Examples.

Since the present invention can apply various transformations and can have various embodiments, specific embodiments are illustrated in the drawings and described in detail in the detailed description. Effects and features of the present invention, and a method for achieving them will become apparent with reference to the embodiments described below in detail in conjunction with the drawings. However, the present invention is not limited to the embodiments disclosed below and may be implemented in various forms.

In the following embodiments, terms such as first, second, etc. are used for the purpose of distinguishing one component from another, not in a limiting sense.

In the following examples, the singular expression includes the plural expression unless the context clearly dictates otherwise.

In the following embodiments, terms such as include or have means that the features or components described in the specification are present, and the possibility that one or more other features or components will be added is not excluded in advance.

In the following embodiments, when it is said that a part such as a film, region, or component is on or on another part, not only when it is directly on the other part, but also another film, region, component, etc. is interposed therebetween. Including cases where there is

In the drawings, the size of the components may be exaggerated or reduced for convenience of description. For example, since the size and thickness of each component shown in the drawings are arbitrarily indicated for convenience of description, the present invention is not necessarily limited to the illustrated bar.

In cases where certain embodiments may be implemented otherwise, a specific process sequence may be performed different from the described sequence. For example, two processes described in succession may be performed substantially simultaneously, or may be performed in an order opposite to the order described.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS Hereinafter, embodiments of the present invention will be described in detail with reference to the accompanying drawings, and the same or corresponding components are given the same reference numerals when described with reference to the drawings.

1 is a cross-sectional view schematically showing a part of a member for an automobile structure according to an embodiment of the present invention, FIG. 2 is a cross-sectional view schematically showing part A of FIG. 1, and FIG. 3 is martensite in the base steel plate of FIG. A plan view showing a part of the phase, and FIG. 4 is a plan view schematically showing a part of the base steel plate of FIG. 2 .

1 to 4 , the member 100 for an automobile structure according to an embodiment of the present invention includes at least one bent portion C, and may have a tensile strength of 1350 MPa or more and a yield strength of 900 MPa or more.

The member 100 for an automobile structure may include a base steel plate 110 and a plating layer 120 covering at least one surface of the base steel plate 110 .

The base steel sheet 110 may be a steel sheet manufactured by performing a hot rolling process and/or a cold rolling process on a slab cast to include a predetermined alloying element in a predetermined content. Such a base steel sheet 110 may exist as a full austenite structure at a hot stamping heating temperature, and then may be transformed into a martensitic structure upon cooling.

The average size of the initial austenite grains of the base steel sheet 110 may be 10 μm to 45 μm. Accordingly, the area of the grain boundary, which is a nucleation site of recrystallized grains, is increased, thereby promoting dynamic recrystallization behavior. In addition, the base steel sheet 110 is made of a component system that can have a microstructure including a martensite phase of 80% or more by area fraction. In addition, the base steel sheet 110 may include a bainite phase of less than 20% by area fraction.

For example, the base steel sheet 110 may include carbon (C), silicon (Si), manganese (Mn), phosphorus (P), sulfur (S), chromium (Cr), boron (B), and the remainder iron (Fe). and other unavoidable impurities. In addition, the base steel sheet 110 may further include at least one alloy element of titanium (Ti), niobium (Nb), and vanadium (V) as an additive. In addition, the base steel plate 110 may further include a predetermined amount of calcium (Ca).

Carbon (C) acts as an austenite stabilizing element in the base steel plate 110 . Carbon is a major element that determines the strength and hardness of the base steel sheet 110, and after the hot stamping process, to secure the tensile strength (eg, tensile strength of 1,350 MPa or more) of the base steel sheet 110, and to ensure hardenability added for the purpose of Such carbon may be included in an amount of 0.19 wt% to 0.25 wt% based on the total weight of the base steel sheet 110 . When the carbon content is less than 0.19 wt%, it is difficult to secure a hard phase (such as martensite), so it is difficult to satisfy the mechanical strength of the base steel sheet 110 . Conversely, when the content of carbon exceeds 0.25 wt%, brittleness of the base steel sheet 110 or a problem of reducing bending performance may occur.

Silicon (Si) acts as a ferrite stabilizing element in the base steel plate 110 . Silicon (Si) as a solid solution strengthening element improves the ductility of the base steel sheet 110, and suppresses the formation of carbides in the low-temperature region, thereby improving the carbon concentration in the austenite. In addition, silicon is a key element in hot rolling, cold rolling, hot pressing, homogenizing the structure (perlite, manganese segregation zone control), and fine dispersion of ferrite. Silicon serves as a martensitic strength heterogeneity control element to improve collision performance. Such silicon may be included in an amount of 0.1 wt% to 0.6 wt% based on the total weight of the base steel sheet 110 . When the content of silicon is less than 0.1wt%, it is difficult to obtain the above-described effect, and cementite formation and coarsening may occur in the final hot stamping martensite structure, and the uniforming effect of the base steel sheet 110 is insignificant and the V-bending angle is secured. can't do it Conversely, when the silicon content exceeds 0.6wt%, hot-rolling and cold-rolling loads increase, hot-rolled red scale is excessive, and plating properties of the base steel sheet 110 may be deteriorated.

Manganese (Mn) acts as an austenite stabilizing element in the base steel sheet 110 . Manganese is added to increase hardenability and strength during heat treatment. Such manganese may be included in 0.5 wt% to 2.0 wt% with respect to the total weight of the base steel sheet 110 . When the manganese content is less than 0.5 wt%, the hardenability effect is not sufficient, and the hard phase fraction in the molded article after hot stamping may be insufficient due to insufficient hardenability. On the other hand, when the content of manganese exceeds 2.0 wt%, ductility and toughness due to manganese segregation or pearlite bands may be reduced, which may cause deterioration in bending performance and may cause a heterogeneous microstructure.

Phosphorus (P) may be included in an amount greater than 0 and 0.03 wt% or less based on the total weight of the base steel plate 110 in order to prevent deterioration of the toughness of the base steel plate 110 . When the phosphorus content exceeds 0.03 wt%, an iron phosphide compound is formed to deteriorate toughness and weldability, and cracks may be induced in the base steel sheet 110 during the manufacturing process.

Sulfur (S) may be included in an amount of more than 0 and 0.003 wt% or less based on the total weight of the base steel sheet 110 . If the sulfur content exceeds 0.003 wt%, hot workability, weldability, and impact properties are deteriorated, and surface defects such as cracks may occur due to the formation of large inclusions.

Chromium (Cr) is added for the purpose of improving the hardenability and strength of the base steel sheet 110 . Chromium makes it possible to refine grains and secure strength through precipitation hardening. Such chromium may be included in 0.05wt% to 0.6wt% based on the total weight of the base steel plate 110 . When the content of chromium is less than 0.05wt%, the precipitation hardening effect is low, and on the contrary, when the content of chromium exceeds 0.6wt%, the Cr-based precipitates and the matrix solid solution increase to decrease toughness, and production cost increases due to cost increase may increase.

Boron (B) is added for the purpose of securing the hardenability and strength of the base steel sheet 110 by suppressing the transformation of ferrite, pearlite, and bainite to secure a martensitic structure. In addition, boron segregates at grain boundaries to increase hardenability by lowering grain boundary energy, and has an effect of grain refinement due to an increase in austenite grain growth temperature. Such boron may be included in an amount of 0.0005 wt% to 0.005 wt% based on the total weight of the base steel sheet 110 . When boron is included in the above range, it is possible to prevent the occurrence of brittleness at the hard phase grain boundary, and secure high toughness and bendability. When the content of boron is less than 0.0005wt%, the hardenability effect is insufficient. Conversely, when the content of boron exceeds 0.005wt%, the solid solubility is low, so that the hardenability is deteriorated or the hardenability is deteriorated or It may cause embrittlement at high temperatures, and toughness and bendability may be reduced due to the occurrence of hard phase intergranular embrittlement.

Meanwhile, other unavoidable impurities may include nitrogen (N) and the like. When a large amount of nitrogen (N) is added, the amount of dissolved nitrogen may increase to decrease impact properties and elongation of the base steel sheet 110 . Nitrogen may be included in an amount greater than 0 and 0.001 wt% or less based on the total weight of the base steel sheet 110 . When the nitrogen content exceeds 0.001 wt%, the impact properties and elongation of the base steel sheet 110 may be reduced.

The additive is a carbide generating element that contributes to the formation of the precipitates 112 in the base steel sheet 110 . Specifically, the additive may include at least one of titanium (Ti), niobium (Nb), and vanadium (V).

Titanium (Ti) forms precipitates 112 such as TiC and/or TiN at a high temperature, thereby effectively contributing to austenite grain refinement. The titanium may be included in an amount of 0.005 wt% to 0.05 wt% based on the total weight of the base steel sheet 110 . When titanium is included in the content range, poor performance and coarsening of precipitates can be prevented, and physical properties of the steel can be easily secured, and defects such as cracks can be prevented on the steel surface. On the other hand, when the content of titanium exceeds 0.05wt%, the precipitate 112 may be coarsened, and elongation and bendability may decrease.

Niobium (Nb) and vanadium (V) may increase strength and toughness according to a decrease in martensite packet size. Each of niobium and vanadium may be included in 0.01wt% to 0.1wt% based on the total weight of the base steel sheet 110 . When niobium and vanadium are included in the above range, the crystal grain refinement effect of the base steel sheet 110 is excellent in the hot rolling and cold rolling process, and it prevents cracks in the slab during steelmaking/playing, and brittle fracture of the product, and prevents the occurrence of brittle fracture. The generation of coarse precipitates can be minimized.

Calcium (Ca) may be added to control the shape of the inclusions. Such calcium may be included in an amount of 0.003 wt % or less based on the total weight of the base steel plate 110 .

The base steel sheet 110 is formed of a composite structure of a martensite phase having an area fraction of 80% or more and a bainite phase having an area fraction of less than 20%, thereby having a tensile strength of 1350 MPa or more and a yield strength of 900 MPa or more.

The martensitic phase is the result of diffusionless transformation of austenite γ under the initiation temperature (Ms) of martensitic transformation during cooling. Martensite may have a rod-shaped lath phase oriented in one direction (d) within each initial grain of austenite.

In addition, iron-based carbides may be generated inside the martensite phase during the manufacturing process of the plated steel sheet to be described later. The iron-based carbide may have a needle-like shape, and the needle-like iron-based carbide may have a diameter of less than 0.2 μm and a length of less than 10 μm. Here, the diameter of the acicular iron-based carbide may mean a minor axis length of the iron-based carbide, and the length of the needle-type iron-based carbide may mean a major axis length of the iron-based carbide.

If the diameter of the iron-based carbide is 0.2 μm or more or the length is 10 μm or more, it remains without melting even at a temperature of Ac3 or more in the annealing heat treatment process, and thus the bendability and yield ratio of the base steel sheet 110 may be reduced. On the other hand, when the diameter of the iron-based carbide is less than 0.2 μm and the length is less than 10 μm, the balance of strength and formability of the base steel sheet 110 may be improved.

Such iron-based carbides may have an area fraction of less than 5% based on the martensite phase. When the area fraction of the iron-based carbide is 5% or more based on the martensite phase, it may be difficult to secure the strength or bendability of the base steel sheet 110 .

In addition, among the iron-based carbides, the area fraction of the iron-based carbide (C1) horizontal to the longitudinal direction (d) of the lath is larger than the area fraction of the iron-based carbide (C2) perpendicular to the longitudinal direction (d) of the lath. The bendability of (110) may be improved. Here, 'horizontal' includes the longitudinal direction (d) of the lath and an angle of 20° or less, and 'vertical' includes forming an angle of 70° or more and 90° or less with the longitudinal direction (d) of the lath. can Specifically, the area fraction of the iron-based carbide (C1) forming an angle of 20° or less with the longitudinal direction (d) of the lath phase may be 50% or more, preferably 60% or more, and 70° with the longitudinal direction (d) of the lath phase. The area fraction of the iron-based carbide (C2) forming an angle of at least 90° may be less than 50%, preferably less than 40%.

Cracks generated during bending deformation may be generated as dislocations move in the martensite phase. In this case, it can be understood that the higher the local strain rate among the given plastic deformations, the higher the energy absorption for the plastic deformation of martensite, and thus the collision performance increases.

On the other hand, if the area fraction of the iron-based carbide (C1) horizontal to the longitudinal direction (d) of the lath is larger than the area fraction of the iron-based carbide (C2) perpendicular to the longitudinal direction (d) of the lath, the dislocation during bending deformation is In the process of internal movement, dynamic strain aging (DSA), ie, indentation dynamic strain aging, due to the local strain rate difference may appear. Indentation dynamic deformation aging is a concept of plastic deformation absorption energy, and since it means resistance to deformation, the more frequent the indentation dynamic deformation aging phenomenon, the better the resistance performance to deformation.

That is, according to the present invention, the area fraction of iron-based carbide (C1) forming an angle of 20° or less with the longitudinal direction (d) of the lath phase is 50% or more, and the length direction (d) of the lath phase and 70° or more 90° As the area fraction of iron-based carbide (C2) forming an angle of ° or less is less than 50%, indentation dynamic deformation aging may occur frequently. Collision performance can be improved.

On the other hand, since the bainite phase having an area fraction of less than 20% in the base steel sheet 110 has a uniform hardness distribution, it has an excellent balance between strength and ductility. However, since bainite is softer than martensite, in order to secure the strength and bending properties of the base steel sheet 110, it is preferable that the bainite has an area fraction of less than 20%.

The needle-type iron-based carbide described above may also be precipitated inside the bainite phase. Since the iron-based carbide inside bainite increases the strength of bainite and reduces the strength difference between bainite and martensite, the yield ratio and bendability of the base steel sheet 110 may be increased. In this case, the iron-based carbide may be present in less than 20% of the bainite phase based on the bainite phase. If the iron-based carbide is 20% or more based on the bainite phase, voids may be generated, which may lead to a decrease in bendability.

The plating layer 120 may be formed with an adhesion amount of 20-100 g/m ² based on one side. As an example, the plating layer 120 is immersed in the base steel sheet 110 in a plating bath containing at least one of molten aluminum and aluminum alloy at 600 to 800 ° C., and then cooled at an average cooling rate of 1 to 50 ° C./s. After forming the pre-coating layer, it may be formed by alloying by mutual diffusion between the base steel plate 110 and the pre-coating layer in the process of hot stamping the base steel plate 110 on which the pre-coating layer is formed.

In addition, after immersing the base steel sheet 110 in the plating bath, one or more of air and gas is sprayed on the surface of the base steel sheet 110 to wipe the hot-dip plated layer, and to adjust the plating adhesion amount of the precoat layer by controlling the spray pressure can

The precoat layer is formed on the surface of the base steel sheet 110 and is formed between a surface layer and a surface layer containing 80% by weight or more of aluminum (Al), and between the surface layer and the base steel sheet 110, aluminum-iron (Al-Fe) and aluminum- An alloying layer including an iron-silicon (Al-Fe-Si) intermetallic compound may be included. The alloying layer may contain 20 wt% to 70 wt% of iron (Fe). For example, the surface layer may contain 80 to 100% by weight of aluminum, and may have an average thickness of 10 μm to 40 μm.

On the other hand, when the base steel sheet 110 with the precoat layer is heated in a heating furnace to perform hot stamping for press-forming the base steel sheet 110 with the precoat layer at a high temperature, the base steel sheet 110 and the precoat layer in the heating process Inter-diffusion occurs between the layers, and the pre-coating layer may be alloyed to form the plating layer 120 .

Meanwhile, in the process of forming the plating layer 120 , hydrogen may be introduced into the base steel sheet 110 from the heating furnace, and hydrogen delayed destruction may be induced in the base steel sheet 110 by the hydrogen introduced into the base steel sheet 110 . can In addition, the member 100 for a vehicle structure according to the present invention may include at least one bent portion (C) depending on the applied position. The bent portion (C) is a portion that is excessively formed compared to a flat region, and the stress is relatively concentrated during press molding, and by using this concentrated stress as a driving force, a partial precipitate behavior change may occur, and the residual stress may be relatively large. . In this case, activated hydrogen that is not captured together with the residual stress of the bent portion C may affect the bent portion C, and thus the bent portion C may become a weak point of delayed hydrogen destruction. However, according to the present invention, at least some of the alloy elements included as additives exist as precipitates 112 in the base steel sheet 110 , and these precipitates 112 capture hydrogen distributed in the base steel sheet 110 . By doing so, the hydrogen delayed fracture resistance can be improved. The precipitates 112 may include at least one carbide of titanium (Ti), niobium (Nb), and vanadium (V), and the size equivalent to the original size of the precipitates 112, the spacing between the precipitates 112, the precipitates ( By controlling the number of 112), it is possible to improve the hydrogen trapping ability of the precipitate 112.

For example, by adjusting the coiling temperature (CT) range of the process conditions, it is possible to control the precipitation behavior such as the number of precipitates 112 , the average distance between the precipitates 112 , and the diameter of the precipitates 112 . In particular, the diameter of the precipitates 112 has a great influence on the improvement of the hydrogen delayed fracture characteristics. Hereinafter, with reference to FIGS. 5 and 6 , the difference in the effect of improving the delayed hydrogen destruction characteristics according to the diameter of the precipitates 112 will be described.

5 and 6 are plan views schematically illustrating a state in which hydrogen is trapped by the precipitates.

Specifically, FIG. 5 shows a state in which hydrogen is trapped in the precipitates 112 with a relatively large diameter, and FIG. 6 shows a state in which hydrogen is trapped in the precipitates 112 with a relatively small diameter. has been

5 , when the diameters of the precipitates 112 are relatively large, the number of hydrogen atoms trapped in one precipitate 112 increases. That is, the hydrogen atoms introduced into the base steel sheet 110 are not evenly dispersed, a plurality of hydrogen atoms are trapped in one hydrogen trap site, and a plurality of hydrogen atoms trapped in one hydrogen trap site are bonded to each other to form hydrogen. The probability of forming a molecule (H2) increases. The formed hydrogen molecules increase the probability of generating pressure inside the base steel sheet 110 , and as a result, the hydrogen delayed fracture characteristics of the base steel sheet 110 may be deteriorated.

On the other hand, when the diameters of the precipitates 112 are relatively small as shown in FIG. 6 , a larger number of the precipitates 112 may be dispersed in the base steel plate 110 than when the diameters of the precipitates 112 are large. In addition, since the size of the precipitate 112 is reduced, the number of hydrogen atoms that can be trapped by one precipitate 112 is also reduced, and as a result, the possibility of forming hydrogen molecules by bonding between hydrogen atoms is also reduced. . That is, since the hydrogen atoms introduced into the base steel sheet 110 are trapped at different hydrogen trap sites, the probability of generating internal pressure due to the hydrogen molecules is reduced, so that the hydrogen delayed destruction characteristics of the base steel sheet 110 can be improved.

That is, by controlling the equivalent size of the precipitates 112 , the spacing between the precipitates 112 , and the number of the precipitates 112 , the hydrogen trapping ability of the precipitates 112 can be improved. For example, by adjusting the coiling temperature (CT) range of the process conditions, it is possible to control the precipitation behavior such as the number of precipitates 112 , the average distance between the precipitates 112 , and the diameter of the precipitates 112 .

In an embodiment, the number of precipitates 112 formed in the base steel plate 110 may be controlled to satisfy a preset range. Specifically, the precipitates 112 are 72 pieces / μm ² (7.2 × 10 ³ pieces / 100 μm ² ) to 165 pieces / μm ² (1.65 × 10 ⁴ pieces / 100 μm ² ) in the base steel plate 110 may be included. . In particular, among the precipitates 112 distributed in the base steel sheet 110 , the fine precipitates having a diameter of 0.01 μm or less are 45 pieces/μm ² (4.5×10 ³ pieces/100 μm ² ) to 160 in the base steel sheet 110 . pieces/μm ² (1.6×10 ⁴ pieces/100 μm ² ) may be included.

When the number of precipitates 112 is formed within the above-described range, it is possible to secure the required tensile strength (eg, 1,350 MPa) after hot stamping and improve formability or bendability. For example, when the number of precipitates 112 having a diameter of 0.01 μm or less is less than 45/μm ² (4.5×10 ^3/100 μm ² ), the strength may be lowered, whereas 160/μm ² (1.6×10 ⁴ pieces/100 μm ² ) If it exceeds, moldability or bendability may be reduced.

In another embodiment, the average distance between the adjacent precipitates 112 may be controlled to satisfy a preset range. Here, the “average distance” may mean a mean free path of the precipitates 112 , and details of a method of measuring this will be described later.

Specifically, the average distance between the precipitates 112 may be 0.4 μm or more and 0.8 μm or less. When the average distance between the precipitates 112 is less than 0.4 μm, the formability or bendability may decrease, whereas, if it exceeds 0.8 μm, the strength may decrease.

In another embodiment, the equivalent circle diameter of the precipitates 112 may be controlled to satisfy a preset condition. Specifically, 60% or more of the precipitates 112 formed in the base steel plate 110 may be formed to have a diameter of 0.01 μm or less. In addition, 25% or more of the precipitates 112 formed in the base steel plate 110 may be formed to have a diameter of 0.005 μm or less. In addition, in an optional embodiment, the average diameter of all the precipitates 112 formed in the base steel sheet 110 may be 0.007㎛ or less.

Meanwhile, the precipitation behavior of the precipitates 112 may be measured by analyzing a transmission electron microscopy (TEM) image. Specifically, TEM images are acquired for arbitrary areas as many as a preset number for the specimen. From the obtained images, the precipitates 112 are extracted through an image analysis program, etc., and the number of precipitates 112 for the extracted precipitates 112, the average distance between the precipitates 112, and the diameter of the precipitates 112 etc can be measured.

In one embodiment, in order to measure the precipitation behavior of the precipitates 112, a surface replication method may be applied as a pretreatment to the measurement target specimen. For example, a one-step replica method, a two-step replica method, an extraction replica method, etc. may be applied, but the exemplary embodiment is not limited thereto.

In another embodiment, when the diameter of the precipitates 112 is measured, the diameter of the precipitates 112 may be calculated by converting the shape of the precipitates 112 into a circle in consideration of the non-uniformity of the shape of the precipitates 112 . Specifically, the area of the extracted precipitate 112 is measured using a unit pixel having a specific area, and the diameter of the precipitate 112 is calculated by converting the precipitate 112 into a circle having the same area as the measured area. can

In another embodiment, the average distance between the precipitates 112 may be measured through the aforementioned mean free path. Specifically, the average distance between the precipitates 112 may be calculated using the particle area fraction and the number of particles per unit length. For example, the average distance between the precipitates 112 may have a correlation as in Equation 1 below.

[Equation 1]

λ=(1-AA)/NL

(*?*: average distance between particles, AA: particle area fraction, NL: number of particles per unit length)

A method of measuring the precipitation behavior of the precipitates 112 is not limited to the above-described example, and various methods may be applied.

7 is a flowchart schematically illustrating an example of a method of manufacturing a member for an automobile structure of FIG. 1 .

7, the method for manufacturing a member for an automobile structure according to an embodiment of the present invention includes a reheating step (S100), a hot rolling step (S200), a cooling/winding step (S300), a cold rolling step (S400) ), an annealing heat treatment step (S500) and a plating step (S600) may be included. Meanwhile, although steps S100 to S600 are illustrated as independent steps in FIG. 7 , some of steps S100 to S600 may be performed in one process, and some may be omitted if necessary.

First, a slab in a semi-finished state to be subjected to the process of forming a base steel sheet is prepared. The slab contains 0.19 wt% to 0.25 wt% of carbon (C), 0.1 wt% to 0.6 wt% of silicon (Si), 0.5 wt% to 2.0 wt% of manganese (Mn), and 0.03 wt% of phosphorus (P) based on the total weight of the base steel sheet. wt% or less, sulfur (S) 0.003 wt% or less, chromium (Cr) 0.05 wt% to 0.6 wt%, boron (B) 0.0005 wt% to 0.005 wt%, the balance may include iron and unavoidable impurities. In addition, the slab may include at least one of titanium (Ti), niobium (Nb), and vanadium (V). For example, the content of titanium (Ti) may be 0.005 wt% to 0.05 wt%, and the content of each of niobium (Nb) and/or vanadium (V) may be 0.01 wt% to 0.1 wt%.

The reheating step ( S100 ) is a step of reheating the slab for hot rolling. In the reheating step (S100), the segregated components are re-dissolved during casting by reheating the slab secured through the continuous casting process in a predetermined temperature range.

The slab reheating temperature (SRT) may be controlled within a preset temperature range to maximize the austenite refining and precipitation hardening effects. In this case, the slab reheating temperature (SRT) range may be included in a temperature range (about 1,000° C. or higher) at which the additives (Ti, Nb, and/or V) are fully solidified during reheating of the slab. If the slab reheating temperature (SRT) is less than the total solid solution temperature range of additives (Ti, Nb, and/or V), the driving force required for microstructure control is not sufficiently reflected during hot rolling. No securement effect can be obtained.

In one embodiment, the slab reheating temperature (SRT) may be controlled to 1,200 ℃ to 1,300 ℃. When the slab reheating temperature (SRT) is less than 1,200 ° C, the components segregated during casting are not sufficiently re-dissolved, so it is difficult to see the effect of homogenizing the alloying elements greatly, and there is a problem in that it is difficult to see the effect of solid solution of titanium (Ti) significantly. On the other hand, the higher the slab reheating temperature (SRT) is, the higher the homogenization is, but when it exceeds 1,300 ° C, the austenite grain size increases, making it difficult to secure strength, and only the manufacturing cost of the steel sheet may increase due to the excessive heating process. .

The hot rolling step (S200) is a step of manufacturing a steel sheet by hot rolling the slab reheated in step S100 at a predetermined finishing delivery temperature (FDT) range. In one embodiment, the finish rolling temperature (FDT) range may be controlled to 840 ℃ to 920 ℃. When the finish rolling temperature (FDT) is less than 840℃, it is difficult to secure the workability of the steel sheet due to the occurrence of a mixed structure due to rolling in an abnormal region, and there is a problem in that workability is deteriorated due to microstructure non-uniformity as well as hot rolling due to rapid phase change. There may be a problem of marketability in the middle. Conversely, when the finish rolling temperature (FDT) exceeds 920° C., the austenite grains are coarsened. In addition, there is a risk that the TiC precipitates are coarsened to deteriorate the final member performance.

The cooling/winding step (S300) is a step of cooling and winding the steel sheet hot-rolled in step S200 within a predetermined coiling temperature (CT) range, and forming precipitates in the steel sheet. That is, in step S300, precipitates are formed by forming carbides of additives (Ti, Nb, and/or V) included in the slab. In one embodiment, the coiling temperature (CT) may be 700 ℃ to 780 ℃. The coiling temperature (CT) affects the redistribution of carbon (C). When the coiling temperature (CT) is less than 700 ° C., the low-temperature phase fraction increases due to overcooling, which may increase strength and increase the rolling load during cold rolling, and there is a problem in that ductility is rapidly reduced. Conversely, when the coiling temperature exceeds 780° C., there is a problem in that formability and strength deteriorate due to abnormal crystal grain growth or excessive crystal grain growth.

On the other hand, by controlling the coiling temperature (CT) range, it is possible to control the precipitation behavior of the precipitate. Experimental examples of changes in properties of the material 1 for hot stamping according to the winding temperature (CT) range will be described later with reference to FIGS. 8 to 10 .

Cold rolling step (S400) is a step of cold rolling after uncoiling (uncoiling) the steel sheet wound in step S300, pickling treatment. At this time, pickling is performed for the purpose of removing the scale of the wound steel sheet, that is, the hot rolled coil manufactured through the above hot rolling process. Meanwhile, in one embodiment, the rolling reduction during cold rolling may be controlled to 30% to 70%, but is not limited thereto.

The annealing heat treatment step (S500) is a step of annealing the cold rolled steel sheet in step S400 at a temperature of 700° C. or higher. In one embodiment, the annealing heat treatment includes heating the cold rolled sheet and cooling the heated cold rolled sheet at a predetermined cooling rate. For example, the heated cold-rolled sheet material is cooled at an average cooling rate of 5° C./s or more up to about 300° C., and then subjected to automatic tempering up to about 100° C. to control the size, area fraction, and directionality of the iron-based carbide.

The plating step ( S600 ) is a step of forming a plating layer on the annealed heat-treated steel sheet. In one embodiment, in the plating step ( S600 ), an Al-Si plating layer may be formed on the steel sheet annealed in the step S500 .

Specifically, in the plating step (S600), the steel sheet is immersed in a plating bath having a temperature of 650°C to 700°C to form a hot-dip plated layer on the surface of the steel sheet, and cooling the steel sheet on which the hot-dip plated layer is formed to form a plating layer. may include steps. In this case, the plating bath may include Si, Fe, Al, Mn, Cr, Mg, Ti, Zn, Sb, Sn, Cu, Ni, Co, In, Bi, etc. as an additive element, but is not limited thereto.

As such, by performing the hot stamping process on the steel sheet manufactured through steps S100 to S600, it is possible to manufacture a member for an automobile structure satisfying the required strength and bendability. In one embodiment, the member for an automobile structure manufactured to satisfy the above-described content conditions and process conditions may have a tensile strength of 1350 MPa or more and a yield strength of 900 MPa or more.

Hereinafter, the present invention will be described in more detail through Examples and Comparative Examples. However, the following Examples and Comparative Examples are for explaining the present invention in more detail, and the scope of the present invention is not limited by the following Examples and Comparative Examples. The following Examples and Comparative Examples can be appropriately modified and changed by those skilled in the art within the scope of the present invention.

8 is a graph showing the comparison of tensile strength and bending stress according to the winding temperature of Examples and Comparative Examples of the present invention, and FIGS. 9 and 10 are 4-point bending tests according to the winding temperature of Examples and Comparative Examples (4 point bending) These are images showing the results of test).

Example (CT700) and Comparative Example (CT800) are specimens prepared by hot stamping the hot stamping material (1) prepared by performing the above-described steps S100 to S600 with respect to the slab having the composition shown in Table 1 below. At this time, the example (CT700) and the comparative example (CT800) are specimens manufactured by applying the same content conditions and process conditions in the manufacturing process of the hot stamping material (1), but differentially applying only the coiling temperature (CT) as a variable. .

Ingredients (wt%) C Si Mn P S Cr B additive 0.19~0.25 0.1~0.6 0.5~2.0 0.03 or less 0.003 or less 0.05~0.6 0.0005~0.005 0.05 or less

Specifically, Example (CT700) is a specimen prepared by hot stamping the material for hot stamping (1) prepared by applying a coiling temperature (CT) of 700 °C, Comparative Example (CT800) is a coiling temperature of 800 °C It is a specimen manufactured by hot stamping the material (1) for hot stamping manufactured by applying (CT).

Meanwhile, FIG. 8 is a graph showing tensile strength and bending stress of Example (CT700) and Comparative Example (CT800).

Referring to FIG. 8 , in the case of tensile strength, the tensile strength of Example (CT700) is greater than that of Comparative Example (CT800), and the bending stress affecting impact properties is also comparable to that of Example (CT700). It can be seen that the improvement was compared with the bending stress of the example (CT800).

This is because, as can be seen in Table 2 below, in the case of Example (CT700) compared to Comparative Example (CT800), the amount of precipitation of the fine precipitates 20 increased and the hydrogen trapping ability was improved accordingly.

Table 2 below shows the measured values of the equilibrium precipitation amount and the amount of activated hydrogen in Example (CT700) and Comparative Example (CT800) and the results of a 4-point bending test. Here, the equilibrium precipitation amount means the maximum number of precipitates that can be precipitated when thermodynamically equilibrium is achieved, and the larger the equilibrium precipitation amount, the greater the number of precipitates precipitated. In addition, the amount of activated hydrogen means the amount of hydrogen excluding hydrogen trapped in the fine precipitates 20 among the hydrogen introduced into the steel sheet 10 .

Such an amount of activated hydrogen may be measured using a thermal desorption spectroscopy method. Specifically, while heating the specimen at a preset heating rate to increase the temperature, it is possible to measure the amount of hydrogen emitted from the specimen at a specific temperature or less. In this case, hydrogen emitted from the specimen at a temperature below a certain temperature is not trapped among the hydrogen introduced into the specimen and may be understood as activated hydrogen that affects delayed hydrogen destruction.

sample name Equilibrium Precipitation
(wt%) 4 point bending test result amount of activated hydrogen
(wppm) CT700 0.028 non-breaking 0.780 CT800 0.009 break 0.801

Table 2 shows the results of performing a four-point bending test on each of the samples with different equilibrium precipitation amounts of fine precipitates, and the amount of activated hydrogen measured using the thermal desorption spectroscopy method. indicates.

Here, the 4-point bending test is a test to check whether stress corrosion cracking occurs by applying a stress at a level below the elastic limit to a specific point on a specimen manufactured by reproducing the state in which the specimen is exposed to a corrosive environment way. At this time, stress corrosion cracking means a crack that occurs when corrosion and continuous tensile stress act simultaneously.

Specifically, the 4 point bending test results of Table 2 are results of checking whether fracture occurs by applying a stress of 1,000 MPa in air for 100 hours to each of the samples. In addition, the amount of activated hydrogen was measured using the above-described thermal desorption spectroscopy method, and the sample was heated at 350° C. or less while raising the temperature from room temperature to 500° C. at a heating rate of 20° C./min for each of the samples. It is a measure of the amount of hydrogen emitted from

Referring to Table 2, in the case of the equilibrium precipitation amount of the fine precipitates 20, the equilibrium precipitation amount of Example (CT700) was 0.028 wt%, and the equilibrium precipitation amount of the comparative example (CT800) was measured to be 0.009 wt%. That is, it can be confirmed that more fine precipitates 20 are formed in Example CT700 as compared to Comparative Example CT800, and thus more hydrogen trap sites can be provided.

On the other hand, in the case of the 4-point bending test result, the Example (CT700) did not break and the comparative example (CT800) did not break. In addition, in the case of the amount of activated hydrogen, the amount of activated hydrogen in Example (CT700) was about 0.780 wppm, and the amount of activated hydrogen in Comparative Example (CT800) was measured to be about 0.801 wppm. In this regard, it can be seen that the Example (CT700) having a relatively low amount of activated hydrogen was not broken, and the Comparative Example (CT800) having a relatively high amount of activated hydrogen was broken. It can be understood that the delayed hydrogen fracture characteristic of the embodiment (CT700) is improved compared to the comparative example (CT800).

That is, in Example (CT700) compared to Comparative Example (CT800), the amount of precipitation of the fine precipitates 20 increased, and, accordingly, the amount of activated hydrogen decreased. This means that the amount of hydrogen trapped inside in the embodiment (CT700) is increased compared to the comparative example (CT800), and as a result, it can be understood that the hydrogen delayed destruction characteristic is improved.

9 and 10 are images showing the results of performing a 4-point bending test with respect to Example (CT700) and Comparative Example (CT800), respectively.

Specifically, FIG. 9 is a result of a four-point bending test performed on Example (CT700), and FIG. 10 is a four-point bending test performed on Comparative Example (CT800) under the same conditions as Example (CT700). corresponds to a result.

9 and 10 , in the case of Example (CT700), the specimen was not fractured as a result of the four-point bending test, whereas in the case of Comparative Example (CT800), it was confirmed that the specimen was fractured.

This is, in the case of the embodiment (CT700) of FIG. 9, a specimen prepared by hot stamping the material for hot stamping (1) prepared by applying a coiling temperature (CT) of 700° C., and a microparticle having a diameter of 0.01 μm or less. The precipitates 20 are formed in 450 or more and 160 or less per unit area (μm ² ), and the average distance between the fine precipitates 20 satisfies 0.4 μm or more and 0.8 μm or less. Therefore, it can be seen that the embodiment CT700 efficiently disperses and traps hydrogen introduced into the steel sheet 10 to improve hydrogen delayed fracture characteristics, and improve tensile strength and bending characteristics.

On the contrary, in the case of the comparative example (CT800) of FIG. 10, as a specimen prepared by hot stamping the material for hot stamping (1) prepared by applying a coiling temperature (CT) of 800° C., the precipitation of fine precipitates 20 The amount is not sufficient, and the diameter of the fine precipitates 20 is coarsened, so that the probability of generating internal pressure due to hydrogen bonding increases. Therefore, it can be seen that Comparative Example CT800 cannot efficiently disperse and trap hydrogen introduced into the steel sheet 10, and the tensile strength, bending characteristics, and hydrogen delayed fracture characteristics are lowered.

That is, even though it is composed of the same components, due to the difference in the coiling temperature (CT), differences occur in strength, bendability, and hydrogen-delayed fracture characteristics of the hot stamping material 1 after the hot stamping process. This is because a difference occurs in the precipitation behavior of the fine precipitates 20 according to the coiling temperature (CT). Therefore, when the content conditions and process conditions according to the above-described embodiments of the present invention are applied, high strength can be secured, and bendability and hydrogen delayed fracture characteristics can be improved.

Table 3 below quantifies the tensile strength, bendability, and hydrogen delayed fracture characteristics according to the difference in the precipitation behavior of the fine precipitates 20 for a plurality of specimens. Specifically, Table 3 shows the measured values of the precipitation behavior (the number of fine precipitates, the average distance between the fine precipitates, the diameter of the fine precipitates, etc.) for a plurality of specimens, and the characteristics (tensile strength, bendability and amount of activated hydrogen) are described.

On the other hand, each of the plurality of specimens is heated to a temperature above Ac3 (the temperature at which the transformation from ferrite to austenite is completed) and cooled down to 300°C or less at a cooling rate of 30°C/s or more, and then the tensile strength, bendability, and activation number A small amount was measured.

At this time, the tensile strength and the amount of activated hydrogen were measured based on the above-described four-point bending test and thermal desorption spectroscopy method, and the bendability was determined by the German Automobile Industry Association (VDA: Verband Der). The V-bending angle is measured according to VDA238-100, the standard of Automobilindustrie).

In addition, the precipitation behavior of the fine precipitates (the number of fine precipitates, the average distance between the fine precipitates, the diameter of the fine precipitates, etc.) was measured through the above-described TEM image analysis. In addition, the precipitation behavior of the fine precipitates was measured for arbitrary areas having an area of 0.5 μm * 0.5 μm, and converted based on the unit area (100 μm ² ).

Psalter all
Number of fine precipitates
(pcs/100㎛ ² ) diameter
10nm or less
fine precipitate all
fine precipitate
average distance
(μm) diameter
5nm or less
fine precipitate Total fine precipitates
average diameter
(μm) after hot stamping
The tensile strength
(MPa) after hot stamping
bendability
(º) after hot stamping
amount of activated hydrogen
(wppm) Count
(pcs/100㎛ ² )
/ ratio(%) Count
(pcs/100㎛ ² )
/ ratio(%) A 7,210 4,514/
62.6% 0.7 1,810/
25.1% 0.0065 1389 53 0.782 B 7,257 6,511/
89.7% 0.66 2,750/
37.9% 0.0068 1402 58 0.791 C 8,290 5,301/
63.9% 0.54 2,329/
28.1% 0.0051 1396 59 0.795 D 11,935 10,712/
89.8% 0.51 7,889/
66.1% 0.0047 1416 61 0.779 E 14,872 14,152/
95.2% 0.52 11,332/
76.2% 0.0041 1440 59 0.761 F 16,499 9,907/
60% 0.57 4,158/
25.2% 0.0056 1511 57 0.725 G 16,499 15,989/
96.9% 0.41 4,158/
25.2% 0.0045 1538 63 0.791 H 9,833 7,520/
76.5% 0.8 4,602/
46.8% 0.0046 1405 54 0.779 I 13,900 13,792/
99.2% 0.4 11,523/
82.9% 0.0041 1409 58 0.751 J 10,031 8,458/
84.3% 0.59 5,166/
51.5% 0.007 1430 55 0.779 K 7,212 4,499/
62.4% 0.76 1,817/
25.2% 0.0068 1329 52 0.794 L 7,198 4,600/
63.9% 0.75 1,857/
25.8% 0.0063 1321 51 0.782 M 16,456 16,022/
97.4% 0.51 14,053/
85.4% 0.0042 1533 42 0.751 N 16,512 10,589/
64.1% 0.45 4,260/
25.8% 0.0044 1484 40 0.794 O 14,522 12,675/
87.3% 0.43 4,458/
30.7% 0.0071 1444 56 0.889 P 16,499 9,870/
59.8% 0.71 4,735/
28.7% 0.0058 1531 61 0.831 Q 7,201 5,080/
70.5% 0.72 1,786/
24.8% 0.0062 1401 54 0.824 R 16,480 15,070/
91.4% 0.44 4,104/
24.9% 0.0059 1510 65 0.864 S 16,435 14,989/
91.2% 0.39 14,808/
90.1% 0.0042 1479 46 0.791 T 13,384 12,598/
94.1% 0.81 12,006/
89.7% 0.0045 1340 54 0.786

Table 3 shows the measured values of the precipitation behavior of the fine precipitates (the number of fine precipitates, the average distance between the fine precipitates, the diameter of the fine precipitates, etc.) for specimens A to T, and the properties (tensile strength, bending) after hot stamping and the amount of activated hydrogen).

Specimens A to J of Table 3 are specimens prepared by hot stamping the material for hot stamping prepared through steps S100 to S600 by applying the above-described process conditions to a slab that satisfies the above-described content conditions (see Table 1). admit. That is, specimens A to J are specimens satisfying the precipitation behavior conditions of the fine precipitates described above. Specifically, in the specimens A to J, fine precipitates are formed in the steel sheet at 72 pieces/㎛ ² (7.2×10 ³ pieces/100 μm ² ) or more and 165 pieces/μm ² (1.65×10 ⁴ pieces/100 μm ² ) or less, and , the average diameter of all fine precipitates is 0.007 μm or less, 60% or more of the fine precipitates formed in the steel sheet have a diameter of 0.01 μm or less, and 25% or more have a diameter of 0.005 μm or less, and the average distance between the fine precipitates is 0.4 µm or more and 0.8 µm or less are satisfied.

It can be seen that the specimens A to J satisfying the precipitation behavior conditions of the present invention have improved tensile strength, bendability and hydrogen delayed fracture characteristics. Specifically, for specimens A to J, the tensile strength after hot stamping satisfies 1,350 MPa or more, the bendability after hot stamping satisfies 50 degrees or more, and the activated hydrogen content after hot stamping satisfies 0.8 wppm or less .

On the other hand, specimens K to T are specimens that do not satisfy at least some of the precipitation behavior conditions of the above-mentioned fine precipitates, and the tensile strength, bendability and/or delayed hydrogen fracture characteristics are inferior compared to specimens A to J. can be checked

In the case of specimen K, the number of fine precipitates with a diameter of 10 nm or less was 4,499. This is less than the lower limit of the condition for the number of fine precipitates with a diameter of 10 nm or less. Accordingly, it can be confirmed that the tensile strength of specimen K is only 1,329 MPa, which is relatively low.

In the case of specimen L, the total number of fine precipitates was 7,198. This is less than the lower limit of the total number of fine precipitates. Accordingly, it can be confirmed that the tensile strength of specimen L is only 1,321 MPa, which is relatively low.

In the case of specimen M, the number of fine precipitates with a diameter of 10 nm or less was 16,022. This exceeds the upper limit of the condition for the number of fine precipitates with a diameter of 10 nm or less. Accordingly, it can be confirmed that the bendability of specimen M is only 42 degrees, which is relatively low.

In the case of specimen N, the total number of fine precipitates is 16,512. This exceeds the upper limit of the total number of fine precipitates. Accordingly, it can be confirmed that the bendability of specimen N is only 40 degrees, which is relatively low.

In the case of specimen O, the average diameter of the total fine precipitates was 0.0071 μm. This exceeds the upper limit of the overall microprecipitate average diameter condition. Accordingly, the amount of activated hydrogen in specimen O was measured as a relatively high 0.889 wppm, confirming that the delayed hydrogen fracture characteristics were relatively deteriorated.

In the case of specimen P, the proportion of fine precipitates with a diameter of 10 nm or less was 59.8%. This is less than the lower limit of the ratio condition of fine precipitates with a diameter of 10 nm or less. Accordingly, the amount of activated hydrogen in the specimen P was measured as a relatively high 0.831 wppm, confirming that the delayed hydrogen fracture characteristics were relatively deteriorated.

In the case of specimen Q, the proportion of fine precipitates with a diameter of 5 nm or less was 24.8%. This is less than the lower limit of the ratio condition of fine precipitates with a diameter of 5 nm or less. Accordingly, the amount of activated hydrogen in the specimen Q was measured as a relatively high 0.824 wppm, confirming that the delayed hydrogen fracture characteristics were relatively deteriorated.

In the case of specimen R, the proportion of fine precipitates with a diameter of 5 nm or less was 24.9%. This is less than the lower limit of the ratio condition of fine precipitates with a diameter of 5 nm or less. Accordingly, the amount of activated hydrogen in the specimen R was measured as a relatively high 0.864 wppm, confirming that the delayed hydrogen fracture characteristics were relatively deteriorated.

In the case of specimen S, the average distance of all fine precipitates was 0.39 μm. This is less than the lower limit of the average distance condition of all fine precipitates. Accordingly, it can be confirmed that the bendability of specimen S is only 46 degrees, which is relatively low.

In the case of specimen T, the average distance of all fine precipitates was 0.81 μm. This exceeds the upper limit of the overall microprecipitate average distance condition. Accordingly, it can be confirmed that the tensile strength of the specimen T is only 1,340 MPa, which is relatively low.

As a result, the material for hot stamping manufactured by the method for manufacturing the material for hot stamping to which the content conditions and process conditions of the present invention are applied as described above satisfies the precipitation behavior conditions of the fine precipitates after hot stamping, and such fine precipitates It was confirmed that the tensile strength, bendability, and hydrogen-delayed fracture characteristics were improved in the hot stamping products satisfying these precipitation behavior conditions.

Table 5 below shows the results of measuring the V-bending angles of Example (CT700) and Comparative Example 2 (CT700). Comparative Example 2 is a specimen manufactured in the same manner as in Example (CT700), but is a case in which auto-tempering is not performed from 300°C to 100°C in the annealing heat treatment step. In the case of the following specimens, by observing the microstructure at a point 1/4 of the thickness of the specimen from the surface of the specimen, the average size, area fraction, and angle of the acicular carbides in martensite in the longitudinal direction of the lath are 20° or less. The area fraction of was measured.

carbide
average diameter
(μm) carbide
average length
(μm) carbide
area fraction
(%) Area fraction (%) of needle-shaped carbide whose angle with the longitudinal direction of the lath is 20° or less V-bending
(°) Example 0.17 7.3 4.5 57 50 Comparative Example 2 0.17 5.0 5.2 45 44

As can be seen from Table 4, in the case of the embodiment in which auto-tempering was performed from 300° C. to 100° C. during the manufacturing process of the plated steel sheet, the needle-type iron-based carbide in martensite was less than 5% of the area based on the martensite phase. The area fraction of the needle-shaped carbide having a fraction, having a diameter of less than 0.2 μm and a length of less than 10 μm, and having an angle of less than 20° with the longitudinal direction of the lath is formed to be 50% or more, thereby reducing the V-bending angle to 50 As it can be seen that more than ° can be secured, it can be confirmed that the tensile strength and bendability are improved. In contrast, in Comparative Example 2, although the iron-based carbide was formed to be relatively small in size, more iron-based carbide was formed in a direction perpendicular to the longitudinal direction of the lath, and it can be seen that the bendability of the plated steel sheet was lowered compared to that of Example. . That is, it can be confirmed that the area fraction of the needle-shaped carbide having an angle of 20° or less with the longitudinal direction of the lath is 50% or more, thereby improving tensile strength and bendability.

Although the present invention has been described with reference to the embodiments shown in the drawings, which are merely exemplary, those skilled in the art will understand that various modifications and equivalent other embodiments are possible therefrom. Therefore, the true technical protection scope of the present invention should be determined by the technical spirit of the appended claims.

Claims (12)

It is a member for an automobile structure comprising a base steel plate and a plating layer covering at least one surface of the base steel plate, and has a tensile strength of 1350 MPa or more and a yield strength of 900 MPa or more,
The base steel sheet is, based on the total weight of the base steel sheet, carbon (C) 0.19 wt% to 0.25 wt%, silicon (Si) 0.1 wt% to 0.6 wt%, manganese (Mn) 0.5 wt% to 2.0 wt%, phosphorus ( P) 0.03 wt% or less, sulfur (S) 0.003 wt% or less, chromium (Cr) 0.05 wt% to 0.6 wt%, boron (B) 0.0005 wt% to 0.005 wt%, the balance iron and unavoidable impurities,
The base steel plate,
a martensite phase having an area fraction of 80% or more;
an iron-based carbide having a diameter of less than 0.2 μm and an area fraction of less than 5% based on the martensite phase, positioned inside the martensite phase in the form of needles having a length of less than 10 μm; and
Precipitates that are distributed within the base steel sheet and contain at least one carbide of titanium (Ti), niobium (Nb), and vanadium (V);
The precipitates are 7.2×10 ³ to 1.65×10 ⁴ or more per unit area (100 μm ² ) for automobile structure members. According to claim 1,
60% or more of the precipitates have an equivalent circle diameter of 10 nm or less. 3. The method of claim 2,
A member for an automobile structure in which the precipitates having an equivalent circle diameter of 10 nm or less are present in 4.5×10 ³ to 1.6×10 ⁴ per unit area (100 μm ² ). 3. The method of claim 2,
25% or more of the precipitates have an equivalent circle diameter of 5 nm or less. According to claim 1,
An average distance between two adjacent precipitates among the precipitates is 0.4 μm or more and 0.8 μm or less. delete delete According to claim 1,
The martensite phase includes a Lath phase,
Among the iron-based carbides, an area fraction of iron-based carbides horizontal to the longitudinal direction of the lath phase is greater than an area fraction of iron-based carbides perpendicular to the longitudinal direction of the lath phase. 9. The method of claim 8,
Among the iron-based carbides, an area fraction of iron-based carbides forming an angle of 20° or less with the longitudinal direction of the lath phase is 50% or more for automobile structures. 9. The method of claim 8,
Among the iron-based carbides, the area fraction of the iron-based carbides forming an angle of 70° or more and 90° or less with the longitudinal direction of the lath is less than 50% of an automobile structure member. delete According to claim 1,
The plating layer is a member for an automobile structure including aluminum (Al).

Images