KR20230087773A - Steel sheet having excellent strength and ductility, and manufacturing method thereof - Google Patents

Abstract

The present invention relates to a steel sheet having excellent strength suitable for automobile materials and having excellent ductility at the same time, and a manufacturing method thereof.

Description

Steel sheet with excellent strength and ductility and its manufacturing method {STEEL SHEET HAVING EXCELLENT STRENGTH AND DUCTILITY, AND MANUFACTURING METHOD THEREOF}

The present invention relates to a steel sheet having excellent strength suitable for automobile materials and having excellent ductility at the same time, and a manufacturing method thereof.

Recently, as interest in reducing carbon dioxide and corresponding regulatory standards have increased, automobile manufacturers are making great efforts to improve fuel efficiency through weight reduction of vehicle bodies.

In addition, as the conversion rate from the existing engine room-centered internal combustion engine to an electric vehicle free from exhaust gas emissions gradually increases, the trend is shifting to the development of automotive steel sheets that can be applied to structural members centered on electric vehicle component parts.

In particular, as regulations on weight reduction and passenger safety of electric vehicles are strengthened, the application rate of high-strength steel for structural members of battery cases is continuously increasing.

In order to reduce the weight of the car body, the thickness of the steel plate must be reduced within the range of securing passenger safety, and to realize this, automobile manufacturers have applied a ratio of high-strength steel with excellent yield strength to tensile strength, that is, the yield ratio (YS/TS). are expanding

In general, an increase in strength is accompanied by a decrease in formability due to a decrease in ductility. In the case of structural member parts, most molding processes are performed, so in order to stably secure the target shape during molding, not only crash stability but also sufficient ductility are required. must be secured

Accordingly, it is necessary to develop a steel sheet for automobiles having excellent strength and ductility at the same time in accordance with these technical requirements.

On the other hand, Patent Document 1 discloses a technique for manufacturing a steel sheet having a martensite phase with an area fraction of 80% or more by subjecting a steel material having a carbon (C) content of 0.18% or more to continuous annealing, water cooling, and overaging treatment. In this case, the yield ratio is increased due to the tempering effect of martensite, but the shape quality of the coil deteriorates due to the temperature deviation in the width and length directions of the steel plate, or cracks occur during molding due to material deviation problems, resulting in poor workability. There is a risk of becoming

Patent Document 2 discloses a high-strength steel sheet having a tensile strength of 1400 MPa or more and a yield ratio of 0.7 or more and excellent crash characteristics. However, this technology has a high amount of alloys such as carbon and manganese, so there is a risk of deterioration in spot weldability and playability of the steel sheet, and poor formability due to low ductility.

In addition, Patent Document 3 discloses a technique for manufacturing a high-strength steel sheet having a yield strength of 1300 MPa or more and a tensile strength of 1500 MPa or more. This steel sheet has the advantage of excellent strength and excellent crash resistance for automobiles, but has a low elongation of less than 8%, so there is a limit in that it can be used only for the production of simple shaped parts such as roll forming.

Accordingly, it is required to develop a steel sheet capable of securing strength and ductility at the same time while solving the above-mentioned problems of the existing technologies.

Japanese Unexamined Patent Publication No. 1992-289120 Korean Patent Publication No. 2020-0027387 Korean Patent Publication No. 2014-0097332

One aspect of the present invention is to provide a steel sheet excellent in strength as well as ductility and a method for manufacturing the same, as a steel sheet suitable for automobile materials, particularly materials such as members requiring high formability.

The object of the present invention is not limited to the above. A person skilled in the art will have no difficulty understanding the further subject matter of the present invention from the general content of this specification.

One aspect of the present invention, by weight, carbon (C): 0.10 ~ 0.22%, manganese (Mn): 2.3 ~ 3.0%, silicon (Si): 0.3 ~ 0.8%, aluminum for acid value (Sol.Al): 0.1% or less (excluding 0%), Phosphorus (P): 0.05% or less (excluding 0%), Sulfur (S): 0.01% or less (excluding 0%), Nitrogen (N): 0.01% or less (excluding 0%) and chromium (Cr): 0.8% or less, niobium (Nb): 0.1% or less, titanium (Ti): 0.1% or less, molybdenum (Mo): 0.4% or less, and boron (B): 0.003% or less. As above, including the balance iron (Fe) and unavoidable impurities, and satisfying the following relational expression 1,

Provided is a steel sheet having excellent strength and ductility, including martensite and tempered martensite with an area fraction of 90% or more (excluding 100%) and the balance of ferrite, bainite, and retained austenite as a microstructure.

[Relationship 1]

4.5 ≤ C + 1.5Mn + 2Cr + B ≤ 6.0

(Here, each element means a weight content.)

Another aspect of the present invention comprises the steps of reheating a steel slab satisfying the above-described alloy composition and relational expression 1 in a temperature range of 1000 to 1350 ° C; manufacturing a hot-rolled steel sheet by hot-rolling the heated steel slab in a temperature range of Ar3 to Ar3+50°C; winding the hot-rolled steel sheet in a temperature range of 450 to 700° C.; manufacturing a cold-rolled steel sheet by cold-rolling the wound hot-rolled steel sheet; Continuously annealing the prepared cold-rolled steel sheet in a temperature range of 760 to 850 ° C.; firstly cooling the continuously annealed cold-rolled steel sheet to a temperature range of 550 to 700° C. at an average cooling rate of 1 to 10° C./s; Secondary cooling at an average cooling rate of 2 to 20 ° C / s to a temperature range of 400 to 600 ° C after the first cooling; and overaging for 100 to 300 seconds after the secondary cooling.

According to the present invention, a steel sheet having high strength and excellent ductility can be provided. In particular, the steel sheet of the present invention has excellent formability and has an effect that it can be suitably applied to automotive structural parts requiring processing into complex shapes.

1 shows a microstructure photograph of an inventive steel according to an embodiment of the present invention.
2 and 3 show pictures of microstructures of comparative steels according to an embodiment of the present invention.

The inventors of the present invention have studied in depth to provide a steel sheet with improved ductility so as to have formability that can be processed into parts requiring complex shapes while having strength suitable for automobile materials.

As a result, it was confirmed that by optimizing the alloy composition system and manufacturing conditions of the steel, it was possible to have a structure advantageous to securing the target physical properties, and from this, it was possible to provide a steel plate suitable for automotive structural members requiring processing into complex shapes. and came to complete the present invention.

Hereinafter, the present invention will be described in detail.

The steel sheet having excellent strength and ductility according to one aspect of the present invention contains, by weight, carbon (C): 0.10 to 0.22%, manganese (Mn): 2.3 to 3.0%, silicon (Si): 0.3 to 0.8%, aluminum ( Sol.Al): 0.1% or less (excluding 0%), Phosphorus (P): 0.05% or less (excluding 0%), Sulfur (S): 0.01% or less (excluding 0%), Nitrogen (N): 0.01% or less (except 0%).

Hereinafter, the reason for limiting the alloy composition of the steel sheet provided in the present invention as above will be described in detail.

Meanwhile, in the present invention, unless otherwise specified, the content of each element is based on weight, and the ratio of tissue is based on area.

Carbon (C): 0.10 to 0.22%

Carbon (C) is an element that is advantageous for securing strength by forming martensite in steel, and is an essential element for manufacturing high-strength steel. In general, the higher the content of C, the easier it is to form martensite, which is advantageous for the formation of a complex structure required for manufacturing high-strength steel. However, in order to simultaneously control the intended strength and elongation, it is necessary to control the content at an appropriate level.

In the present invention, the C may be added in an amount of 0.10% or more in order to secure a target strength and form an appropriate level of martensite. However, when the content is excessive and exceeds 0.22%, there is a problem in that weldability and formability are inferior.

Therefore, in the present invention, the C may be included in 0.10 to 0.22%, more advantageously, 0.12% or more and 0.20% or less.

Manganese (Mn): 2.3 to 3.0%

Manganese (Mn) is an element that improves hardenability of steel, and plays an important role in forming martensite. In addition, the Mn contributes to the increase in strength of steel by the solid solution strengthening effect, and plays a role in suppressing plate breakage and high-temperature brittleness caused by S during hot rolling by precipitating sulfur (S), which is inevitably added in steel, as MnS. do.

In order to sufficiently obtain the above-mentioned effect, it is preferable to add the Mn in an amount of 2.3% or more. However, if the content exceeds 3.0%, not only does weldability deteriorate, martensite is excessively formed, making the material unstable, and a band-shaped oxide band is formed, which increases the risk of processing cracks and plate breakage. There is a growing problem. In addition, during annealing, manganese oxide is eluted from the surface of the steel sheet and there is a possibility of impairing plating properties.

Therefore, in the present invention, the Mn may be included in 2.3 to 3.0%, more advantageously, 2.5% or more and 2.8% or less.

Silicon (Si): 0.3 to 0.8%

Silicon (Si) is a useful element capable of securing strength without reducing the ductility of the steel sheet. Such Si promotes ferrite formation and promotes C enrichment into untransformed austenite to promote martensite formation.

In order to sufficiently obtain the above-mentioned effect, it is advantageous to add Si at 0.3% or more, but when the content exceeds 0.8%, there is a risk of causing hydrogen embrittlement and poor weldability as well as plating properties.

Therefore, in the present invention, the Si may be included in an amount of 0.3 to 0.8%, and more advantageously, it may be included in an amount of 0.4% or more and 0.7% or less.

Acid soluble aluminum (Sol.Al): 0.1% or less (excluding 0%)

Acid-soluble aluminum (Sol.Al) is an element added for deoxidation and particle size refinement of steel. If the content of Sol.Al is excessive and exceeds 0.1%, not only the castability deteriorates during continuous casting, but also excessive inclusions. There is a problem in that the possibility of generating defects in the material of the annealed material and surface defects of the plating material increases due to formation.

Therefore, the Sol.Al may be included in an amount of 0.1% or less, but 0% may be excluded in consideration of the level inevitably added to the steel. More advantageously, the Sol.Al may be included in an amount of 0.01% or more.

Phosphorus (P): 0.05% or less (excluding 0%)

Phosphorus (P) is the most advantageous element for securing the strength of steel without significantly impairing formability, but when added excessively, the possibility of brittle fracture greatly increases, increasing the possibility of plate breakage of slabs during hot rolling. is also In addition, there is a problem that P also acts as an element that inhibits plating surface characteristics.

In consideration of this, the P may be included in an amount of 0.05% or less, and 0% may be excluded in consideration of a level inevitably added to steel.

Sulfur (S): 0.01% or less (excluding 0%)

Sulfur (S) is an impurity element that is inevitably added to steel, and it is desirable to manage its content as low as possible. In particular, since S in steel can increase the possibility of glowing brittleness, it is preferable to limit its content to 0.01% or less. However, 0% can be excluded considering the level that is unavoidably added during the manufacturing process.

Nitrogen (N): 0.01% or less (excluding 0%)

Nitrogen (N) is an impurity element that is unavoidably added to steel, and it is important to manage its content as low as possible. However, since there is a problem that the refining cost of steel rises rapidly for this purpose, it is preferable to control the operating conditions to 0.01% or less, which is a possible range. However, 0% can be excluded considering the level that is unavoidably added during the manufacturing process.

On the other hand, in order to advantageously secure the intended physical properties of the steel sheet of the present invention having the above-described alloy composition, chromium (Cr): 0.8% or less, niobium (Nb): 0.1% or less, titanium (Ti): 0.1% or less, molybdenum (Mo): 0.4% or less and boron (B): may further include one or more selected from 0.003% or less.

Chromium (Cr): 0.8% or less

Chromium (Cr) is an element added to secure high strength by improving the hardenability of steel. Cr is not only effective in forming martensite, but also advantageous in manufacturing high-strength steel having high ductility by minimizing the decrease in elongation compared to the increase in strength.

However, when the Cr content exceeds a certain level, there is a problem of excessively increasing the formation rate of martensite and increasing the fraction of coarse Cr-based carbides, resulting in a decrease in elongation. In addition, there is a possibility of deteriorating hydrogen embrittlement and weldability.

Therefore, the addition of Cr may be limited to 0.8% or less, and more advantageously, 0.6% or less.

Niobium (Nb): 0.1% or less

Niobium (Nb) is an element that is segregated at austenite grain boundaries to suppress coarsening of austenite crystal grains during annealing heat treatment and contributes to improvement of yield strength and tensile strength by forming fine carbides.

However, when the Nb content is excessive, strength and elongation may decrease due to precipitation of coarse carbides and reduction in carbon content in steel, and manufacturing costs may increase, resulting in poor economic feasibility.

Therefore, when adding the Nb, it may be limited to 0.1% or less, more advantageously, it may be included to 0.05% or less.

Titanium (Ti): 0.1% or less

Titanium (Ti) not only contributes to securing yield strength and tensile strength by forming fine carbides, but also suppresses the precipitation of AlN by precipitating N in steel as TiN, thereby effectively reducing the risk of cracking during play.

However, if the content is excessive, the strength and elongation may be reduced due to the precipitation of coarse carbides and the reduction of the carbon content in the steel, and nozzle clogging may also be caused during playing.

Therefore, when the Ti is added, it may be limited to 0.1% or less, and more advantageously, it may be 0.05% or less.

Molybdenum (Mo): 0.4% or less

Molybdenum (Mo) is an element that improves hardenability of steel and contributes to strength improvement by forming fine carbides in steel by combining with Ti, Nb, etc. as well as refinement of ferrite.

The Mo is an expensive element, and if its content is excessive, manufacturing costs may increase significantly, and crystal grain refinement and solid solution strengthening effects may be excessively increased, which may rather decrease ductility.

Therefore, when adding the Mo, it may be limited to 0.4% or less, and more advantageously, it may be 0.2% or less.

Boron (B): 0.003% or less

Boron (B) is effective in delaying the transformation of austenite into pearlite in a cooling process after continuous annealing.

However, if the content of B is excessive, B may be concentrated on the surface of the steel sheet, resulting in deterioration of coating adhesion.

Therefore, the addition of B may be limited to 0.003% or less.

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

In the steel sheet of the present invention that satisfies the above-described alloy composition, the relationship between the contents of C, Mn, Cr, and B preferably satisfies the following relational expression 1.

[Relationship 1]

4.5 ≤ C + 1.5Mn + 2Cr + B ≤ 6.0

(Here, each element means a weight content.)

Equation 1 above is obtained as an empirical value for the content relationship of specific components to secure the basic material of the steel sheet desired in the present invention.

If the value of relational expression 1 is less than 4.5, since the hardenability of the steel is lowered, the phase fraction of the low-temperature transformation structure such as martensite or bainite is lowered, making it impossible to secure the desired strength in the present invention. On the other hand, if the value exceeds 6.0, the hardenability of the steel is excessively high, and thus a material having desired formability cannot be obtained in the present invention.

More advantageously, the value of relational expression 1 may be 4.7 or more and 5.8 or less.

The steel sheet of the present invention having the above-described alloy component system includes a martensite phase and a tempered martensite phase as a main phase, and may include ferrite, bainite, and retained austenite as other structures.

In the present invention, the martensite phase generated during the steel manufacturing process is mostly transformed into the tempered martensite phase during the subsequent heat treatment (corresponding to the post-heat treatment (tempering) process of the present invention). At this time, the rate of transformation into the tempered martensite phase will be determined according to the post-heat treatment conditions of the present invention, and it should be noted that almost all martensite phases can be transformed into the tempered martensite phase under the conditions of the present invention.

More specifically, the steel sheet of the present invention preferably contains a martensite phase and a tempered martensite phase in an area fraction of 90% or more. If the fraction is less than 90%, the strength of the target level cannot be secured. More advantageously, 95% or more of the martensite phase and the tempered martensite phase may be included.

In addition, when the fraction of the remaining structure excluding the martensite phase exceeds 10%, not only ductility of the target level but also hole expandability cannot be secured. In this case, preferably, it is advantageous to include 5% or less of ferrite, 4% or less of bainite, and 1% or less of retained austenite.

The steel sheet of the present invention may contain fine precipitates in the above-described microstructure, and specifically, at least one type of fine precipitates selected from the group consisting of Ti, Nb, and Mo having an average diameter (equivalent circle diameter) of 50 nm or less per unit area (m ² ) per 10 ¹² or more may be included.

In the present invention, by including the fine precipitates as described above, there is an effect of securing the desired high-strength characteristics more advantageously.

More specifically, the steel sheet of the present invention has a yield strength of 1000 MPa or more, a tensile strength of 1470 MPa or more, a yield ratio of 0.7 or more, and an elongation of 10% or more, which can ensure excellent strength and elongation at the same time.

In addition, the steel sheet of the present invention is characterized in that the product of yield strength and elongation (YS×El) is 11000 MPa·% or more, and the product of tensile strength and elongation (TS×El) is 15000 MPa·% or more. In addition, the present invention can secure hole expandability (HER) of 30% or more by effectively lowering the hardness difference between phases by martensitic tempering.

In general, during phase transformation of ferrite, bainite, etc., since carbon concentration in austenite remaining in the vicinity increases, strength of steel increases due to martensite with increased solid-solution carbon in the final structure. At this time, when the TRIP effect is accompanied by the stabilization of retained austenite, the elongation can be improved at the same time. In addition, an increase in elongation can be expected even when the fraction of soft phase ferrite increases.

However, in the case of the present invention, not only the ferrite fraction, but also the high elongation even though the fraction of retained austenite is relatively low, along with the homogeneous structure of martensite (and tempered martensite) formed with an area fraction of 90% or more This is because the hardness difference between the phases is greatly reduced. In other words, due to the high fraction of martensite (and tempered martensite), it is possible to secure high elongation due to tissue homogeneity and reduction of hardness difference between phases, along with ultra-high tensile strength of 1470 MPa or more.

Meanwhile, the steel sheet of the present invention may be a cold-rolled steel sheet, and may be a hot-dip galvanized steel sheet including a zinc-based plating layer on at least one surface of the cold-rolled steel sheet, and an alloyed hot-dip galvanized steel sheet obtained by alloying the hot-dip galvanized steel sheet.

Although not particularly limited, the zinc-based plating layer may be a zinc plating layer mainly containing zinc or a zinc alloy plating layer containing aluminum and/or magnesium in addition to zinc.

Hereinafter, a method for manufacturing a steel sheet having excellent strength and ductility provided by the present invention, which is another aspect of the present invention, will be described in detail.

Briefly, the present invention can manufacture a desired steel sheet through the process of [steel slab reheating - hot rolling - winding - cold rolling - continuous annealing - cooling - overaging], and then [dip galvanizing - alloying heat treatment], etc. Further steps can be performed.

Conditions for each step are described in detail below.

[Reheat steel slab]

A steel slab having predetermined components is prepared. Since the steel slab of the present invention has a composition corresponding to the alloy component system of the above-described steel sheet, the description of the alloy component system of the steel slab is replaced with the above description.

The prepared steel slab may be heated to a certain temperature range, and the heating temperature of the steel slab at this time may be in the range of 1000 to 1350 ° C. If the heating temperature of the steel slab is less than 1000 ° C, there is a risk of hot rolling in the temperature range below the target finishing hot rolling temperature range, and if the heating temperature of the steel slab exceeds 1350 ° C, the steel reaches the melting point and melts. because it has potential.

[Hot rolling]

The steel slab heated according to the above may be hot-rolled to produce a hot-rolled steel sheet. During the hot rolling, the finishing hot rolling may be performed in a temperature range of Ar3˜Ar3+50° C.

If the temperature during the finish hot rolling is less than Ar3, there is a high possibility that the hot deformation resistance will rapidly increase. On the other hand, if the temperature exceeds Ar3 + 50 ° C, an excessively thick oxide scale will occur, and the crystal grains of the hot-rolled steel sheet will be coarsely formed. This may cause deterioration of physical properties of the final steel sheet.

Here, the finishing hot rolling temperature may be the temperature at the exit side of the finishing mill.

[wind up]

The hot-rolled steel sheet manufactured according to the above may be cooled and wound in a temperature range of 450 to 700 ° C.

If the coiling temperature is less than 450 ° C., martensite or bainite is excessively generated, and manufacturing problems such as shape defects due to load during subsequent cold rolling may occur. On the other hand, when the temperature exceeds 700 ° C., there is a problem in that pickling property deteriorates due to an increase in surface scale.

[Cold Rolling]

The coiled hot-rolled steel sheet may be cold-rolled to produce a cold-rolled steel sheet. At this time, the cold rolling conditions (eg, reduction ratio) are not particularly limited, but may be performed at a reduction ratio of 35 to 65%.

If the reduction ratio during the cold rolling is less than 35%, the recrystallization driving force is weakened, making it difficult to secure good recrystallized grains, and shape correction may be difficult. On the other hand, if the reduction ratio exceeds 65%, there is a possibility that cracks may occur at the edge of the steel sheet, and the rolling load may increase rapidly.

Meanwhile, prior to performing the cold rolling, a pickling process for removing scale generated on the surface of the hot-rolled steel sheet may be further performed. The pickling process can be performed under normal conditions, and the conditions are not particularly limited.

[Continuous Annealing]

The cold-rolled steel sheet manufactured according to the above may be continuously annealed, and through the continuous annealing treatment, a foundation for the microstructure desired in the present invention may be provided.

Specifically, the continuous annealing may be performed at a temperature range of 760 to 850 ° C, and it is preferable to maintain the temperature for 30 to 300 seconds.

When the temperature during the continuous annealing is less than 760° C., the fraction of ferrite is increased so that not only the strength is lowered, but also the hole expandability may be deteriorated due to the increase in the hardness difference between the phases. On the other hand, when the temperature exceeds 850 ° C., the fraction of ferrite is greatly reduced and ductility is lowered, and there is a possibility that an annealed oxide is formed on the surface of the steel sheet. More preferably, the continuous annealing may be performed at a temperature of 770° C. or higher and 820° C. or lower.

When maintaining the cold-rolled steel sheet in the continuous annealing temperature range, if the time is less than 30 seconds, recrystallization of ferrite is not sufficiently performed, making it difficult to secure elongation, and the possibility of material deviation in the length / width direction of the steel sheet increases. On the other hand, when the time exceeds 300 seconds, the annealing effect is saturated, and there is a problem in that productivity is lowered.

[Cooling]

The cold-rolled steel sheet subjected to the continuous annealing process may be cooled, and at this time, it is preferable to perform cooling in stages according to the temperature range.

Specifically, the cooling is performed first at an average cooling rate of 1 to 10°C/s to a temperature range of 550 to 700°C, and then to an average cooling rate of 2 to 20°C/s to a temperature range of 400 to 600°C. Secondary cooling is possible.

1st cooling

The present invention is carried out by slow cooling compared to secondary cooling, which is a subsequent process during primary cooling, and can suppress plate-shaped inferiority due to temperature drop during secondary cooling, which is a relatively rapid cooling section.

When the cooling end temperature during the primary cooling is less than 550° C. or exceeds 700° C., it is out of an appropriate temperature gradient range with subsequent secondary cooling, making it difficult to secure stable cooling capacity.

In addition, when the average cooling rate during the primary cooling exceeds 10° C./s, there is a concern that sufficient concentration of C and Mn in austenite may not occur. In the present invention, the lower limit of the average cooling rate during the primary cooling is not particularly limited, but since cooling is performed at the cooling rate of normal slow cooling, it may be 1 ° C. / s or more.

secondary cooling

During the secondary cooling of the primary cooled cold-rolled steel sheet, an optimal plate shape can be secured by appropriately adjusting the cooling rate and cooling end temperature according to the width and thickness of the steel sheet to be obtained.

When the cooling end temperature during the secondary cooling is less than 400° C., the total martensite fraction is relatively reduced as the fraction of bainite to be transformed increases, making it impossible to secure the target level of strength. On the other hand, when the temperature exceeds 600 ° C., it is not possible to obtain a bainite phase, and some pearlite is introduced into the structure, so that the target level of strength cannot be secured. More advantageously, the cooling end temperature of the secondary cooling may be 430°C or more and 570°C or less.

In addition, when the average cooling rate during the secondary cooling exceeds 20° C./s, there is a concern that sufficient concentration of C and Mn in austenite may not occur. The average cooling rate during the secondary cooling may vary depending on the temperature deviation from the primary cooling end temperature, and may be preferably 2° C./s or more. More advantageously, the average cooling rate may be 4° C./s or more and 15° C./s or less.

[Over-age]

The cold-rolled steel sheet subjected to secondary cooling may be subjected to an overaging treatment, and at this time, the overaging treatment may be performed within a temperature range at which the secondary cooling is terminated.

In the present invention, the overaging treatment may be a process of maintaining the temperature for 100 to 300 seconds. At this time, if the holding time is excessive and exceeds 300 seconds, excessive bainite transformation occurs during the holding process, and there is a concern that the fraction of the bainite phase in the final structure may increase. Since this leads to a decrease in the fraction of martensite, it is impossible to effectively secure the desired strength. On the other hand, if the time is less than 100 seconds, the bainite phase may not be contained in the final structure as the bainite nose is avoided, which causes the ductility of the steel to decrease.

[Hot-dip galvanizing, alloyed hot-dip galvanizing]

The cold-rolled steel sheet subjected to the overaging treatment according to the above may be immersed in a plating bath to perform plating. At this time, although not limited thereto, a hot-dip zinc-based plating bath may be used, and after such hot-dip galvanizing, alloying heat treatment may be selectively performed.

In the present invention, the plating method for hot-dip galvanizing or hot-dip galvanizing alloy is not particularly limited, and process conditions for hot-dip galvanizing or alloyed hot-dip galvanizing may be applied. As a non-limiting example, the hot-dip galvanizing may be performed in a zinc-based plating bath at 430 to 490° C., and the alloying heat treatment may be performed at a temperature range of 480 to 600° C.

In addition, the composition of the plating bath in the hot-dip galvanizing or alloyed hot-dip galvanizing is not particularly limited, but may be a pure zinc plating bath or a zinc-based alloy plating bath containing Si, Al, Mg, and the like.

The steel sheet completed with hot-dip galvanization and, if necessary, alloying heat treatment as described above may be cooled to room temperature under normal conditions, and the cooling process is not particularly limited. However, it will be obvious that it can be replaced with known cooling methods such as water cooling, oil cooling, and furnace cooling.

[Temper rolling]

Temper rolling may be further performed on the plated steel sheet according to the above, and at this time, it may be performed at a reduction ratio of 0.1 to 1.0%.

It is possible to obtain the effect of increasing the yield strength without accompanying the increase in tensile strength and rainfall, which is usually performed on the steel sheet by temper rolling. However, if the rolling reduction of the temper rolling is less than 0.1%, the effect of increasing the yield strength is insignificant, and it is difficult to control the shape. On the other hand, when the reduction ratio exceeds 1.0%, there is a concern that operability may be significantly inferior due to high elongation work.

[Post heat treatment (tempering)]

In the present invention, a post-heat treatment process may be further performed after the temper rolling treatment. In this case, the post-heat treatment process is a tempering process, and tempered martensite may be introduced into the steel sheet.

Temperature and holding time during the post-heat treatment are major process factors for controlling the ratio of tempered martensite to total martensite. By introducing tempered martemsite through the post-heat treatment, it is possible to effectively secure desired yield strength and hole expandability at the same time while maintaining ductility.

Specifically, the post-heat treatment process may be a process of heating the temper-rolled steel sheet to a temperature range of 100 to 400° C., maintaining the temperature at the temperature for 20 to 50 seconds, and then cooling the steel sheet to room temperature.

If the temperature during the post-heat treatment is low or the holding time is not sufficient, the ratio of the tempered martensite fraction to the total martensite fraction does not reach a desired level. On the other hand, when the temperature is too high or the holding time is long, the ratio of the tempered martensite fraction to the total martensite fraction exceeds the desired level. In this case, a decrease in strength due to coarsening of the structure and a decrease in ductility due to precipitation of carbides such as cementite may be induced.

Therefore, in the present invention, the post-heat treatment process may be performed at 100 to 400 ° C for 20 to 50 seconds, more advantageously at 150 ° C or higher and 300 ° C or lower. More advantageously, it may be 220 ° C or less.

The steel sheet of the present invention manufactured through the above-described series of processes is characterized by a yield ratio of 0.7 or more and an elongation of 10.0% or more, along with ultra-high strength of tensile strength of 1470 MPa or more and yield strength of 1000 MPa or more, by forming the intended microstructure. In addition, the steel sheet has a product of yield strength and elongation of 11000 MPa %, a product of tensile strength and elongation of 15000 MPa % or more, and a hole expandability (HER) of 30% or more, having ultra-high strength and excellent ductility at the same time. can be provided

Hereinafter, the high-strength steel sheet excellent in formability and the manufacturing method of the present invention will be described in detail through specific examples. It should be noted that the following examples are only for understanding of the present invention, and are not intended to specify the scope of the present invention. The scope of the present invention is determined by the matters described in the claims and the matters reasonably inferred therefrom.

(Example)

A steel slab having the alloy composition shown in Table 1 (the rest being Fe and unavoidable impurities) was vacuum melted, heated in a temperature range of 1200 ° C, and then finish rolling was completed at an outlet temperature of 880 to 920 ° C, 600 Coiled at °C. Thereafter, after removing surface scale by pickling, cold rolling was performed at a cold rolling reduction ratio of 50% to prepare a cold rolled steel sheet. Then, the cold-rolled steel sheet was subjected to continuous annealing treatment, stepwise cooling, and overaging treatment under the conditions shown in Table 2 below.

Hot-dip galvanizing was performed on the cold-rolled steel sheet that had been subjected to the overaging treatment, and temper rolling was performed at a reduction ratio of 0.3%, and then post-heat treatment (tempering) was completed under the conditions shown in Table 2 below to prepare a final steel sheet. At this time, hot-dip galvanizing was performed in a galvanizing bath at 430 to 490 ° C.

steel grade Alloy composition (% by weight) relational expression 1 C Si Mn Cr Mo Nb Ti B Al P S N A 0.16 0.6 2.6 0.5 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 5.1 B 0.16 0.6 2.8 0.5 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 5.4 C 0.24 0.6 2.6 0.5 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 5.1 D 0.16 0.6 2.1 0.5 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 4.3 E 0.16 0.6 3.3 0.5 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 6.1 F 0.16 0.1 2.6 0.5 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 5.1 G 0.16 1.0 2.6 0.5 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 5.1 H 0.16 0.6 2.6 0 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 4.1 I 0.16 0.6 2.6 1.0 0.1 0.03 0.02 0.002 0.03 0.01 0.002 0.003 6.1 J 0.16 0.6 2.6 0.5 0.5 0.03 0.02 0.002 0.03 0.01 0.002 0.003 5.1

steel grade continuous annealing 1st cooling secondary cooling overaging treatment post heat treatment division temperature
(℃) maintain
hour
(s) end
temperature
(℃) Cooling
speed
(℃/s) end
temperature
(℃) Cooling
speed
(℃/s) temperature
(℃) maintain
hour
(s) temperature
(℃) maintain
hour
(s) A 790 42 650 3.1 490 9.5 460 150 60 30 Comparative Example 1 A 790 42 650 3.1 490 9.5 460 150 170 30 Invention example 1 A 790 42 650 3.1 490 9.5 460 150 450 30 Comparative Example 2 A 780 42 650 2.9 520 7.7 460 180 200 30 Invention example 2 A 870 42 700 3.8 500 11.9 460 150 200 30 Comparative Example 3 A 750 42 650 2.2 500 8.9 460 150 200 30 Comparative Example 4 A 790 42 530 5.8 500 1.8 460 150 200 30 Comparative Example 5 A 790 42 650 3.1 370 16.6 460 150 200 30 Comparative Example 6 A 790 42 650 3.1 500 8.9 460 70 200 30 Comparative Example 7 A 790 42 650 3.1 500 8.9 460 330 200 30 Comparative Example 8 B 780 42 650 2.9 530 7.1 460 200 200 30 Inventive example 3 C 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 9 D 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 10 E 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 11 F 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 12 G 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 13 H 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 14 I 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 15 J 790 42 650 3.1 500 8.9 460 150 200 30 Comparative Example 16

The microstructure of the steel sheet thus prepared was observed and shown in Table 3 below. Among the microstructures, martensite (M), tempered martensite (T-M), ferrite (F), and bainite (B) were observed through SEM after nital etching of the polished specimen cross section, and retained austenite (R-γ) was measured through XRD analysis.

Precipitates in the microstructure were observed using TEM. At this time, after observing the image at 30000 magnification, the number and average diameter of the fine precipitates per unit area were measured.

In addition, the physical property values for each specimen were measured, and the results are described in Table 3 below. Yield strength (YS), tensile strength (TS), and elongation (El) were evaluated through tensile tests, and each mechanical Physical properties were measured.

In addition, the hole expansion rate (HER) was evaluated through a hole expansion test, and after forming a punched hole of 10 mmØ (die inner diameter 10.3 mm, clearance 12.5%), a conical punch with a vertex angle of 60 ° was used to remove the burr of the punched hole. It was inserted into the punched hole in an outward direction, and after pressing and expanding the punched hole periphery at a moving speed of 12 mm/min, it was calculated using the following [Equation].

[ceremony]

Hole expansion rate (HER, %) = {(D - D ₀ ) / D ₀ } × 100

(Here, D means the hole diameter (mm) when the crack penetrates the steel sheet along the thickness direction, and D ₀ means the initial hole diameter (mm).)

division Microstructure (area fraction%) fine precipitate mechanical properties M
(+TM) F B R-γ density
(/m ² ) average
diameter
(μm) YS
(MPa) TS
(MPa) El
(%) YR YS×El
(MPa %) TS×El
(MPa %) HER
(%) Comparative Example 1 93 2 4 One 10 ¹³ 14 1060 1530 10.1 0.69 10706 15453 28 Invention example 1 93 2 4 One 10 ¹³ 13 1127 1512 10.3 0.75 11608 15574 34 Comparative Example 2 93 2 4 One 10 ¹³ 14 1205 1460 7.4 0.83 8917 10804 30 Invention example 2 92 4 3 One 10 ¹³ 15 1101 1507 10.4 0.73 11450 15673 32 Comparative Example 3 94 0 5 One 10 ¹⁵ 17 1244 1526 9.2 0.82 11445 14039 37 Comparative Example 4 87 9 3 One 10 ¹² 12 962 1450 8.0 0.66 7696 11600 23 Comparative Example 5 93 2 4 One 10 ¹³ 14 1065 1501 9.8 0.71 10437 14710 30 Comparative Example 6 85 2 11 2 10 ¹³ 13 1025 1488 8.7 0.69 8918 12946 32 Comparative Example 7 97 2 0 One 10 ¹³ 13 1011 1503 9.5 0.67 9605 14279 34 Comparative Example 8 83 2 13 2 10 ¹³ 15 987 1467 9.0 0.67 8883 13203 26 Inventive example 3 93 3 3 One 10 ¹⁴ 15 1140 1521 10.1 0.75 11514 15362 31 Comparative Example 9 93 2 3 2 10 ¹⁵ 19 1241 1529 8.4 0.81 10424 12844 25 Comparative Example 10 94 2 3 One 10 ¹³ 13 997 1485 9.6 0.67 9571 14256 32 Comparative Example 11 93 2 3 2 10 ¹⁴ 15 1152 1531 8.0 0.75 9216 12248 22 Comparative Example 12 79 2 17 2 10 ¹⁵ 16 971 1473 9.2 0.66 8933 13552 35 Comparative Example 13 94 2 3 One 10 ¹³ 13 1130 1519 9.8 0.74 11074 14886 32 Comparative Example 14 94 2 3 One 10 ¹³ 13 988 1452 10.1 0.68 9979 14665 28 Comparative Example 15 94 2 3 One 10 ¹⁴ 14 1078 1538 7.8 0.70 8408 11996 25 Comparative Example 16 96 2 One One 10 ¹⁵ 12 1034 1551 7.3 0.67 7548 11322 23

(In Table 3, fine precipitates mean one or more types of fine precipitates selected from the group consisting of Ti, Nb, and Mo.)

As shown in Tables 1 to 3, Inventive Examples 1 to 3 satisfying both the alloy composition system and manufacturing conditions proposed in the present invention are martensite phase (+ tempered martensite phase) formed at 90 area% or more as intended In addition, ferrite, bainite, and retained austenite phases were appropriately formed as other structures. In addition, as fine precipitates were formed as intended, not only ultra-high strength of 1470 MPa or more, but also elongation of 10% or more and hole expandability of 30% or more were secured.

That is, it can be seen that the strength and ductility of the steel sheet manufactured according to the present invention are greatly improved at the same time.

On the other hand, in Comparative Examples 1 to 7, while the alloy components satisfy the proposal of the present invention, the manufacturing conditions deviate from the present invention.

Among them, Comparative Example 4 is a case where the temperature during the continuous annealing treatment is considerably low, and despite a relatively high ferrite fraction, it can be confirmed that the ductility is poor. This can be attributed to the fact that, due to the high fraction of ferrite, the hardness difference with the martensite phase having a high carbon concentration greatly increases, and thus strain localization occurs intensively at the interface.

In comparison, as mentioned above, it can be seen that Inventive Examples 1 to 3 ensured high elongation even though the ferrite fraction was low. This is because the microstructure is mostly composed of martensite (+ tempered martensite), which not only increases the homogeneity of the structure, but also shows ultra-high strength by reducing the difference in hardness between the ferrite and martensite phases, and at the same time has an elongation of 10% or more. that can represent

On the other hand, in Comparative Example 1, as the post-heat treatment was performed at too low a temperature, the tempering of martensite was not sufficient, and the hole expandability was relatively poor due to the limitation of the effect of reducing the hardness difference between phases.

Comparative Example 2 is a case in which the temperature during the post-heat treatment was excessively high, and as the martensite tempering occurred excessively, it was not possible to secure the tensile strength at the level of wooden cloth, and the desired ductility could not be secured due to the formation of coarse carbides.

In Comparative Example 3, since the temperature during the continuous annealing treatment was excessively high, annealing was initially performed in a single phase region, so that ferrite was not introduced, and thus, ductility was inferior.

In Comparative Example 5, as the primary cooling was terminated at too low a temperature during cooling after the continuous annealing treatment, a bainite phase was additionally formed, so that the target level of ductility could not be secured.

In Comparative Example 6, as the secondary cooling was terminated at too low a temperature after the continuous annealing treatment, a large amount of bainite phase was formed, so that the target level of ductility could not be secured.

In Comparative Example 7, the overaging treatment time was not sufficient, and as a result, martensite phase was excessively formed and bainite phase was not introduced at all, resulting in high strength but poor ductility.

In Comparative Example 8, as the overaging treatment was performed for an excessively long time, the bainite phase was excessively formed, while the martensite phase was insufficient, so that desired strength, ductility, and hole expandability could not be secured.

Comparative Examples 9 to 16 are cases where the manufacturing conditions satisfy the present invention, but the alloy component system deviate from the proposal of the present invention.

Among them, Comparative Example 9 was inferior in ductility and hole expandability as the martensite phase formed by the excessive addition of C in the steel became relatively harder and the hardness difference with the surrounding phases increased.

Comparative Example 10 was a case where the content of Mn in the steel was not sufficient, and strength and ductility were inferior. In particular, in Comparative Example 10, the value of relational expression 1 was obtained as 4.3, so it can be understood that the strength intended in the present invention was not secured in Comparative Example 10 as the hardenability of the steel decreased.

In Comparative Example 11, it can be seen that the Mn content in the steel is excessive, and in particular, the value of relational expression 1 exceeds 6.0. That is, in Comparative Example 10, the ductility and formability (hole expandability) targeted in the present invention were not secured due to excessively high hardenability of the steel.

In Comparative Example 12, when the content of Si in the steel was not sufficient, the martensite phase was not sufficiently formed and the bainite phase was excessively formed, so that strength and ductility could not be excellently secured at the same time.

In Comparative Example 13, the content of Si in the steel was excessive, and the solid solution strengthening effect was greatly increased, so that strength was secured, but ductility was inferior accordingly.

Comparative Example 14 is a case in which Cr is not added, and the target level of strength and hole expandability cannot be secured. In particular, in Comparative Example 14, the value of relational expression 1 was obtained as low as 4.1, and thus, in Comparative Example 14, the strength intended in the present invention was not secured because the hardenability of the steel was lowered.

In Comparative Example 15, it can be seen that the Cr content is excessive, and in particular, the value of relational expression 1 exceeds 6.0. That is, in Comparative Example 15, since the hardenability of the steel was excessively high, the ductility and formability (hole expandability) targeted in the present invention could not be secured.

Comparative Example 16 is a case in which the content of Mo is excessive, and as the hardenability is greatly increased and the bainite nose is pushed further back, the bainite phase in the final structure is relatively lowered, and the hardness difference between the phases increases. It was not possible to secure the level of ductility and hole expandability.

1 shows a microstructure photograph (SEM, 5000 times) of Inventive Example 1, and it can be seen that ferrite, bainite, and retained austenite phases are properly formed in addition to the martensite phase.

Meanwhile, FIGS. 2 and 3 show microstructural photographs (SEM, 5000 times) of Comparative Example 5 and Comparative Example 11, respectively. As shown in FIGS. 2 and 3 , it can be confirmed that the bainite phase is excessively formed.

Claims (12)

In % by weight, carbon (C): 0.10 to 0.22%, manganese (Mn): 2.3 to 3.0%, silicon (Si): 0.3 to 0.8%, acid soluble aluminum (Sol.Al): 0.1% or less (excluding 0%) ), phosphorus (P): 0.05% or less (excluding 0%), sulfur (S): 0.01% or less (excluding 0%), nitrogen (N): 0.01% or less (excluding 0%), and chromium (Cr): 0.8% or less, niobium (Nb): 0.1% or less, titanium (Ti): 0.1% or less, molybdenum (Mo): 0.4% or less, and boron (B): 0.003% or less, balance iron (Fe) and unavoidable impurities, and satisfies the following relational expression 1,
A steel sheet with excellent strength and ductility including martensite and tempered martensite with an area fraction of 90% or more (excluding 100%) and the balance ferrite, bainite, and retained austenite in microstructure.

[Relationship 1]
4.5 ≤ C + 1.5Mn + 2Cr + B ≤ 6.0
(Here, each element means a weight content.)
According to claim 1,
The steel sheet includes one or more types of fine precipitates selected from the group consisting of Ti-based, Nb-based, and Mo-based,
The fine precipitates have an average diameter of 50 nm or less, and a unit area (m ² ) per unit area (m 2 ) 10 ¹² steel sheet having excellent strength and ductility.
According to claim 1,
The steel sheet has a yield strength of 1000 MPa or more, a tensile strength of 1470 MPa or more, a yield ratio of 0.7 or more, and an elongation of 10% or more.
According to claim 1,
The steel sheet has excellent strength and ductility, wherein the product of yield strength and elongation (YS × El) is 11000 MPa % or more, and the product of tensile strength and elongation (TS × El) is 15000 MPa % or more.
According to claim 1,
The steel sheet has excellent hole expandability (HER) of 30% or more in strength and ductility.
In % by weight, carbon (C): 0.10 to 0.22%, manganese (Mn): 2.3 to 3.0%, silicon (Si): 0.3 to 0.8%, acid soluble aluminum (Sol.Al): 0.1% or less (excluding 0%) ), phosphorus (P): 0.05% or less (excluding 0%), sulfur (S): 0.01% or less (excluding 0%), nitrogen (N): 0.01% or less (excluding 0%), and chromium (Cr): 0.8% or less, niobium (Nb): 0.1% or less, titanium (Ti): 0.1% or less, molybdenum (Mo): 0.4% or less, and boron (B): 0.003% or less, balance iron (Fe) and reheating a steel slab containing unavoidable impurities and satisfying the following relational expression 1 in a temperature range of 1000 to 1350° C.;
manufacturing a hot-rolled steel sheet by hot-rolling the heated steel slab in a temperature range of Ar3 to Ar3+50°C;
winding the hot-rolled steel sheet in a temperature range of 450 to 700° C.;
manufacturing a cold-rolled steel sheet by cold-rolling the wound hot-rolled steel sheet;
Continuously annealing the prepared cold-rolled steel sheet in a temperature range of 760 to 850° C.;
firstly cooling the continuously annealed cold-rolled steel sheet to a temperature range of 550 to 700° C. at an average cooling rate of 1 to 10° C./s;
Secondary cooling at an average cooling rate of 2 to 20 ° C / s to a temperature range of 400 to 600 ° C after the first cooling; and
Step of overaging treatment for 100 to 300 seconds after the secondary cooling
Method for producing a steel sheet having excellent strength and ductility comprising a.

[Relationship 1]
4.5 ≤ C + 1.5Mn + 2Cr + B ≤ 6.0
(Here, each element means a weight content.)
According to claim 6,
The cold rolling is a method for producing a steel sheet having excellent strength and ductility, which is performed at a cold rolling reduction of 35 to 65%.
According to claim 6,
The continuous annealing process is a method for producing a steel sheet having excellent strength and ductility to be performed for 30 to 300 seconds.
According to claim 6,
Method for producing a steel sheet having excellent strength and ductility, further comprising the step of hot-dip galvanizing the cold-rolled steel sheet subjected to the overaging treatment.
According to claim 9,
A method for producing a steel sheet having excellent strength and ductility, further comprising performing an alloying heat treatment after the hot-dip galvanizing.
According to claim 10,
A method for producing a steel sheet having excellent strength and ductility, further comprising the step of temper rolling at a reduction ratio of 0.1 to 1.0% after the hot-dip galvanizing treatment or alloying heat treatment.
According to claim 11,
Method for producing a steel sheet with excellent strength and ductility further comprising a post-heat treatment (tempering) step of heating to a temperature range of 100 to 400 ° C. after the temper rolling treatment, holding for 20 to 50 seconds, and then cooling to room temperature.

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