KR102255823B1 - High-strength steel having excellent formability and high yield ratio and method for manufacturing same - Google Patents

Abstract

Provided are a steel sheet having a high yield ratio and improved formability and a manufacturing method thereof. The steel sheet having a high yield ratio and improved formability comprises, by wt%, C: 0.04 to 0.15%, Si: 0.5% or less, Mn: 0.5 to 2.0%, Ti: 0.2% or less (including 0%), Nb: 0.1% or less (including 0%), V: 0.2% or less (including 0%), Mo: 0.5% or less (excluding 0%), and the remainder of Fe and inevitable impurities, satisfies the following relational expression 1, and has a microstructure containing 95% or more of ferrite and the remainder of pearlite. The aspect ratio of the ferrite phase is 0.3 or more and 0.75 or less. In the microstructure, Ti, Nb, V, and Mo-based fine precipitates having an average size of 50 nm or less are distributed in an amount of 10^12 or more per square meter.

Description

High-yielding ratio steel sheet with excellent formability and its manufacturing method {HIGH-STRENGTH STEEL HAVING EXCELLENT FORMABILITY AND HIGH YIELD RATIO AND METHOD FOR MANUFACTURING SAME}

The present invention relates to a high-strength steel and a method of manufacturing the same, and more particularly, to a high-strength steel having high yield ratio and excellent formability, and a method of manufacturing the same.

In recent years, according to various environmental regulations and energy use regulations, the use of high-strength steel sheets is required to improve fuel efficiency or durability. In particular, as the impact stability regulation of automobiles has recently expanded, high-strength steel having excellent yield strength has been adopted for structural members such as members, seat rails, and pillars in order to improve the impact resistance of the vehicle body. Structural members have the advantage of absorbing impact energy as the yield strength compared to the tensile strength, that is, the yield ratio (yield strength/tensile strength) is higher. However, in general, as the strength of the steel sheet increases, the elongation decreases, resulting in a problem in that the forming processability decreases, so that the development of a material with improved high yield ratio and formability at the same time is required.

A typical manufacturing method to increase the yield strength is to use water cooling during continuous annealing. That is, a steel sheet having a tempered martensite structure in which the microstructure is tempered with martensite can be manufactured by cracking in the annealing process and then immersing in water for tempering. As a conventional technique, patent document 1 is mentioned. In Patent Document 1, a steel material having 0.18% or more of carbon is continuously annealed and then water-cooled to room temperature, followed by overaging treatment for 1 to 15 minutes at a temperature of 120 to 300°C, so that the martensite volume ratio is 80 to 97% or more. It is to develop site steel. In the case of manufacturing ultra-high strength steel by the tempering method after water cooling as described above, the yield ratio is very high, but there is a problem that the shape quality of the coil is deteriorated due to temperature variations in the width and length directions. Therefore, problems such as material defects and workability decrease depending on the part occur during roll forming processing.

The invention disclosed in Patent Document 2 is mentioned as a prior art for improving workability in the high-tensile steel sheet. Patent Document 2 is a steel sheet composed of a composite structure mainly composed of martensite, and proposes a method of manufacturing a high-tensile steel sheet in which fine precipitated copper particles having a particle diameter of 1 to 100 nm are dispersed inside the structure in order to improve workability. However, Patent Document 2 has a problem in that red heat embrittlement due to Cu may occur by excessively adding Cu content to 2 to 5% in order to precipitate good fine Cu particles, and the manufacturing cost may be excessively increased.

On the other hand, Patent Document 3 uses ferrite as a matrix structure, has a microstructure including 2 to 10 area% pearlite, mainly through the addition of carbon and nitride forming elements such as Nb, Ti, V, etc. A steel plate with improved strength by precipitation strengthening and grain refinement is proposed. Patent Document 3 has the advantage that high strength can be easily obtained compared to a low manufacturing cost, but the recrystallization temperature rises rapidly due to fine precipitates, so that sufficient recrystallization occurs and there is a disadvantage that high temperature annealing must be performed in order to secure ductility. . In addition, there is a problem in that it is difficult to obtain a high strength steel of 600 MPa or higher in the existing precipitation-reinforced steel that is strengthened by depositing carbon and nitride on a ferrite matrix.

Therefore, by solving the above-described problems, there is a demand for a steel material capable of exhibiting ultra-high strength while exhibiting high yield ratio and formability.

Japanese Patent Application Publication No. JP1992-289120 Japanese Patent Application Publication No. JP2005-264176 Korean Patent Application Publication No. KR2015-0073844

The present invention has been devised to solve the above problems, and an object of the present invention is to provide a high-strength steel having a yield ratio of 0.8 or more, an elongation of 10% or more and a strength of 980 MPa or more as an index for evaluating formability, and a method of manufacturing the same.

The subject of the present invention is not limited to the above description. Those of ordinary skill in the art to which the present invention pertains will not have any difficulty in understanding the additional subject of the present invention from the general details of the present specification.

One aspect of the present invention,

C: 0.04 to 0.15%, Si: 0.5% or less, Mn: 0.5 to 2.0%, Ti: 0.2% or less (including 0%), Nb: 0.1% or less (including 0%), V: 0.2% or less (Including 0%), Mo: 0.5% or less (excluding 0%), the balance contains Fe and inevitable impurities, and satisfies the following relational formula 1,

It has a microstructure containing 95% or more of ferrite and the balance pearlite,

The aspect ratio of the ferrite phase is 0.3 or more and 0.75 or less,

It relates to a high-yield-ratio steel sheet having excellent formability in which ^{10 12} or more Ti, Nb, V, Mo-based microprecipitates having an average size of 50 nm or less are distributed per square meter in the microstructure.

[Relationship 1]

21[Ti] + 11[Nb] + 20[V] + 10[Mo] ≤ 8

(Here, [Ti], [Nb], [V], and [Mo] mean the weight percent of the element.)

The steel sheet of the present invention may be a plated steel sheet having a zinc-based plating layer formed on one side thereof.

In addition, another aspect of the present invention,

C: 0.04 to 0.15%, Si: 0.5% or less, Mn: 0.5 to 2.0%, Ti: 0.2% or less (including 0%), Nb: 0.1% or less (including 0%), V: 0.2% or less (Including 0%), Mo: 0.5% or less (excluding 0%), a process of producing a steel slab containing the balance Fe and inevitable impurities and satisfying the above relational formula 1;

Manufacturing a hot-rolled steel sheet by hot rolling the steel slab so that the temperature at the outlet side of the finish rolling is Ar ₃ to Ar _{3 +50°C;}

A step of cooling the hot-rolled steel sheet to room temperature at an average cooling rate of 0.1°C/s or less after winding at 400 to 700°C; And

It relates to a method for manufacturing a high-yield-ratio steel sheet having excellent formability, including a step of cold rolling the cooled hot-rolled steel sheet at a reduction ratio of 40 to 70%, and then continuously annealing it at a temperature range of 700 to 790°C.

In the present invention, the process of first cooling the continuously annealed cold-rolled steel sheet to 620 to 700°C at a cooling rate of 1 to 10°C/sec; And a step of secondary cooling the first cooled cold-rolled steel sheet at a cooling rate of 5 to 20°C/sec to 300 to 580°C.

After the secondary cooling process, the step of overaging while maintaining a constant temperature may be further included.

The overaged cold-rolled steel sheet is hot-dip galvanized in a temperature range of 430 to 490°C, then alloying heat treatment is performed as necessary, and then cooled to a temperature of 100°C or less at an average cooling rate of 5°C/s or more. .

Temper rolling of less than 2% may be performed on the cooled cold-rolled steel sheet.

According to the present invention having the configuration as described above, first, a steel material having a high tensile strength of 980 MPa or more can be manufactured, and second, a high-strength steel sheet having a yield ratio of 0.8 or more and an elongation of 10% or more can be manufactured.

Such a steel plate can be used for parts requiring collision safety as structural members for automobiles.

1 is a photograph showing a microstructure according to Inventive Example 3 of the present invention.
2 is a photograph showing fine precipitates according to Inventive Example 3 of the present invention.

Hereinafter, the present invention will be described.

The present inventor confirmed through an experiment that the target tensile properties and microstructure can be realized when the components and the annealing operation conditions satisfy a specific relationship, and the present invention was completed.

Specifically, the high-strength steel according to an embodiment of the present invention, by weight% C: 0.04 to 0.15%, Si: 0.5% or less, Mn: 0.5 to 2.0%, Ti: 0.2% or less, Nb: 0.1% or less, V: 0.2% or less, Mo: 0.5% or less The balance Fe and inevitable impurities are included, and the above relational expression 1 is satisfied.

Hereinafter, the steel sheet composition of the present invention and the reason for limiting the content thereof will be described. Herein, "%" means "% by weight" unless otherwise stated.

C: 0.04 to 0.15%

Carbon (C) in steel is a very important element added for solid solution strengthening. In addition, carbon contributes to the strength by bonding with the precipitated element to form fine carbides. However, if the amount exceeds 0.15%, martensite is formed during cooling due to the above causes and an increase in hardenability, resulting in a sharp increase in strength, resulting in a decrease in elongation, and thus it is impossible to secure the elongation targeted by the present invention. In addition, due to poor weldability, welding defects occur when processing customer parts. On the other hand, if the carbon content is lower than 0.04%, it is very difficult to secure the desired strength, so it is preferable to limit the content to 0.04 to 0.15%. More preferably, it is limited to the range of 0.6 to 0.13%.

Si: 0.5% or less (excluding 0%)

Si is a ferrite stabilizing element and is an element advantageous in securing a ferrite fraction targeted by the present invention by promoting ferrite transformation during cooling. In addition, it is effective in increasing the strength of ferrite because of its good solid solution strengthening ability, and is a useful element capable of securing strength without deteriorating the ductility of the steel sheet. However, if it exceeds 0.5%, the effect of solid solution strengthening increases, resulting in a decrease in elongation, causing surface scale defects, resulting in inferior plating surface quality, and also limiting the amount of addition to 0.5% or less because it degrades chemical conversion treatment. desirable.

Mn: 0.5~2.0%

Manganese (Mn) is an element that prevents hot embrittlement due to the formation of FeS by completely depositing sulfur in the steel as MnS and solid-solution strengthening of steel. If the content is less than 0.5%, it is difficult to secure the target strength in the present invention. On the other hand, if it exceeds 2.0%, there is a high possibility that problems such as weldability and hot rollability may occur, and at the same time, the hardenability is increased to Since the site can be formed more easily, the elongation can be reduced. In addition, there is a problem that Mn-Band (band of Mn oxide) is formed in the structure, thereby increasing the risk of processing cracks and plate breakage, and there is a problem that Mn oxide is eluted on the surface during annealing, which greatly impairs plating properties. Therefore, in the present invention, it is desirable to limit the content of Mn to 0.5 to 2.0%, and more preferably, to limit to 0.8 to 1.8%.

Ti: 0.2% or less

Ti is a fine carbide forming element and contributes to securing yield strength and tensile strength. In addition, Ti, as a nitride forming element, has an effect of inhibiting AlN precipitation by depositing N in the steel as TiN, thereby reducing the risk of cracking during playing. However, when the Ti content exceeds 0.2%, coarse carbides are precipitated, the strength and elongation may be reduced by reducing the amount of carbon in the steel, and nozzle clogging may occur during playing. Therefore, in the present invention, it is preferable to limit the Ti content to 0.2% or less. More preferably, it is limited to 0.15% or less.

Nb: 0.1% or less

Nb is an element that segregates at the austenite grain boundary, suppresses coarsening of austenite grains during annealing heat treatment, and forms fine carbides, contributing to the increase in strength. However, when the Nb content exceeds 0.1%, coarse carbides are precipitated, strength and elongation may be reduced by reducing the amount of carbon in the steel, and manufacturing cost may increase. Therefore, in the present invention, it is preferable to limit the Nb content to 0.1% or less. More preferably, it is limited to 0.08% or less.

V: 0.2% or less

Vanadium (V) is an element that reacts with carbon or nitrogen to form carbon-nitrides, and is an element that plays an important role in increasing the yield strength of steel by forming fine precipitates at low temperatures. However, when the V content exceeds 0.2%, coarse carbides are precipitated, strength and elongation may be reduced by reducing the amount of carbon in the steel, and manufacturing cost may increase. Therefore, in the present invention, it is preferable to limit the Nb content to 0.1% or less. More preferably, it is limited to 0.08% or less.

Mo: 0.5% or less

Molybdenum (Mo) is an element that forms carbides, and plays a role of improving yield strength and tensile strength by maintaining fine size of precipitates when added in combination with carbon-nitride forming elements such as Ti, Nb, and V. In addition, molybdenum (Mo) is an element added to delay the transformation of austenite into pearlite and to improve the fineness and strength of ferrite. Such Mo has the advantage of improving the hardenability of the steel to finely form martensite at grain boundaries, thereby controlling the yield ratio. However, as an expensive element, the higher the content, the more disadvantageous in manufacturing it is, so it is preferable that the content is appropriately controlled. In order to obtain the above-described effect, it is preferable to add at most 0.5%.If the Mo content exceeds 0.5%, it causes a sharp increase in alloy cost, resulting in poor economic efficiency, and due to excessive grain refining effect and solid solution strengthening effect. Rather, there is a problem that the ductility of the steel is deteriorated. More preferably, it is limited to 0.3% or less.

·Relationship 1

In the present invention, it is necessary to contain Ti, Nb, V and Mo so as to satisfy the following relational formula 1.

[Relationship 1]

21[Ti] + 11[Nb] + 20[V] + 10[Mo] ≤ 8

(Here, [Ti], [Nb], [V], and [Mo] mean the weight percent of the element.)

The above relational expression 1 defines the relationship between the precipitated elements necessary to control fine precipitates in the steel. Ti, Nb, and V are widely known as elements that generate carbides, and when a large amount is added, the number of fine precipitates increases, contributing to an increase in yield strength and tensile strength, but there is a problem that the manufacturing cost of the steel sheet increases. When the precipitate is distributed in a large amount in a fine size, the precipitation strengthening effect increases, but since the elongation decreases, it is necessary to limit the content of precipitation elements in manufacturing a high-formed steel sheet with a high yield ratio. The above relation is an atomic fraction relation that considers the atomic weight of each element, the left side of the relational expression means the sum of the added precipitation elements, and the right side is the precipitation that can be added to the maximum to achieve the minimum elongation of 10% target in the present invention. It means the sum of the elements.

The remaining component of the present invention is iron (Fe). However, since unintended impurities from the raw material or the surrounding environment may inevitably be mixed in the normal manufacturing process, this cannot be excluded. Since these impurities are known to anyone of ordinary skill in the manufacturing process, all the contents are not specifically mentioned in the present specification.

Next, the high yield ratio steel sheet having excellent formability of the present invention needs to satisfy the following microstructure and phase fraction control conditions in addition to the alloy composition in order to improve workability such as ductility while satisfying the high yield ratio.

Specifically, the steel sheet of the present invention has a microstructure including at least 95% of ferrite and at least one of martensite, pearlite, and bainite. Ferrite is a soft structure that contributes to the ductility of the steel sheet, and when 980 MPa class strength is achieved with only ferrite and fine precipitates as in the present invention, it is possible to secure a target elongation of 10% when the fraction of ferrite is 95% or more. . In addition, in the final annealed steel sheet, the crystal shape of ferrite must have an elongated shape to secure the target strength.

In the present invention, it is preferable that the aspect ratio of the ferrite phase (Aspect ratio: a value obtained by dividing the short axis of crystal grains by the major axis) is 0.3 or more and 0.75 or less. In the present invention, the shape of ferrite is difficult to secure elongation due to relatively insufficient recrystallization of ferrite when the aspect ratio obtained by dividing the long axis of the ferrite grain by the short axis is less than 0.3, and when it exceeds 0.75, the recrystallization of ferrite is excessive and the target strength Is difficult to secure.

^{In addition, within the microstructure of the steel sheet of the present invention, 10 12} or more Ti, Nb, V, and Mo-based microprecipitates having an average size of 50 nm or less are distributed per square meter. The smaller the size of the precipitate and the higher the density, the larger the bar contributes to the strength. Therefore, when the number of fine precipitates is ^{less than 10 12} per square meter or the average size exceeds 50 nm, it is difficult to secure the target strength.

Next, a method of manufacturing a steel sheet having a high yield ratio excellent in formability according to the present invention will be described.

The method for manufacturing a steel sheet of the present invention includes a process of manufacturing a steel slab having the above-described steel composition component; Manufacturing a hot-rolled steel sheet by hot rolling the steel slab so that the temperature at the outlet side of the finish rolling is Ar ₃ to Ar _{3 +50°C;} A step of cooling the hot-rolled steel sheet after winding it at 400 to 700°C; And a step of cold-rolling the cooled hot-rolled steel sheet at a reduction ratio of 40 to 70%, followed by continuous annealing at a temperature range of 700 to 790°C.

First, in the present invention, after preparing a steel slab having the above-described steel composition, the hot-rolled steel sheet is manufactured by hot rolling it so that the _{temperature at the outlet side of the finish rolling is Ar 3} to Ar _{3 +50°C.}

In the present invention, it is preferable that hot rolling is performed so that the outlet temperature during hot rolling, more specifically, the outlet temperature of the finish rolling mill is Ar ₃ to Ar _{3 +50°C.} If the outlet side temperature is _{less than Ar 3} , the hot deformation resistance is likely to increase rapidly, and the top, tail and edges of the hot rolled coil become single-phase regions, increasing in-plane anisotropy and deteriorating formability. Can be. On the other hand, when Ar ₃ exceeds +50° C., not only too thick oxidized scale occurs, but there is a high possibility that the microstructure of the steel sheet becomes coarse.

Subsequently, in the present invention, the hot-rolled steel sheet is wound at 400 to 700° C. and then cooled to room temperature at an average cooling rate of 0.1° C./s or less.

In the present invention, after the hot rolling is completed, it is preferable to maintain a temperature range of 400 to 700° C. during winding. When the coiling temperature is less than 400°C, excessive martensite or bainite is generated, causing excessive strength increase of the hot-rolled steel sheet, and thus manufacturing problems such as shape defects due to load during cold rolling may occur. On the other hand, if it exceeds 700°C, the pickling property is deteriorated due to an increase in surface scale, so it is preferable to limit it to the above-described winding temperature. Subsequently, in the present invention, the wound hot-rolled steel sheet is cooled.

And in the present invention, after cold-rolling the cooled hot-rolled steel sheet at a reduction ratio of 40 to 70%, it is continuously annealed at a temperature range of 700 to 790°C.

The hot rolled steel sheet is subjected to a cold rolling process after pickling. During the cold rolling, it is preferable to roll at a reduction ratio of 40 to 70%. If the reduction ratio is less than 40%, the driving force for recrystallization is weakened, which is likely to cause problems in obtaining good recrystallized grains, and shape correction is very difficult. On the other hand, when the reduction ratio exceeds 70%, the possibility of cracking at the edge of the steel sheet is high, and the rolling load may increase rapidly.

In addition, the cold-rolled steel sheet undergoes a continuous annealing process to provide the basis for the microstructure desired in the present invention.

At this time, it is preferable to perform continuous annealing at a temperature range of 700°C or more and 790°C or less. If the annealing temperature is less than 700°C, it is difficult to secure the elongation because the ferrite is not sufficiently recrystallized. On the other hand, when the annealing temperature exceeds 790°C, recrystallization of ferrite is excessive, and ferrite grains and precipitates become coarse, so it may be difficult to secure a target strength.

Subsequently, the manufacturing method of the present invention comprises a step of first cooling the continuously annealed cold-rolled steel sheet to 620 to 700°C at a cooling rate of 1 to 10°C/sec; And a step of secondary cooling the first cooled cold-rolled steel sheet at a cooling rate of 5 to 20°C/sec to 300 to 580°C.

After continuous annealing, the steel sheet is first cooled to 650 to 700°C at a cooling rate of 1 to 10°C/s. This first cooling step has an aspect of suppressing the heat level of the plate shape due to a sudden temperature drop in the rapid cooling section by performing slow cooling prior to the subsequent second rapid cooling.

After the first cooling, the second cooling is performed at a cooling rate of 5 to 20°C/s to a temperature of 300°C or more and 580°C or less. At this time, the optimum plate shape can be secured by varying the rapid cooling rate and temperature according to the width and thickness of the steel plate to be manufactured. If the cooling temperature is less than 300℃, there is a possibility that a cooling deviation occurs in the width direction or the length direction of the steel plate and the plate shape is deteriorated.

The present invention may further include a process of overaging while maintaining a constant temperature after the secondary cooling process. By performing such an overaging treatment, it is possible to secure an optimum shape by performing constant heat treatment in the width direction and the length direction of the coil.

In the present invention, after the overaging treatment is completed, skin pass rolling may be performed in the range of 0.1 to 2.0% elongation. Typically, in the case of skin pass rolling, a yield strength increase of at least 50 MPa occurs without an increase in tensile strength. If the elongation is less than 0.1%, it is difficult to control the shape, and if the elongation is less than 2.0%, it is preferable to limit it to the above-described range because the operability is greatly unstable due to high drawing operation.

Further, in the present invention, after hot-dip galvanizing treatment of the overaged cold-rolled steel sheet in a temperature range of 430 to 490°C, then alloying heat treatment is performed as necessary, and an average cooling rate of 5°C/s or more up to a temperature of 100°C or less. It may further include a step of cooling with.

Hereinafter, the present invention will be described in detail through examples.

(Example)

Steel grade C Si Mn Ti Nb V Mo Relation 1 Comparative Steel 1 0.03 0.07 1.4 0.11 0.015 0.045 0.1 4.4 Invention Lesson 1 0.08 0 1.5 0.1 0.05 0.1 0.2 6.7 Invention Lesson 2 0.08 0 1.5 0.1 0.05 0 0.2 4.7 Invention Lesson 3 0.08 0 1.5 0.1 0 0.1 0.2 6.1 Invention Lesson 4 0.08 0 1.5 0.1 0 0 0.2 4.1 Invention Lesson 5 0.08 0 1.5 0.14 0 0 0.3 5.9 Comparative lecture 2 0.08 0 1.5 0.1 0 0.1 0.4 8.1

* In Table 1, the unit of content is% by weight, and the remaining components are Fe and unavoidable impurities.

As shown in Table 1, a steel slab having a composition component was vacuum-melted, heated in a heating furnace at a reheating temperature of 1200° C. for 1 hour, hot-rolled, and then wound. At this time, during the hot rolling operation, the hot rolling was terminated in the temperature range of FDT 880 to 920°C, and the coiling temperature (CT) was controlled at 650°C.

Thereafter, pickling was performed using a hot-rolled steel sheet, and cold rolling was performed with a cold rolling reduction ratio of 45%. The cold-rolled steel sheet was subjected to continuous annealing (SS) under the annealing conditions shown in Table 2, and the first slow cooling and the second rapid cooling were performed. At this time, the cooling end temperature was set to 650°C and the cooling rate was set to 3°C/s in the first slow cooling (SCS), and the cooling end temperature was set to 450°C in the second rapid cooling (RCS), and the cooling rate was set to 10°C. It was set to /s. Thereafter, the cold-rolled steel sheet, which was finally cooled secondarily, was subjected to skin pass rolling at a reduction ratio of 0.1%.

Thereafter, a JIS No. 5 tensile test piece of the cold-rolled steel sheet prepared as described above was prepared to measure the physical properties of the steel sheet, and the results are shown in Table 2 below. The specific measurement method is as follows. A specimen of JIS No. 5 size was taken in the direction perpendicular to the rolling direction of the annealed heat-treated steel sheet, and a tensile test was conducted at a strain rate of 0.01/sec.The test results were yield strength (YS), tensile strength (TS), and yield strength. The ratio (YR, YS/TS) and elongation (El) were shown.

And at this time, the fraction and aspect ratio of ferrite, and the density and average size of the fine precipitates were measured as the microstructure of each manufactured cold-rolled steel sheet, and are also shown in Table 2 below. The ferrite phase fraction and aspect ratio were measured through SEM, and the density and average size of the fine precipitates were measured using TEM. The specific measurement method is as follows. After nital etching on the TD surface (transverse direction) of the annealed heat-treated steel sheet, the aspect ratio was derived by measuring the long axis and the short axis diameter of the observed ferrite crystal grains through SEM at 3000 magnification as shown in Fig.2. . The ferrite phase can be distinguished from other tissues through SEM, but more precisely, it was reconfirmed that it occupies 100% fraction through EBSD and XRD phase analysis. For the microprecipitates, a thin foil was prepared for TEM observation and observed at a magnification of 30000 as shown in FIG.

Steel grade SS
(℃) YS
(MPa) TS
(MPa) YR El(%) ferrite Fine precipitate Remark Fraction (%) Aspect ratio Density (/m ² ) Average size
(nm) Comparative Steel 1 750 990 982 1.01 8 100 0.71 10 ¹³ 15 Comparative Example 1 Invention Lesson 1 750 1045 1072 0.97 10 100 0.65 10 ¹⁴ 12 Invention Example 1 Invention Lesson 2 750 966 1001 0.97 11 100 0.69 10 ¹³ 11 Invention Example 2 Invention Lesson 3 750 1061 1080 0.98 10 100 0.59 10 ¹⁴ 9 Invention Example 3 Invention Lesson 4 750 991 1016 0.98 10 100 0.57 10 ¹⁵ 16 Invention Example 4 Invention Lesson 5 750 1123 1118 1.00 10 100 0.45 10 ¹⁴ 10 Inventive Example 5 Comparative lecture 2 750 1100 1092 1.01 8 100 0.19 10 ¹⁴ 21 Comparative Example 2 Comparative Steel 1 800 867 869 1.00 9 100 0.78 10 ¹⁵ 31 Comparative Example 3 Invention Lesson 1 800 820 916 0.90 9 100 0.76 10 ¹⁴ 10 Comparative Example 4 Invention Lesson 2 800 824 882 0.93 11 100 0.81 10 ¹⁵ 9 Comparative Example 5 Invention Lesson 3 800 854 935 0.91 10 100 0.85 10 ¹⁵ 15 Comparative Example 6 Invention Lesson 4 800 832 870 0.96 9 100 0.77 10 ¹⁴ 12 Comparative Example 7 Invention Lesson 5 800 898 953 0.94 9 100 0.76 10 ¹³ 17 Comparative Example 8 Comparative lecture 2 800 844 978 0.86 9 100 0.79 10 ¹⁴ 9 Comparative Example 9 Comparative Steel 1 850 620 661 0.94 14 100 0.81 10 ¹⁵ 11 Comparative Example 10 Invention Lesson 1 850 561 684 0.82 16 100 0.82 10 ¹⁴ 17 Comparative Example 11 Invention Lesson 2 850 577 654 0.88 15 100 0.84 10 ¹⁵ 12 Comparative Example 12 Invention Lesson 3 850 583 699 0.83 13 100 0.81 10 ¹⁵ 5 Comparative Example 13 Invention Lesson 4 850 573 653 0.88 15 100 0.79 10 ¹⁴ 8 Comparative Example 14 Invention Lesson 5 850 620 711 0.87 15 100 0.77 10 ¹³ 9 Comparative Example 15 Comparative lecture 2 850 566 766 0.74 12 100 0.86 10 ¹⁴ 17 Comparative Example 16 Comparative Steel 1 900 512 545 0.94 24 100 0.87 10 ¹⁵ 12 Comparative Example 17 Invention Lesson 1 900 503 615 0.82 25 100 0.91 10 ¹⁴ 17 Comparative Example 18 Invention Lesson 2 900 548 600 0.91 24 100 0.84 10 ¹⁵ 16 Comparative Example 19 Invention Lesson 3 900 536 624 0.86 24 100 0.88 10 ¹⁴ 19 Comparative Example 20 Invention Lesson 4 900 554 600 0.92 25 100 0.89 10 ¹⁴ 21 Comparative Example 21 Invention Lesson 5 900 555 633 0.88 24 100 0.91 10 ¹⁴ 22 Comparative Example 22 Comparative lecture 2 900 437 670 0.65 23 100 0.84 10 ¹⁵ 17 Comparative Example 23

As shown in Table 1-2, in the case of Inventive Examples 1 to 5 satisfying the composition range and annealing conditions of the present invention, it can be seen that all of the physical properties required by the present invention are satisfied.

1 is a photograph showing a microstructure according to Inventive Example 3 of the present invention, and FIG. 2 is a photograph showing a fine precipitate according to Inventive Example 3 of the present invention.

On the other hand, Comparative Steel 1 is a case where the C composition range of the present invention is not satisfied, and Comparative Steel 2 is a case where the numerical value of the composition relational formula 1 of the present invention is not satisfied. As shown in Table 2, when using these comparative steels, even if the annealing conditions described in the present invention are satisfied, target tensile properties as in Comparative Examples 1, 2, 3, 9, 11, 16, 17 and 23 It can be seen that it is impossible to secure.

In addition, in the case of Comparative Examples 4 to 8, 11 to 15, and 18 to 22, which do not satisfy the annealing condition even if the composition range of the present invention and the relational expression 1 are satisfied, it is difficult to secure the target physical properties. For example, in the case of this example, the annealing temperature exceeded 790°C and the recrystallization of ferrite was excessive. As a result, the aspect ratio exceeded 0.75, and thus the tensile strength was not satisfied with 980 MPa.

The present invention is not limited to the above embodiments and embodiments, but may be manufactured in a variety of different forms, and those of ordinary skill in the art to which the present invention pertains other It will be appreciated that it can be implemented in a specific form. Therefore, it should be understood that the implementation examples and embodiments described above are illustrative in all respects and are not limiting.

Claims (7)

C: 0.04 to 0.15%, Si: 0.5% or less, Mn: 0.5 to 2.0%, Ti: 0.2% or less (including 0%), Nb: 0.1% or less (including 0%), V: 0.2% or less (Including 0%), Mo: 0.5% or less (excluding 0%), the balance contains Fe and inevitable impurities, and satisfies the following relational formula 1,
It has a microstructure containing 95 area% or more of ferrite and the balance pearlite,
The aspect ratio of the ferrite phase is 0.3 or more and 0.75 or less,
A high-yield-ratio steel sheet having excellent formability in which ^{10 12} or more Ti, Nb, V, Mo-based microprecipitates having an average size of 50 nm or less are distributed in the microstructure.
[Relationship 1]
21[Ti] + 11[Nb] + 20[V] + 10[Mo] ≤ 8
(Here, [Ti], [Nb], [V], and [Mo] mean the weight percent of the corresponding element.)
The high yield ratio steel sheet according to claim 1, wherein the steel sheet is a plated steel sheet having a zinc-based plating layer formed on one surface thereof.
C: 0.04 to 0.15%, Si: 0.5% or less, Mn: 0.5 to 2.0%, Ti: 0.2% or less (including 0%), Nb: 0.1% or less (including 0%), V: 0.2% or less (Including 0%), Mo: 0.5% or less (excluding 0%), the balance of Fe and the process of producing a steel slab that satisfies the following relational formula 1 including inevitable impurities;
Manufacturing a hot-rolled steel sheet by hot rolling the steel slab so that the temperature at the outlet side of the finish rolling is Ar ₃ to Ar _{3 +50°C;}
A step of cooling the hot-rolled steel sheet to room temperature at an average cooling rate of 0.1°C/s or less after winding at 400 to 700°C; And
After cold-rolling the cooled hot-rolled steel sheet at a reduction ratio of 40 to 70%, a process of continuously annealing it at a temperature range of 700 to 790°C.
[Relationship 1]
21[Ti] + 11[Nb] + 20[V] + 10[Mo] ≤ 8
(Here, [Ti], [Nb], [V], and [Mo] mean the weight percent of the element.)
The method of claim 3, further comprising: first cooling the continuously annealed cold-rolled steel sheet to 620 to 700°C at a cooling rate of 1 to 10°C/second; And a step of secondary cooling the first cooled cold-rolled steel sheet to 300 to 580 °C at a cooling rate of 5 to 20 °C/sec.
The method of claim 4, further comprising the step of overaging while maintaining a constant temperature after the secondary cooling process.
The method of claim 5, wherein the overaged cold-rolled steel sheet is hot-dip galvanized in a temperature range of 430 to 490°C, then alloying heat treatment is performed as needed, and an average cooling of 5°C/s or more to a temperature of 100°C or less. High yield ratio steel sheet manufacturing method having excellent formability, characterized in that cooling at a speed.
The method of claim 5, wherein temper rolling of less than 2% is performed on the overaged cold-rolled steel sheet.

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