KR20210080676A - High-strength steel sheet with excellent ductability and formability and method of manufacturing the same - Google Patents

Abstract

A method to manufacture a high-strength steel sheet in accordance with an aspect comprises the steps of: preparing a hot rolled steel sheet which includes 0.1-0.3 wt% of carbon (C), 0.5-2.5 wt% of silicon (Si), 1.0-3.0 wt% of manganese (Mn), equal to or less than 0.01 wt% of phosphorus (P), equal to or less than 0.02 wt% of sulfur (S), 0.1-1.0 wt% of aluminum (Al), 0.005-0.1 wt% of titanium (Ti), 0.005-0.1 wt% of niobium (Nb), and the remainder consisting of iron (Fe) and inevitable impurities, wherein the content of titanium (Ti) ([Ti]) and the content of nitrogen (N) ([N]) satisfy [Ti] >= 3.4 x [N]; winding the hot rolled steel sheet at a temperature between 500 and 650℃; cold rolling the steel sheet at a reduction rate of 55-70%; annealing and thermally treating the cold rolled steel sheet at temperature between 800 and 900℃; and thermally treating the steel sheet at temperature between 450 and 650℃ for alloying.

Description

High-strength steel sheet with excellent ductility and formability {HIGH-STRENGTH STEEL SHEET WITH EXCELLENT DUCTABILITY AND FORMABILITY AND METHOD OF MANUFACTURING THE SAME}

The present invention relates to a steel sheet and a method for manufacturing the same, and more particularly, to a high-strength steel sheet having excellent ductility and formability of 780 MPa class suitable for automobile parts and a method for manufacturing the same.

Recently, as the use of ultra-high-strength steel sheets for the purpose of improving fuel efficiency and stability in the automobile market is gradually increasing, the demand for ultra-high-strength steel with high formability is increasing. In response to these demands, various developments have been made on ultra-high strength steel with high ductility and excellent hole expandability.

In general, dual phase steel (DP) is a steel with an excellent balance of strength and ductility, and although its utility is high, there is a limit to securing formability in machining difficult-to-form parts. This is because the difference in hardness between the phases between ferrite and martensite constituting the ideal steel is large. In the case of TRIP (Transformed Induced Plasticity) steel, the hard structure is composed of bainite and martensite, which has relatively superior formability compared to DP steel. . In addition, it is characterized by having a higher elongation compared to an ideal-structured steel due to the improvement of work hardenability by strain-induced transformation of Retained Austenite (RA).

In the case of TRIP steel, it is divided into various types according to the composition and process technology, but basically, a large amount of silicon (Si) must be added to secure retained austenite. However, when 1.5 wt% or more of silicon (Si) is added, various surface quality problems may occur in the production of plating materials. Q&P (Quenching & Partitioning) steel was developed as a way to secure the stability of retained austenite, but for the application of the technology, there is a problem in that the production cost increases due to new facility investment and reheating processes during production.

In Japanese Patent Laid-Open No. 61-157625, a steel sheet containing 0.12 to 0.7% of C, 0.4 to 1.4% of Si, and 0.2 to 2.5% of Mn by weight, by weight, is described as a two-phase region of ferrite-austenite or austenite. Disclosed is a method for manufacturing a high-strength cold-rolled steel sheet in which austenite is stabilized by heating to a single-phase region, annealing, cooling, and performing austenite treatment maintained at 350 to 500°C. It is possible to manufacture a cold-rolled steel sheet excellent in ductility and cold workability by such a technique. However, in the manufacture of a hot-dip galvanized steel sheet, since there is a time limit for sufficiently performing austenite treatment, stable retained austenite as in the prior art cannot be obtained.

In Unexamined Patent Publication No. 10-2019-0040018, in wt%, C: 0.03% to 0.70%, Si: 0.25% to 2.50%, Mn: 1.00% to 5.00%, P: 0.100% or less, S: 0.010% or less, sol .Al: 0.001% to 2.500, N: 0.020% or less, the balance consists of iron and impurities, the metal structure contains more than 5.0 volume% of retained austenite, more than 5.0 volume% of tempered martensite, , a method for manufacturing a high-strength cold-rolled steel sheet excellent in local ductility satisfying that the amount of C in retained austenite is 0.85 mass% or more is disclosed. However, hole expandability for forming difficult-to-form parts is not considered. In addition, when Si is excessively added to stabilize retained austenite, the surface wettability is lowered by the Si oxide, and plating properties may be deteriorated.

In Patent Publication No. 10-2014-7010265, a steel sheet containing 0.10 to 0.6% C, 0 to 3% Si, 0.5 to 3.0% Mn, and 0.005 to 3.0% by weight, Al: 0.005 to 3.0% by weight, satisfies the above component composition A method for manufacturing a TRIP steel sheet excellent in workability by heating a steel material to be used in a ferrite-austenite two-phase region or austenite single-phase region, cooling it to a temperature below Ms, and then reheating it to a temperature of 350°C or higher and 490°C or lower was disclosed. However, if the Mn content is high, the slab heating temperature and time are increased to prevent Mn segregation, and the production cost increases. In order to go through the reheating process after cooling, the equipment investment cost and process cost for this increase.

In Laid-Open Patent Publication No. 10-2014-0057177, it contains No. by weight, C: 0.1 to 0.3%, Si: 0.1 to 2.0%, Mn: 1.5 to 3.0%, Al: 0.005 to 1.5%, and the cold-rolled steel sheet is heated to a temperature of Ac3 or higher. After the primary annealing step of annealing and cooling, and after the primary annealing, heating and maintaining at a temperature in the range of Ac1 to Ac3, then cooling to Ms to Mf at a cooling rate of 20° C./s or more, and then reheating to Ms or more to 1 A method of manufacturing a hot-dip galvanized steel sheet having excellent ductility including a secondary annealing step of cooling after holding for more than a second is disclosed. However, the primary and secondary annealing processes cause an increase in manufacturing cost, and additional equipment investment is required to go through these processes.

The problem to be solved by the present invention is to provide a steel sheet having both ductility and formability as high-strength steel having tensile strength (TS): 780 MPa or more, and a method for manufacturing the same.

High-strength steel sheet according to an aspect of the present invention is, by weight%, carbon (C): 0.1 to 0.3%, silicon (Si): 0.5 to 2.5%, manganese (Mn): 1.0 to 3.0%, phosphorus (P): 0.01 % or less, sulfur (S): 0.02% or less, aluminum (Al): 0.1 to 1.0%, titanium (Ti): 0.005 to 0.1%, niobium (Nb): 0.005 to 0.1%, remaining iron (Fe) and other unavoidable Including impurities, the titanium (Ti) and nitrogen (N) content is controlled by the following formula, [Ti] ≥ 3.4 x [N]

([Ti] and [N] represent the content of titanium and nitrogen in weight percent in the total weight of the steel sheet, respectively), tensile strength (TS): 780 MPa or more, elongation (EL): 25% or more, and hole expandability ( HER): It is characterized in that it has 30% or more of physical properties.

The steel sheet may have a microstructure composed of polygonal ferrite: 30 to 70%, bainite: 5 to 40%, martensite: 0 to 10%, and retained austenite: 5 to 15%.

The steel sheet has an average grain size of 1 to 5 μm, an aspect ratio of polygonal ferrite is 3 or less, an aspect ratio of bainite and martensite is 3 or more, and an aspect ratio of retained austenite is 3 or less. It is desirable to have a tissue with at least 70% phase.

A method of manufacturing a high-strength steel sheet according to an aspect of the present invention, (a) by weight, carbon (C): 0.1 to 0.3%, silicon (Si): 0.5 to 2.5%, manganese (Mn): 1.0 to 3.0% , phosphorus (P): 0.01% or less, sulfur (S): 0.02% or less, aluminum (Al): 0.1 to 1.0%, titanium (Ti): 0.005 to 0.1%, niobium (Nb): 0.005 to 0.1%, remainder Preparing a hot-rolled steel sheet containing iron (Fe) and other unavoidable impurities; (b) winding the hot-rolled steel sheet in a temperature range of 500 to 650°C; (c) cold rolling the steel sheet at a reduction ratio of 55 to 70%; (d) annealing the cold-rolled cold-rolled steel sheet in a temperature range of 800 to 900°C; and (e) heat-treating the steel sheet in a temperature range of 450 to 650° C. for alloying.

The step (a) comprises the steps of reheating the steel slab at 1,150 ~ 1,300 ℃; And finishing rolling temperature (FDT): may include the step of hot rolling at 800 ~ 1,100 ℃.

In step (b), the method may further include cooling the hot-rolled steel sheet at a cooling rate of 1 to 100° C./s before the winding.

The step (e) comprises the steps of: (e-1) raising the temperature of the steel sheet to 800 to 900°C at a temperature increase rate of 1 to 10°C and then maintaining the steel sheet for 60 to 600 seconds; (e-2) maintaining the steel sheet for 10 to 100 seconds after cooling to 400 to 550 °C at a cooling rate of 3 to 20 °C/s; and (e-3) heat-treating the steel sheet in a temperature range of 450 to 650°C.

After step (e), the steel sheet has physical properties such as tensile strength (TS): 780 MPa or more, elongation (EL): 25% or more, and hole expansion (HER): 30% or more, polygonal ferrite: 30 to 70%, It is preferable to have a microstructure composed of bainite: 5 to 40%, martensite: 0 to 10%, and retained austenite: 5 to 15%.

According to the present invention, it can be expected to replace with high-strength steel for difficult-to-form parts among automobile parts by simultaneously securing elongation and hole expandability through control of alloy composition and process conditions.

1 is a process flowchart schematically illustrating a method of manufacturing a high-strength steel sheet having excellent ductility and formability according to an embodiment of the present invention.

Hereinafter, the present invention will be described in detail with reference to the accompanying drawings so that those of ordinary skill in the art to which the present invention pertains can easily practice it. The present invention may be embodied in several different forms, and is not limited to the embodiments described herein. The same reference numerals are assigned to the same or similar components throughout this specification. In addition, detailed descriptions of well-known functions and configurations that may unnecessarily obscure the gist of the present invention will be omitted.

High-strength steel sheet with excellent ductility and formability

The high-strength steel sheet excellent in ductility and formability of the present invention is, by weight, carbon (C): 0.1 to 0.3%, silicon (Si): 0.5 to 2.5%, manganese (Mn): 1.0 to 3.0%, phosphorus (P) ): 0.01% or less, Sulfur (S): 0.02% or less, Aluminum (Al): 0.1 to 1.0%, Titanium (Ti): 0.005 to 0.1%, Niobium (Nb): 0.005 to 0.1%, remaining iron (Fe) and other unavoidable impurities.

Hereinafter, the role and content of each component included in the high-strength steel sheet having excellent ductility and formability according to an embodiment of the present invention will be described in detail.

Carbon (C): 0.1 to 0.3 wt%

Carbon (C) is the most economical and effective alloying component to secure the strength of steel. Carbon (C) was used as a main element to increase strength, such as suppression of austenite phase transformation during cooling, and to stabilize retained austenite. When carbon (C) is added in 0.1 wt% or less, it is difficult to secure a desired material to be obtained, and when the carbon (C) content exceeds 0.3 wt%, the material is inferior due to an increase in the two-phase fraction and generation of pearlite (Understood strength) may occur. In addition, as the abnormal fraction increases, a problem in which hole expandability decreases may also occur. Therefore, in the present invention, the content of carbon (C) is preferably limited to 0.1 to 0.3% by weight.

Silicon (Si): 0.5 to 2.5 wt%

Silicon (Si) is well known as a ferrite stabilizing element and is well known as an element that increases ductility by increasing the ferrite fraction during cooling. In addition, since the carbide formation inhibitory power is very large, it is an essential element to secure the TRIP effect through an increase in the carbon concentration in the retained austenite during the formation of bainite. However, when a large amount of silicon (Si) is added, the concentration of carbon (C) increases toward austenite, which leads to good separation of abnormalities, but increases the hardness difference with the matrix structure ferrite, thereby lowering the hole expandability (HER) and yielding strength (YP) Also, the specification may be lowered. Therefore, it is preferable to add silicon (Si) in an amount of 0.5 to 2.5 wt%.

Manganese (Mn): 1.0 to 3.0 wt%

Manganese (Mn) as an austenite stabilizing element was used as an element to increase the fraction of the low-temperature phase and to increase the strength due to the solid solution strengthening effect. In order to secure strength and toughness, manganese (Mn) is preferably added in an amount of 1.0 wt% or more. However, when excessively added, segregation is easy to occur and martensite is formed in the segregation zone, which greatly inhibits the formation of a desired tissue. In particular, addition of manganese (Mn) is very effective in order to secure the tensile strength to 780 MPa or more, but it is preferable to limit the upper limit to 3.0 wt % or less in order to obtain high ductility at the same time.

Aluminum (Al): 0.1 to 1.0 wt%

Aluminum (Al) is an element that stabilizes ferrite and suppresses the formation of carbides, like silicon (Si), and improves ductility and formability, and is an important element in the present invention. In order to see the effectiveness of the addition of such aluminum (Al), 0.1% by weight or more should be added. However, when excessively added, high-temperature ductility is greatly reduced and productivity is lowered, so the upper limit is limited to 1.0 wt% or less.

Titanium (Ti): 0.005 to 0.1 wt%

Titanium (Ti) forms TRJCNFANF such as Ti (C, N) at high temperatures and exhibits the effect of refining crystal grains. In addition, since it has an effect of suppressing the formation of AlN, it is necessarily included in the steel of the present invention. The amount of addition must satisfy the following formula.

Ti(wt%)≥ 3.4 x N(wt%).

However, when titanium (Ti) is excessively added, the recrystallization temperature is excessively increased to cause a non-uniform structure, and thus the upper limit thereof is limited to 0.1% by weight.

If the content of titanium (Ti) is less than 0.005% by weight, it is difficult to achieve a precipitation effect sufficient to satisfy the material, and the possibility of insufficient strength increases. On the other hand, when the content of titanium (Ti) in the carbon range of the present invention exceeds 0.1% by weight, the above effect may reach a saturated state, rather coarse TiN may be generated, thereby deteriorating toughness. In the case of adding the titanium (Ti), it is possible to finely extract the abnormal tissue based on the precipitation effect, and improve the hole expandability by strengthening the ferrite matrix.

Niobium (Nb): 0.005 to 0.1 wt%

Precipitation strengthening elements such as niobium (Nb), molybdenum (Mo), and vanadium (V) are known to increase the phase degree of the soft phase and bring about grain refinement effects. However, in the case of the element, when excessively added, the recrystallization temperature is excessively increased to cause a non-uniform structure, so the upper limit thereof is limited to 0.1% by weight.

Other unavoidable additions

In the case of phosphorus (P), the addition range was limited to 0.1 wt% or less due to the problem of lowering corrosion resistance due to segregation of the center of the slab, and in the case of sulfur (S), which inhibits toughness and weldability, the content is more stringent to 0.02 wt% or less Nitrogen (N), which lowers the impact properties and elongation, and greatly reduces the toughness of the weld joint, is also strictly limited to 0.04% by weight or less.

The remaining component of the high-strength steel sheet of the present invention is iron (Fe). However, since unintended impurities from raw materials or the surrounding environment may inevitably be mixed in the normal steel manufacturing process, it cannot be excluded. Since these impurities are known to anyone skilled in the steel manufacturing process, all of them are not specifically mentioned in this specification.

The steel sheet of the present invention having the above alloy component content is an ultra-high strength steel sheet of tensile strength (TS): 780 MPa or more, and secures elongation of 25% or more and hole expandability of 35% or more to secure ductility and formability at the same time. The hole expandability (HER) is defined by the following relational expression showing the relationship with the _{hole diameter (D f} , mm) after breaking by expanding the _{initial hole diameter (D 0 , mm).}

Hole extensibility (HER) = (D _f -D ₀ )/D ₀ *100(%)

Microstructure and grain size of steel sheet

The steel sheet of the present invention has a microstructure composed of polygonal ferrite: 30 to 70%, bainite: 5 to 40%, martensite: 0 to 10%, and retained austenite: 5 to 15%. At this time, in the case of the retained austenite fraction, it is based on the result of analyzing the ¼ point in the thickness direction in the direction perpendicular to the rolling direction by EBSD.

When the area ratio of the polygonal ferrite phase is 30% or less, the target high ductility cannot be obtained, and when it exceeds 70%, the tensile strength is lowered. Therefore, it is preferable that the area fraction of the polygonal ferrite phase is 30 to 70%. Bainite has a smaller interphase hardness difference with ferrite than martensite, which effectively disperses the strain stress and helps improve formability. However, when the area ratio of bainite is 5% or less, the effect cannot be obtained, and when it is 40% or more, there is a problem in that ductility is lowered.

Martensite is a structure that can increase strength most effectively, but when the area ratio is 10% or more, ductility and formability are deteriorated. Retained austenite is a necessary phase to obtain high ductility, and it improves ductility by effectively absorbing energy during deformation due to strain-induced transformation. However, if it is less than 1%, this effect is difficult to obtain, and if it is more than 15%, voids are easily induced during the hole expansion test, thereby reducing the formability.

On the other hand, the average grain size of the high strength steel sheet of the present invention is 1 ~ 5㎛. The lower the average grain size, the more advantageous, and an average of 3 μm or less is most advantageous. When the average grain size is 5 μm or more, the effect of stress dispersing during the hole expansion test is reduced and the formability is deteriorated. The average grain size is based on the average value of data obtained by measuring 30 or more grains in the result of photographing ¼ point in the thickness direction at 500 times with an optical microscope.

In addition, the aspect ratio of the polygonal ferrite phase is 3 or less, the aspect ratio of bainite and martensite as the hard phase is 3 or more, and the phase having an aspect ratio of 3 or less among retained austenite has a structure of 70% or more. desirable. At this time, the aspect ratio is based on the average value of the data obtained from 30 or more crystal grains after photographing ¼ point in the thickness direction at 5000 times with a scanning electron microscope.

Next, another aspect of the present invention, a method for manufacturing a high-strength steel sheet having excellent ductility and formability will be described in detail.

Manufacturing method of high-strength steel sheet with excellent ductility and formability

1 is a process flowchart schematically illustrating a method of manufacturing a high-strength steel sheet having excellent ductility and formability according to an embodiment of the present invention.

Referring to FIG. 1 , the method for manufacturing a high-strength steel sheet having excellent ductility and formability of the present invention is, in weight %, carbon (C): 0.1 to 0.3%, silicon (Si): 0.5 to 2.5%, manganese (Mn) : 1.0 to 3.0%, phosphorus (P): 0.01% or less, sulfur (S): 0.02% or less, aluminum (Al): 0.1 to 1.0%, titanium (Ti): 0.005 to 0.1%, niobium (Nb): 0.005 -0.1%, preparing a hot-rolled steel sheet containing the remaining iron (Fe) and other unavoidable impurities (S110), cold rolling the hot-rolled steel sheet (S120); and annealing the cold-rolled cold-rolled steel sheet (S130); and cooling the steel sheet after the annealing heat treatment and post heat treatment (S140).

Preparing a hot-rolled steel sheet (S110)

Prepare a steel ingot or slab (hereinafter referred to as a slab) satisfying the above composition, and heat the slab to 1,150 to 1,300 ° C. to re-dissolve segregated components during casting and re-dissolve precipitates. The reheating temperature is preferably 1,150 ~ 1,300 ℃ to ensure a normal hot rolling temperature. If the reheating temperature is less than 1,150 ℃, there may be a problem that the hot rolling load increases rapidly, and if it exceeds 1,300 ℃, the amount of surface scale increases, which may lead to material loss.

After the reheating, hot rolling is performed in a conventional manner, but finish rolling is preferably performed in a temperature range of 800 to 1,100°C. When the finish rolling temperature exceeds 1,100° C., there is a risk that the quality of the steel sheet may be deteriorated due to the occurrence of scale on the surface of the steel sheet. In addition, at a rolling temperature of less than 800° C., the grains are refined to increase the strength, but it may cause an increase in the rolling load and a decrease in productivity.

Next, the hot-rolled sheet material is cooled to a predetermined coiling temperature. The cooling may be both air cooling or water cooling, and cooling may be performed at a cooling rate of 10 to 100° C./s. Preferably, the cooling is performed to a coiling temperature. In the present invention, the coiling temperature preferably satisfies 500 to 650 °C. The coiling temperature is to secure appropriate amounts of ferrite and pearlite, and when the coiling temperature is too high, coarse ferrite and pearlite are generated, making it difficult to secure strength. When the coiling temperature exceeds 650℃, the yield ratio decreases due to the formation of coarse grains, but there may be a problem that the toughness is lowered and the target strength is not reached. On the other hand, when the coiling temperature is lower than 500℃, the structure is As it becomes finer, strength and toughness can be increased, but it is difficult to satisfy elongation.

Cold rolling the hot-rolled steel sheet (S120)

After preparing the hot-rolled steel sheet, as a step of cold-rolling the hot-rolled steel sheet (S120), using the hot-rolled steel sheet prepared as above, pickling treatment and cold rolling are performed. The reduction ratio during cold rolling is preferably 55 to 70%, and the effect of increasing the hole expandability due to the effect of refining the structure is maximized. When rolling at a reduction ratio of 55% or less during cold rolling, it is difficult to obtain the target hole expandability due to the low effect of refining the structure. When rolling at a reduction ratio of 70% or more, the roll force increases and the process load increases.

Annealing step (S130)

In this step, the cold-rolled steel sheet is heated to a temperature of Ac1 ~ Ac3 at a temperature increase rate of 1 ~ 10 ℃ / s in a continuous annealing furnace having a normal slow cooling section. Preferably, the cold-rolled steel sheet is maintained at a temperature of 800 to 900° C. for 60 to 600 seconds. The annealing temperature is a very important factor, and it is possible to secure stable strength with a tensile strength of 780 MPa or more in a temperature range of 800 ~ 900 °C. If the annealing temperature is less than 800° C., the risk of non-recrystallized grains may increase, and it may be difficult to form sufficient austenite, thereby making it difficult to secure the strength targeted in the present invention. In addition, when the annealing temperature exceeds 900°C, the amount of bainite rapidly increases due to the formation of excessive austenite, which may cause excessive increase in yield strength and deterioration of ductility. After that, the growth of precipitates is suppressed.

The holding time is 60 to 600 seconds. If the temperature increase time is less than 60 seconds, the crystal grains become relatively fine and the strength increase occurs greatly, and if it exceeds 600 seconds, the temperature increase time is too long, causing difficulties in productivity.

Cooling and alloying heat treatment step (S140)

After annealing, cool to 400 ~ 550℃ at an average cooling rate of 3 ~ 20℃/s and hold for 10 ~ 100 seconds. The temperature after annealing affects the phase change of the austenite formed during annealing. When the cooling temperature after annealing is less than 400℃, it is difficult to achieve the target yield ratio due to martensite formation, and when it exceeds 550℃, it is difficult to obtain the target strength because pearlite, not bainite, is formed. After cooling, alloying heat treatment is performed. During hot-dip galvanizing, the plating bath temperature is maintained in the range of 450 to 550 °C, and the alloying heat treatment temperature is preferably 450 to 650 °C.

If the alloying heat treatment temperature is less than 450 ℃, there may be a problem that the tempering effect for carbon (C) diffusion cannot be sufficiently obtained, and when the alloying heat treatment temperature exceeds 550 ℃, the strength of the martensite phase is reduced by the generation of carbides problems may arise.

Hereinafter, the present invention will be described in more detail through examples. However, the following examples are only examples for explaining the present invention in more detail, and do not limit the scope of the present invention. Content not described here will be omitted because it can be technically inferred sufficiently by a person skilled in the art.

Example

After preparing steel slabs A to F having the alloy composition (unit: weight %) of Table 1 below, the process conditions of Table 2 were applied to perform predetermined hot rolling, winding, cold rolling, annealing, alloying heat treatment, etc. Specimens of 1 to 39 were prepared.

steel grade C Si Mn P S Al Ti Nb N remark A 0.18 1.0 1.5 <0.01 <0.002 0.5 0.025 0.04 <0.003 B 0.16 1.2 1.5 <0.01 <0.002 0.5 0.025 0.04 <0.003 C 0.14 1.1 2.0 <0.01 <0.002 0.7 0.02 0.03 <0.003 D 0.18 1.2 1.5 <0.01 <0.002 0.03 0.01 004 <0.003 Al-free E 0.2 1.5 1.5 <0.01 <0.002 0.3 0.02 - <0.003 No Nb F 0.16 1.0 1.5 <0.01 <0.002 0.3 0.02 0.06 <0.003

Psalter steel grade Thickness after rolling (mm) FDT(℃) Coiling temperature (℃) Annealing temperature (℃) Cold Rolled (%) One

A














3.2









900 550 800

50 2 550 830 3 520 830 4 550 860 5 480 830 6

B 550 800


50 7 550 830 8 580 860 9 620 860 10 680 830 11


A 550 800
60 12 550 830 13 520 840 14 550 860 15 560 840 70 16 620 820 17


B 550 800
60 18 550 830 19 580 830 20 550 860 21 560 840 70 22 680 810 23
C



930 550 760
50 24 550 800 25 520 830 26 550 860 27
C 550 740
60 28 550 830 29 580 860 30 550 860 31
F
3.4



900 460 830
60 32 520 830 33 620 830 34
D


3.2 550 760
50 35 550 800 36 520 830 37
E 550 740
60 38 550 830 39 580 860

Yield strength (YP), tensile strength (TS), elongation (EL), and hole expandability (HER) were measured for each of the specimens 1 to 39, respectively, and are shown in Table 3 below. In addition, the phase fraction and average grain size were measured for each of the specimens 1 to 39, respectively, and are shown in Table 4.

Psalter steel grade YP(MPa) TS(MPa) Elongation (%) Hole expandability (%) invention remark One

A 490 814 26.5 24.5 X
Insufficient hole expandability during 50% cold rolling 2 485 798 27.3 26.5 X 3 482 802 27.1 25.5 X 4 478 786 28.3 27 X 5 452 769 27.1 Coiling temperature 500℃ or less 6

B 465 806 26.4 24 X
Insufficient hole expandability during 50 % cold rolling
7 469 804 27.3 26.5 X 8 471 808 26.9 26 X 9 468 787 27.9 26.5 X 10 429 761 26.8 22.5 X Coiling temperature 500℃ or less 11

A 472 802 27.8 31 O 12 470 802 28.2 35.5 O 13 465 809 28.1 35 O 14 471 806 27.7 33 O 15 476 811 26.9 34 O 16 458 788 26.8 32 O 17


B 485 792 26.7 38.5 O 18 485 797 26.5 34 O 19 475 791 27.3 34.5 O 20 489 798 25.2 35 O 21 478 804 27.3 36 O 22 443 764 25.7 27 X Coiling temperature 650℃ or higher 23
C 427 754 25.5 22.5 X Below annealing temperature 24 494 815 25 23.5 X Insufficient hole expandability during 50 % cold rolling 25 491 806 24.9 25.5 X 26 483 800 26.5 25 X 27

C 436 774 23.6 25 X Below annealing temperature 28 486 819 26.3 30.5 O 29 473 816 26.7 30 O 30 476 810 27.8 31.5 O 31
F 461 776 22.8 26.5 X Coiling temperature 500℃ or less 32 487 802 26.3 32 O 33 491 816 27.1 31.5 O 34
D 502 762 22.4 19 X Lack of elongation and hole expandability when Al is not added 35 548 809 21.5 24 X 36 537 798 23.2 24 X 37
E 416 802 25.4 21 X Insufficient hole expandability when Nb is not added 38 424 792 29.4 24 X 39 420 787 30.4 23 X

Psalter steel grade Phase fraction (%) Average grain size (㎛) invention remark ferrite bainite marten
site Residual Austenite One A 62 21 10 7 7.5 X grain size coarse 2 65 16 9 8 8.2 X 3 67 14 10 9 6.9 X 4 65 19 8 8 6.3 X 5 68 12 10 10 6.3 X 6 B 68 15 8 9 7.2 X 7 67 17 6 10 6.8 X 8 65 19 8 8 6.4 X 9 66 21 6 11 6.4 X 10 69 15 8 8 6.9 X 11 A 63 22 7 8 2.4 O 12 62 21 8 9 2.6 O 13 63 23 5 9 2.4 O 14 65 22 5 8 2.3 O 15 61 19 10 10 2.2 O 16 65 17 10 8 2.7 O 17 B 65 21 9 10 2.6 O 18 64 19 7 10 2.4 O 19 61 23 6 11 2.5 O 20 66 19 5 10 2.4 O 21 68 20 4 8 2.1 O 22 72 12 8 8 2.9 X ferrite excess 23 C 76 14 5 5 9.5 X More than 70% ferrite 24 65 16 9 8 7.1 X grain size coarse 25 67 14 10 9 7.2 X 26 65 19 8 8 6.8 X 27 C 74 210 9 7 9.7 X More than 70% ferrite 28 64 16 9 8 3.2 O 29 67 14 10 9 2.8 O 30 65 19 8 8 2.6 O 31 F 72 12 8 8 2.1 X More than 70% ferrite 32 64 16 10 10 2.3 O 33 65 16 9 10 2.3 O 34 D 72 16 6 6 8.9 X More than 70% ferrite 35 65 15 14 6 6.9 X grain size coarse 36 63 15 14 8 7.1 X 37 E 73 11 9 7 12.4 X More than 70% ferrite 38 67 14 10 9 11.2 X grain size coarse 39 65 19 8 8 11.6 X

Referring to Tables 1 to 4, specimens 11 to 21, 28 to 30, 32, and 33 are specimens that satisfy all of the alloy components, winding temperature, annealing temperature and cold rolling reduction rate suggested in the present invention, and yield Strength (YP): 450 MPa or more, tensile strength (TS): 780 MPa or more, elongation (EL): 25% or more, and hole expansion (HER): 30% or more. It has been shown to have a texture and size. In particular, it can be expected to be replaced with high-strength steel for difficult-to-form parts among automobile parts by simultaneously exhibiting excellent elongation and hole expandability.

On the other hand, in the case of the remaining specimens, as shown in Tables 3 and 4, some alloy components, coiling temperature, and annealing temperature range presented in the present invention were out of range, and elongation and/or hole expandability did not meet the standards. , it can be confirmed that the microstructure and/or grain size are also out of the target range.

As described above, according to the present invention, high-strength steel for difficult-to-form parts among automobile parts by simultaneously securing elongation (≥ 25 %) and hole expansion (≥ 30 %) through control of alloy composition and process conditions. can be expected to be replaced by

Although the above description has been focused on the embodiments of the present invention, various changes or modifications may be made at the level of those skilled in the art. Such changes and modifications can be said to belong to the present invention without departing from the scope of the present invention. Accordingly, the scope of the present invention should be judged by the claims described below.

Claims (8)

By weight%, carbon (C): 0.1 to 0.3%, silicon (Si): 0.5 to 2.5%, manganese (Mn): 1.0 to 3.0%, phosphorus (P): 0.01% or less, sulfur (S): 0.02% Hereinafter, aluminum (Al): 0.1 to 1.0%, titanium (Ti): 0.005 to 0.1%, niobium (Nb): 0.005 to 0.1%, the remainder including iron (Fe) and other unavoidable impurities,
The titanium (Ti) and nitrogen (N) content is controlled by the following formula,
[Ti] ≥ 3.4 x [N]
([Ti] and [N] represent the content of titanium and nitrogen in weight percent in the total weight of the steel sheet, respectively),
Tensile strength (TS): 780 MPa or more, elongation (EL): 25% or more, and hole expansion (HER): characterized in that it has physical properties of 30% or more,
High-strength alloyed hot-dip galvanized steel sheet. According to claim 1,
The steel sheet is characterized in that it has a microstructure consisting of polygonal ferrite: 30 to 70%, bainite: 5 to 40%, martensite: 0 to 10%, and retained austenite: 5 to 15%,
High-strength alloyed hot-dip galvanized steel sheet. According to claim 1,
The average grain size of the steel sheet is 1 ~ 5 ㎛,
The aspect ratio of the polygonal ferrite phase is 3 or less,
The aspect ratio of bainite and martensite is 3 or more,
Characterized in that the phase having an aspect ratio of 3 or less among retained austenite has a structure of 70% or more,
High-strength alloyed hot-dip galvanized steel sheet. (a) by weight, carbon (C): 0.1 to 0.3%, silicon (Si): 0.5 to 2.5%, manganese (Mn): 1.0 to 3.0%, phosphorus (P): 0.01% or less, sulfur (S) : 0.02% or less, aluminum (Al): 0.1 to 1.0%, titanium (Ti): 0.005 to 0.1%, niobium (Nb): 0.005 to 0.1%, remaining iron (Fe) and other unavoidable impurities, including titanium ( preparing a hot-rolled steel sheet in which the content of Ti) ([Ti]) and the content of nitrogen (N) ([N]) satisfy [Ti] ≥ 3.4 x [N];
(b) winding the hot-rolled steel sheet in a temperature range of 500 to 650°C;
(c) cold rolling the steel sheet at a reduction ratio of 55 to 70%;
(d) annealing the cold-rolled cold-rolled steel sheet in a temperature range of 800 to 900°C; and
(e) characterized in that it comprises the step of heat-treating the steel sheet in a temperature range of 450 ~ 650 ℃ for alloying,
A method for manufacturing a high-strength alloyed hot-dip galvanized steel sheet. 5. The method of claim 4,
The step (a) is,
reheating the steel slab at 1,150 to 1,300 °C; and
Finishing rolling temperature (FDT): characterized in that it comprises the step of hot rolling at 800 ~ 1,100 ℃,
A method for manufacturing a high-strength alloyed hot-dip galvanized steel sheet. 6. The method of claim 5,
In the step (b), it characterized in that it further comprises the step of cooling the hot-rolled steel sheet at a cooling rate of 1 ~ 100 ℃ / s before the winding,
A method for manufacturing a high-strength alloyed hot-dip galvanized steel sheet. 6. The method of claim 5,
Step (e) is,
(e-1) maintaining the steel sheet for 60 to 600 seconds after heating the steel sheet to 800 to 900 °C at a temperature increase rate of 1 to 10 °C;
(e-2) maintaining the steel sheet for 10 to 100 seconds after cooling to 400 to 550 °C at a cooling rate of 3 to 20 °C/s; and
(e-3) characterized in that it comprises the step of heat-treating the steel sheet in a temperature range of 450 ~ 650 ℃,
A method for manufacturing a high-strength alloyed hot-dip galvanized steel sheet. 6. The method of claim 5,
The steel sheet after step (e) has physical properties of tensile strength (TS): 780 MPa or more, elongation (EL): 25% or more, and hole expansion (HER): 30% or more,
Polygonal ferrite: 30 to 70%, bainite: 5 to 40%, martensite: 0 to 10%, and retained austenite: characterized in that it has a microstructure consisting of 5 to 15%,
A method for manufacturing a high-strength alloyed hot-dip galvanized steel sheet.

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