CN117043382A - High-strength steel sheet excellent in hole expansibility and ductility and method for producing same - Google Patents

Abstract

The present invention relates to steel suitable as an automobile material, and specifically to a high-strength steel plate with excellent hole expandability and ductility and a manufacturing method thereof. The microstructure of the high-strength steel plate of the present invention is composed of a hard phase and a soft phase. Through optimized cold rolling and annealing processes, the martensite phase as the hard phase is evenly distributed on the recrystallized ferrite matrix, and The crack resistance during processing can be improved by introducing a non-equilibrium ferrite phase at the interface of the hard phase and the soft phase.

Description

High-strength steel sheet with excellent hole expandability and ductility and method for producing the same

Technical Field

The present invention relates to a steel suitable for automobile materials, and more particularly to a high-strength steel sheet having excellent hole expandability and ductility and a method for manufacturing the same.

Background Art

In recent years, in the automotive industry, due to environmental regulations related to _CO2 emissions and energy use regulations, the use of high-strength steel is required to improve fuel efficiency and durability.

In particular, as regulations on impact stability of automobiles are expanded, high-strength steel having excellent strength is used as a material for structural parts such as members, seat rails, and pillars in order to improve the impact resistance of vehicle bodies.

Such automotive parts have complex shapes in terms of stability and design, and are mainly formed and manufactured using press dies, and therefore require high strength and a high level of formability.

However, as the strength of steel increases, it has the characteristic of absorbing impact energy, but generally, as the strength increases, the elongation decreases, so there is a problem of reduced formability. In addition, when the yield strength is too high, the introduction of material into the mold during molding will decrease, so there is a problem of poor formability and increased manufacturing costs.

In addition, there are many parts of automobile parts that expand after processing holes, so hole expandability (HER) is required for smooth molding, but high-strength steel has low hole expandability, and there are problems such as cracks during the molding process. As mentioned above, if the hole expandability is poor, cracks will be generated in the molded part of the part when the car collides, making the part easy to be damaged, so the safety of passengers may be threatened.

In addition, representative high-strength steels used as automotive materials include dual phase steel (DP steel), transformation induced plasticity steel (TRIP steel), complex phase steel (CP steel), ferrite bainite steel (FB steel), etc.

DP steel, as an ultra-high tensile steel, has a low yield strength ratio of about 0.5 to 0.6, so it is easy to process and has the advantage of high elongation second only to TRIP steel. Therefore, it is mainly used in outer doors, seat rails, seat belts, suspensions, arms, wheel discs, etc.

TRIP steel has a yield ratio in the range of 0.57 to 0.67, and thus has the characteristic of exhibiting excellent formability (high ductility), and is therefore suitable for parts requiring high formability, such as frames, roofs, seat belts, and bumpers.

CP steel has a low yield strength ratio, high elongation and bending workability, so it is used in side panels, underbody reinforcements, etc., and FB steel has excellent hole expandability, so it is mainly used in suspension lower arms or wheel discs, etc.

Among them, DP steel is mainly composed of ferrite with excellent ductility and hard phases (martensite phase and bainite phase) with high strength, and there may be a trace of retained austenite. This DP steel has low yield strength and high tensile strength, thus having excellent characteristics such as low yield ratio (YR), high work hardening rate, high ductility, continuous yield behavior, room temperature aging resistance, and bake hardenability. In addition, by controlling the fraction of each phase, recrystallization degree, distribution uniformity, etc., high-strength steel with high hole expansion can be manufactured.

However, in order to ensure ultra-high tensile strength of 1100 MPa or more, the fraction of hard phases such as martensite phases that contribute to improving strength needs to be increased. In this case, the yield strength increases, so there is a problem of defects such as cracks occurring during stamping.

Generally, DP steel for automobiles is manufactured by using steelmaking and continuous casting processes to make slabs, which are then subjected to [heating-rough rolling-hot finishing rolling] to obtain hot-rolled coils, and then annealing to manufacture the final product.

Among them, the annealing process is mainly a process carried out when manufacturing cold-rolled steel sheets. The cold-rolled steel sheets are manufactured as follows: the hot-rolled coils are pickled to remove the surface oxide scale, and cold-rolled at a specified reduction rate at room temperature, and then annealed and further flattened as needed.

The cold-rolled steel sheet (cold-rolled material) obtained by cold rolling is in a very hardened state and is not suitable for making parts that require workability. Therefore, the workability can be improved by softening it through heat treatment in a continuous annealing furnace as a subsequent process.

As an example, the annealing process is to heat a steel plate (cold rolled material) to about 650-850° C. in a heating furnace and then hold the steel plate for a certain period of time, thereby reducing hardness and improving workability through recrystallization and phase transformation phenomena.

The steel sheet that has not been annealed has high hardness, especially surface hardness, and insufficient workability. However, the steel sheet that has been annealed has a recrystallized structure, so that the hardness, yield point, and tensile strength are reduced, which can help improve the workability.

As a representative method for reducing the yield strength of DP steel, in the heating process during continuous annealing, ferrite is completely recrystallized to form an equiaxed crystal form, so that when austenite is formed and grown in the subsequent process, it becomes an equiaxed crystal form, which is conducive to the formation of an austenite phase with a small grain size and uniformity.

In addition, as a prior art for improving the workability of high-strength steel, Patent Document 1 proposes a method based on microstructure refinement, specifically, a method for dispersing fine precipitated copper particles with a particle size of 1-100 nm in the microstructure of a multi-phase steel plate mainly composed of a martensite phase. However, this technology requires the addition of 2-5% Cu to obtain good fine precipitated phase particles, so red hot brittleness caused by a large amount of Cu may occur, and there is a problem of excessive increase in manufacturing costs.

Patent document 2 discloses a steel plate having a ferrite matrix structure and a structure containing 2-10% pearlite by area, and improving strength by precipitation strengthening and grain refinement caused by adding carbonitride-forming elements (e.g., Ti, etc.). The steel plate has good hole expandability, but has limitations in further improving tensile strength, and has a problem of cracking during stamping due to high yield strength and low ductility.

Patent document 3 discloses a method for manufacturing a cold-rolled steel sheet that simultaneously obtains high strength and high ductility utilizing the tempered martensite phase and has an excellent sheet shape after continuous annealing. However, the carbon (C) content in the steel is as high as 0.2% or more, so in addition to the problem of poor weldability, there is also the possibility of furnace pit defects due to the addition of a large amount of Si.

In view of the above-mentioned prior art, in order to improve the formability such as hole expandability of high-strength steel that satisfies the physical properties required for utilization such as weldability, it is necessary to develop a method that can form a uniform structure in steel while lowering the yield strength and improving the ductility.

[Prior art literature]

[Patent Document]

(Patent Document 1) Japanese Patent Publication No. 2005-264176

(Patent Document 2) Korean Patent Publication No. 2015-0073844

(Patent Document 3) Japanese Patent Publication No. 2010-090432

Summary of the invention

Technical issues to be solved

One aspect of the present invention provides a high-strength steel sheet which has a low yield ratio and high strength and is excellent in formability such as hole expandability due to improved ductility, and a method for producing the same, as a material suitable for automobile structural parts and the like.

The technical problems to be solved by the present invention are not limited to the above contents. The technical problems to be solved by the present invention can be understood from the overall content of this specification, and those skilled in the art of the present invention can easily understand the additional technical problems of the present invention.

Technical Solution

One aspect of the present invention provides a high-strength steel sheet having excellent hole expandability, wherein the high-strength steel sheet comprises, by weight%, carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): 1.2% or less (except 0%), chromium (Cr): 0.1% or less (except 0%), molybdenum (Mo): 0.1% or less (except 0%), niobium (Nb): 0.1% or less (except 0%), titanium (Ti): 0.1% or less (except 0%). %), boron (B): less than 0.002% (except 0%), aluminum (sol.Al): 0.02-0.05%, phosphorus (P): less than 0.05% (except 0%), sulfur (S): less than 0.01% (except 0%), nitrogen (N): less than 0.01% (except 0%), iron (Fe) and other inevitable impurities, the fine structure contains ferrite with an area fraction of 20-30%, non-equilibrium ferrite of 5-15% and the balance of martensite.

Another aspect of the present invention provides a method for manufacturing a high-strength steel plate with excellent hole expansion performance, characterized in that the method includes the following steps: preparing a steel billet containing the above-mentioned alloy composition; heating the steel billet within a temperature range of 1100-1300°C; hot rolling the heated steel billet to manufacture a hot-rolled steel plate; coiling the hot-rolled steel plate within a temperature range of 400-700°C; cooling the coiled hot-rolled steel plate to room temperature; cold rolling the cooled hot-rolled steel plate to manufacture a cold-rolled steel plate; continuously annealing the cold-rolled steel plate; after the continuous annealing, performing a primary cooling at an average cooling rate of 1-10°C/second to a temperature range of 570-630°C; and after the primary cooling, performing a secondary cooling at an average cooling rate of 5-50°C to a temperature range of 300-400°C, wherein the continuous annealing is performed in an apparatus provided with a heating zone, a soaking zone and a cooling zone, and the heating zone and the soaking zone are controlled in a temperature range of 810-850°C.

Beneficial Effects

According to the present invention, it is possible to provide a steel sheet having excellent hole expandability even though having high strength, thereby improving formability and collision resistance.

As described above, the steel sheet of the present invention with improved formability can prevent processing defects such as cracks and wrinkles during press forming, and thus has the effect of being suitable for use in parts that need to be processed into complex shapes, etc. Furthermore, it is also effective to manufacture materials with improved collision resistance, so that when a car using such a part inevitably collides, defects such as cracks are not easily formed.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows the thermal history and phase change history during continuous annealing according to one embodiment of the present invention.

FIG. 2 (a) shows a void formation mechanism in a structure, and FIG. 2 (b) shows an interface strengthening mechanism in a structure of an inventive example according to an embodiment of the present invention.

FIG. 3 shows microstructure photographs of an inventive example and a comparative example according to an embodiment of the present invention.

Best Mode for Carrying Out the Invention

The inventors of the present invention have conducted intensive studies to develop a material having a moldability at a level suitable for use in automotive parts that need to be processed into a complex shape.

In particular, the inventors have derived a microstructure that can eliminate the hardness difference between the soft phase and the hard phase that affects the crack resistance of steel, and at the same time confirmed that the goal can be achieved by controlling the refinement of the hard phase and the shape of the grains that are conducive to preventing the generation and propagation of pores, thereby completing the present invention.

In particular, the present invention introduces an intermediate phase, preferably a non-equilibrium ferrite phase, to eliminate the hardness difference between the soft phase and the hard phase, and the technical significance of forming such a structure lies in optimizing the alloy composition and manufacturing conditions.

Hereinafter, the present invention will be described in detail.

According to one aspect of the present invention, a high-strength steel plate with excellent hole expandability and ductility may contain, in terms of weight%, carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): 1.2% or less (except 0%), chromium (Cr): 0.1% or less (except 0%), molybdenum (Mo): 0.1% or less (except 0%), niobium (Nb): 0.1% or less (except 0%), titanium (Ti): 0.1% or less (except 0%), boron (B): 0.002% or less (except 0%), aluminum (sol.Al): 0.02-0.05%, phosphorus (P): 0.05% or less (except 0%), sulfur (S): 0.01% or less (except 0%), and nitrogen (N): 0.01% or less (except 0%).

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

In addition, unless otherwise specified, the content of each element in the present invention is based on weight, and the proportion of the structure is based on area.

Carbon (C): 0.05-0.12%

Carbon (C) is an important element added for solid solution strengthening. This C combines with precipitation elements to form fine precipitates, thereby contributing to improving the strength of steel.

When the C content exceeds 0.12%, the hardenability increases and martensite is formed during the cooling process when the steel is manufactured, so the strength increases excessively, and there is also a problem of causing a decrease in elongation. In addition, due to the deterioration of weldability, there is a possibility of welding defects when processing into parts. In addition, when the C content is less than 0.05%, it is difficult to ensure the target level of strength.

Therefore, the C content may be 0.05-0.12%. More preferably, the C content may be 0.06% or more and 0.10% or less.

Manganese (Mn): 2.5-3.0%

Manganese (Mn) is an element that precipitates sulfur (S) in steel as MnS, thereby preventing hot brittleness caused by the formation of FeS and contributing to solid solution strengthening of steel.

When the content of such Mn is less than 2.5%, not only the above-mentioned effects cannot be obtained, but also it is difficult to ensure the strength at the target level. On the other hand, when the content of Mn exceeds 3.0%, the possibility of problems such as weldability and hot rolling is high, and martensite is more likely to be formed due to the increase in hardenability, so ductility may be reduced. In addition, too many Mn bands (Bands) (Mn oxide bands) are formed in the structure, so there is a problem of increasing the risk of defects such as processing cracks. In addition, Mn oxides are dissolved from the surface during annealing, so there is a problem of greatly hindering the plating property.

Therefore, the content of Mn may be 2.5-3.0%.

Silicon (Si): 1.2% or less (except 0%)

Silicon (Si) is a ferrite stabilizing element and helps to ensure the target level of ferrite fraction by promoting ferrite transformation. In addition, due to its excellent solid solution strengthening ability, it is effective in improving the strength of ferrite and is a useful element for ensuring strength without reducing the ductility of steel.

When the Si content exceeds 1.2%, the solid solution strengthening effect is excessive, which in turn reduces ductility and causes surface scale defects, thereby adversely affecting the quality of the plated surface. In addition, there is a problem of hindering chemical treatability.

Therefore, the content of Si may be 1.2% or less, and 0% may be excluded. More preferably, the content of Si may be 0.1% or more.

Chromium (Cr): 0.1% or less (except 0%)

Chromium (Cr) is an element that contributes to the formation of the desired structure of the present invention, and suppresses the formation of martensite and bainite phases during annealing heat treatment, while forming fine carbides, thereby contributing to improving strength. That is, Cr has the effect of suppressing the formation of bainite that competes with non-equilibrium ferrite, so when an appropriate level of Cr is contained, it is conducive to the formation of non-equilibrium ferrite phase at high temperature.

When the Cr content exceeds 0.1%, the non-equilibrium ferrite phase cannot be formed, so the ductility and hole expandability of the steel decrease, and when carbides are formed at the grain boundaries, the strength and elongation may deteriorate. In addition, there is a problem of increased manufacturing cost.

Therefore, the Cr content may be 0.1% or less, and 0% may be excluded. More preferably, the Cr content may be 0.01% or more.

Molybdenum (Mo): 0.1% or less (except 0%)

Molybdenum (Mo) is an element that promotes the formation of a non-equilibrium ferrite phase by suppressing the phase transformation of pearlite, suppresses the formation of a martensite phase during annealing heat treatment, and forms fine carbides to contribute to improving the strength.

When the Mo content exceeds 0.1%, the hardenability is excessive and the non-equilibrium ferrite phase cannot be formed. Therefore, the ductility and hole expandability of the steel may be reduced, and there is a problem of increased manufacturing cost.

Therefore, the content of Mo may be 0.1% or less, and 0% may be excluded. More preferably, the content of Mo may be 0.01% or less.

Niobium (Nb): 0.1% or less (except 0%)

Niobium (Nb) is an element that segregates at austenite grain boundaries to suppress coarsening of austenite grains during annealing heat treatment and forms fine carbides to contribute to improving strength.

When the content of such Nb exceeds 0.1%, coarse carbides are precipitated, and the strength and elongation may be poor due to the reduction in the carbon content in the steel, and there is a problem of increased manufacturing cost.

Therefore, the Nb content may be 0.1% or less, and 0% may be excluded. More preferably, the Nb content may be 0.01% or less.

Titanium (Ti): 0.1% or less (except 0%)

Titanium (Ti) is an element that forms fine carbides and helps to ensure yield strength and tensile strength. In addition, Ti precipitates N in steel as TiN, thereby having the effect of suppressing the formation of AlN caused by Al that inevitably exists in steel, and thus has the effect of reducing the possibility of cracking during continuous casting.

When the Ti content exceeds 0.1%, coarse carbides are precipitated, and the strength and elongation may be reduced due to the reduction in the carbon content in the steel. In addition, there is a possibility of causing nozzle clogging during continuous casting, and there is a problem of increased manufacturing cost.

Therefore, the content of Ti may be 0.1% or less, and 0% may be excluded. More preferably, the content of Ti may be 0.01% or less.

Boron (B): 0.002% or less (except 0%)

Boron (B) is an element that delays the transformation of austenite into pearlite during cooling after annealing heat treatment, but when the B content exceeds 0.002%, excessive B is concentrated on the surface, thereby possibly causing deterioration in plating adhesion.

Therefore, the content of B may be 0.002% or less, and 0% may be excluded.

Aluminum (sol.Al): 0.02-0.05%

Aluminum (sol.Al) is an element added for the purpose of grain refinement and deoxidation of steel. When the content of aluminum (sol.Al) is less than 0.02%, aluminum-killed steel cannot be produced in a stable state. On the other hand, when the content of aluminum (sol.Al) exceeds 0.05%, grain refinement occurs, thereby having the effect of improving strength, but excessive inclusions are formed during the steelmaking and continuous casting operation, thereby increasing the possibility of surface defects of the plated steel sheet.

Therefore, the content of the acid-soluble aluminum (sol.Al) may be 0.02-0.05%.

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

Phosphorus (P) is a substitutional element with the greatest solid solution strengthening effect, and is an element that improves in-plane anisotropy without significantly reducing formability while ensuring strength. However, when too much P is added, the possibility of brittle failure increases significantly, resulting in an increased possibility of plate breakage during hot rolling, and there is a problem of hindering the surface properties of the plated surface.

Therefore, in the present invention, the P content may be controlled to be less than 0.05%, except for 0% in consideration of the level that is inevitably added.

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

Sulfur (S) is an impurity element in steel and an element that is inevitably added. Since it hinders ductility, it is preferable to control the sulfur content to a minimum. In particular, S has the problem of increasing the possibility of red-hot brittleness, so it is preferable to control the sulfur content to 0.01% or less. However, considering the level that is inevitably added during the manufacturing process, 0% may be excluded.

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

Nitrogen (N) is a solid solution strengthening element, but when the content of the nitrogen (N) exceeds 0.01%, the risk of generating brittleness increases, and it combines with Al in the steel to precipitate excessive AlN, thus possibly hindering the continuous casting quality.

Therefore, the N content may be 0.01% or less, and 0% may be excluded in consideration of the level that is inevitably added.

The remaining component of the present invention is iron (Fe). However, undesirable impurities may inevitably be mixed from the raw materials or the surrounding environment during the conventional manufacturing process, so this impurity cannot be excluded. These impurities are well known to technicians in the conventional manufacturing process, so all of them are not specifically described in this specification.

The fine structure of the steel sheet of the present invention having the above alloy composition can be composed of ferrite as a soft phase, a martensite phase as a hard phase, and a non-equilibrium ferrite phase formed at the interface between them.

Specifically, the steel sheet of the present invention includes a ferrite phase with an area fraction of 20-30%, a non-equilibrium ferrite phase with an area fraction of 5-15%, and may include a martensite phase as a residual structure. In addition, a trace amount of retained austenite phase may be included.

In the present invention, the non-equilibrium ferrite phase is a structure that is conducive to minimizing the hardness difference between the soft phase and the hard phase, and is a structure different from the existing equilibrium ferrite (polygonal ferrite). The non-equilibrium ferrite may be acicular ferrite or bainitic ferrite. In addition, depending on the cooling conditions, Widmanstattenferrite, massive ferrite, etc. may be included. Specifically, the non-equilibrium ferrite is affected by the components constituting the mother phase, and contains relatively high C and Mn compared to the equilibrium ferrite. For example, in the case of equilibrium ferrite, assuming that the C concentration is 0.02%, the non-equilibrium ferrite has a higher C content of 0.03-0.04%.

Therefore, the C concentration and Mn concentration of the hard phase formed near (around) the non-equilibrium ferrite are relatively reduced, so the hardness difference between the soft phase and the hard phase is reduced and the effect of improving the hole expandability can be obtained. In addition, when the Si concentration in the non-equilibrium ferrite is as low as less than 1%, the stacking fault energy increases, and the cross slip becomes difficult, and the formation of voids due to deformation is hindered (Figure 2).

When the fraction of such non-equilibrium ferrite phase is too high, the fraction of the hard phase is relatively reduced, so the strength at the target level cannot be ensured. In view of this, the non-equilibrium ferrite phase may be contained at 15% or less. On the other hand, when the fraction of the non-equilibrium ferrite phase is less than 5%, the above-mentioned effect (minimizing the hardness difference between the hard phase and the soft phase) cannot be fully obtained, so the hole expandability deteriorates.

When the fraction of the ferrite phase is less than 20%, it is disadvantageous to ensure the ductility of the steel, but when the fraction of the ferrite phase exceeds 30%, the fraction of the hard phase is relatively reduced, so it is difficult to ensure the target level of strength.

In the structure other than the ferrite phase and the non-equilibrium ferrite phase, the fraction of the martensite phase is not particularly limited, but in order to ensure ultra-high strength of a tensile strength of 1100 MPa or more, the martensite phase may be included with an area fraction of 50% or more. However, when the fraction of the martensite phase exceeds 75%, ductility decreases, making it difficult to ensure the target level of formability.

In addition, with respect to the retained austenite phase, it is favorable for the fraction to be not more than 3%, and even if the fraction is 0%, there is no problem in ensuring the desired physical properties.

The steel sheet of the present invention having the above-mentioned fine structure may have a tensile strength of 1100 MPa or more, a yield strength of 550-700 MPa, and an elongation (total elongation) of 12% or more, and thus has the characteristics of high strength and high ductility.

In addition, the steel plate has a hole expansion ratio (HER) of more than 25%, and thus has excellent resistance to cracks that may be generated during processing and resistance to collision fracture.

Hereinafter, a method for producing a high-strength steel sheet having excellent hole expandability and ductility according to another aspect of the present invention will be described in detail.

In short, the present invention can manufacture the desired steel plate through the process of [steel billet heating-hot rolling-coiling-cold rolling-continuous annealing], and each process is described in detail below.

[Heating of steel billet]

First, a steel billet satisfying the above-mentioned alloy composition may be prepared and then heated.

This process is carried out in order to smoothly carry out the subsequent hot rolling process and fully obtain the desired physical properties of the steel plate. In the present invention, there is no particular restriction on the conditions of this heating process, as long as it is a common condition. As an example, the heating process can be carried out in a temperature range of 1100-1300°C.

[Hot rolling]

The steel slab heated as described above may be hot rolled to produce a hot rolled steel sheet, and at this time, hot finish rolling may be performed at an outlet side temperature of Ar3 or more and 1000° C. or less.

When the outlet temperature during the hot finish rolling is lower than Ar3, the thermal deformation resistance increases sharply, and the top, tail and edge of the hot rolled coil become single-phase regions, so the in-plane anisotropy increases, which may deteriorate the formability. In addition, when the outlet temperature during the hot finish rolling exceeds 1000°C, the rolling load is relatively reduced, which is beneficial to productivity, but there is a possibility of generating thick oxide scale.

More specifically, the hot finish rolling may be performed within a temperature range of 760-940°C.

[Collect]

The hot-rolled steel sheet manufactured as described above can be rolled into a coil shape.

The coiling can be performed at a temperature range of 400-700° C. When the coiling temperature is lower than 400° C., excessive martensite phase or non-equilibrium ferrite phase is formed, causing excessive increase in the strength of the hot-rolled steel sheet, so that problems such as shape defects due to load may occur during subsequent cold rolling. On the other hand, when the coiling temperature exceeds 700° C., there is a problem of deterioration in pickling properties due to increase in surface oxide scale.

[cool down]

Preferably, the coiled hot-rolled steel sheet is cooled to room temperature at an average cooling rate of 0.1°C/s or less (except 0°C/s). At this time, the coiled hot-rolled steel sheet may be cooled after being transported, placed, etc., and the process before cooling is not limited thereto.

As described above, by cooling the coiled hot-rolled steel sheet at a predetermined speed, a hot-rolled steel sheet in which carbides serving as nucleation sites of austenite are finely dispersed can be obtained.

[Cold rolling]

The hot-rolled steel sheet coiled as described above may be cold-rolled to manufacture a cold-rolled steel sheet.

In the present invention, the cold rolling can be performed at a cold rolling reduction ratio of 55-70%. When the cold rolling reduction ratio is less than 55%, the recrystallization driving force is weakened, and it is difficult to obtain good recrystallized grains. On the other hand, when the cold rolling reduction ratio exceeds 70%, the risk of cracks at the edge of the steel plate increases, and there is a possibility that the rolling load increases sharply.

In the present invention, when an appropriate level of cold rolling reduction ratio is applied during cold rolling, the recrystallization of ferrite can be further promoted in the heating zone during the subsequent continuous annealing process, thereby inducing the formation of fine ferrite, so that small and uniform austenite can be formed at the ferrite grain boundary. This affects the size or distribution of the non-equilibrium structure during the cooling process, and is conducive to maintaining the strength of the final product while improving the elongation, hole expansion and other processing properties.

In addition, the cold rolling reduction ratio can be achieved by only one cold rolling, that is, by only one stand, and the steel can be reduced as described above, which has an economical effect.

However, in the case of a thick plate steel material having a thickness of 6 mm or more of the hot rolled steel sheet before cold rolling, the target reduction ratio can be achieved by repeatedly rolling using a reversing mill. In this case, the number of all passes of repeated rolling can be set to 1 stand. A reversing mill is a rolling mill for rolling thin plate steel materials, which refers to a rolling mill that rolls a material by reciprocating between a pair of rolls, and a single pass of the material during reciprocation can be set to 1 pass.

The present invention can perform pickling treatment on the hot-rolled steel sheet before the cold rolling, and the pickling treatment process can be performed by a conventional method.

[Continuous annealing]

Preferably, the cold rolled steel sheet manufactured as described above is subjected to a continuous annealing process. As an example, the continuous annealing process can be performed in a continuous annealing furnace (CAL).

Generally, a continuous annealing furnace (CAL) may be composed of [a heating zone - a soaking zone - a cooling zone (a slow cooling zone and a rapid cooling zone) - (an aging zone as required)], and undergoes the following process: a cold-rolled steel sheet is loaded into the continuous annealing furnace as described above, and then heated at a specific temperature in the heating zone, and maintained in the soaking zone for a specified time after reaching the target temperature.

In the present invention, the temperatures of the heating zone and the soaking zone during the continuous annealing can be controlled to be the same, which means that the end temperature of the heating zone and the start temperature of the soaking zone are controlled to be the same ( FIG. 1 ).

Specifically, the temperature of the heating zone and the soaking zone can be controlled at 810-850°C.

When the temperature of the heating zone is lower than 810°C, sufficient heat input for recrystallization cannot be applied. On the other hand, when the temperature of the heating zone exceeds 850°C, productivity decreases and excessive austenite phase is formed, so that the fraction of the hard phase after subsequent cooling is greatly increased, and there is a possibility that the ductility of the steel will deteriorate.

In addition, when the temperature of the soaking zone is lower than 810°C, excessive cooling is required at the end temperature of the heating zone, which is disadvantageous in economy and the heat for recrystallization may be insufficient. On the other hand, when the temperature of the soaking zone exceeds 850°C, the fraction of austenite is too large, the hard phase increases during cooling, and thus the formability may be reduced.

When the temperature of the soaking zone is increased within the above temperature range, the stability of austenite can be reduced, thereby promoting the formation of a non-equilibrium ferrite phase in the subsequent cooling process.

Although detailed description will be given later, in the present invention, cooling is performed in stages when cooling after passing through the heating zone and the soaking zone, the non-equilibrium ferrite phase is introduced after one cooling, and the final structure can be composed of a certain fraction of soft phase, hard phase and non-equilibrium ferrite phase. Therefore, in the steel plate of the present invention, not only the strength and ductility can be improved, but also the effect of improving the workability can be obtained at the same time due to the interface strengthening of the non-equilibrium ferrite phase.

Therefore, in order to obtain the desired fine structure of the present invention, it is preferable to control the heat input to the steel sheet in the heating zone consisting of the heating zone and the soaking zone during the continuous annealing.

[Stage Cooling]

As described above, a desired structure can be formed by cooling the cold-rolled steel sheet subjected to the heat treatment described above, and in this case, stepwise cooling is preferably performed.

In the present invention, the staged cooling may consist of primary cooling-secondary cooling. Specifically, after the continuous annealing, primary cooling may be performed at an average cooling rate of 1-10°C/second to a temperature range of 570-630°C, and then secondary cooling may be performed at an average cooling rate of 5-50°C/second to a temperature range of 300-400°C.

At this time, the primary cooling is performed more slowly than the secondary cooling, so that defects in the sheet shape caused by a rapid temperature drop during the secondary cooling which is the subsequent relatively rapid cooling period can be suppressed.

When the termination temperature during the primary cooling is lower than 570°C, the diffusion activity of carbon is low due to the too low temperature, and the carbon concentration in ferrite increases. On the other hand, since the carbon concentration in austenite becomes low, the fraction of the hard phase becomes too large, thereby increasing the yield ratio, thereby increasing the tendency of cracks to occur during processing. In addition, since the cooling rate between the soaking zone and the cooling zone (slow cooling zone) is too fast, the shape of the plate becomes uneven. When the termination temperature exceeds 630°C, an excessively high cooling rate is required during subsequent cooling (secondary cooling), and it is difficult to introduce a non-equilibrium ferrite phase.

Furthermore, when the average cooling rate during the primary cooling exceeds 10° C./sec, carbon diffusion cannot be sufficiently performed. In addition, in consideration of productivity, the primary cooling may be performed at an average cooling rate of 1° C./sec or more.

As described above, after the primary cooling is completed, rapid cooling (secondary cooling) may be performed at a cooling rate above a certain level. At this time, when the secondary cooling termination temperature is lower than 300°C, cooling deviation occurs in the width direction and the length direction of the steel plate, and there is a possibility that the plate shape may deteriorate. On the other hand, when the secondary cooling termination temperature exceeds 400°C, the hard phase cannot be fully ensured, so the strength may be reduced, and due to the formation of bainite, the yield strength ratio may increase and the elongation may decrease.

When the average cooling rate during the secondary cooling is less than 5° C./sec, the fraction of the hard phase may be too high, whereas when the average cooling rate during the secondary cooling exceeds 50° C./sec, the hardness may be insufficient.

In addition, as required, after the step-by-step cooling is completed, an aging treatment may be performed.

The over-aging treatment is a process of maintaining the temperature for a certain period of time after the secondary cooling termination temperature, and uniformly heat treating the coil along the width and length directions, thereby having the effect of improving the shape quality. To this end, the over-aging treatment can be carried out for 200-800 seconds.

The overaging treatment may be performed after the secondary cooling is terminated, so the temperature of the overaging treatment may be the same as the secondary cooling termination temperature, or the overaging treatment may be performed within the secondary cooling termination temperature range.

The microstructure of the high-strength steel sheet of the present invention manufactured as described above is composed of a hard phase and a soft phase. In particular, the recrystallization of ferrite is maximized by optimizing the cold rolling and annealing processes, so that a structure in which martensite as a hard phase is uniformly distributed on the final recrystallized ferrite matrix can be obtained. In addition, by introducing a non-equilibrium ferrite phase at the interface between the hard phase and the soft phase, the crack resistance during processing is improved.

Therefore, even if the steel sheet of the present invention has a high strength of 1100 MPa or more in tensile strength, it is possible to ensure excellent formability such as excellent hole expandability by ensuring a low yield ratio and high ductility.

Hereinafter, the present invention will be described in more detail by way of examples. However, the description of these examples is only for illustrating the implementation of the present invention, and the present invention is not limited to the description of these examples. This is because the scope of the present invention is determined by the contents described in the claims and the contents reasonably derived therefrom.

DETAILED DESCRIPTION

(Example)

Steel billets having the alloy composition shown in Table 1 below were manufactured, and then each steel billet was heated at 1200°C for 1 hour, and then hot-rolled at a finishing temperature of 880-920°C to manufacture hot-rolled steel sheets. After that, each hot-rolled steel sheet was coiled at 650°C and cooled to room temperature at a cooling rate of 0.1°C/second. Thereafter, the coiled hot-rolled steel sheet was cold-rolled and continuously annealed according to the conditions shown in Table 2 below, and then cooled in stages (primary cooling-secondary cooling), and then aged at 360°C for 520 seconds to manufacture the final steel sheet.

At this time, the primary cooling in the step cooling was performed at an average cooling rate of 3° C./sec, and the secondary cooling was performed at an average cooling rate of 20° C./sec. In addition, the cold rolling was performed by one stand.

The microstructure of each steel plate manufactured as described above was observed, and physical property indices used in processing such as tensile and processing characteristics and hole expansion ratio were evaluated. The results are shown in Table 3 below.

At this time, the tensile test on each test piece was performed by sampling a tensile test piece of JIS No. 5 size in a direction perpendicular to the rolling direction and then performing a tensile test at a strain rate of 0.01/sec.

In addition, the hole expansion (HER, %) measurement test was carried out according to ISO16630. Specifically, when a circular hole is punched on the test piece and then expanded using a conical punch, it is expressed as the ratio of the hole expansion amount until the crack generated at the edge of the hole penetrates in the thickness direction to the initial hole. At this time, the test piece size is 120mm×120mm, the clearance is 12%, the punching diameter is 10mm, the punching holding load is 20 tons (ton), and the test speed is set to 12mm/minute.

In addition, the martensite phase and the non-equilibrium ferrite phase corresponding to the hard phase in the organizational phase were observed by SEM at 2000 times and 5000 times magnification after etching with nitric acid etching solution. At this time, the size, fraction, etc. of each observed phase were measured. In addition, for the phase, each fraction was measured using SEM and an image analyzer program after etching with nitric acid etching solution.

[Table 1]

[Table 2]

[Table 3]

As shown in Tables 1 to 3, it can be seen that the steel alloy composition and manufacturing conditions meet the conditions proposed by the present invention, and in particular, the invention examples 1 to 6, in which both the cold rolling and continuous annealing processes meet the processes proposed by the present invention, achieve sufficient recrystallization of ferrite and form a fine hard phase during the annealing process after cold rolling, and are connected by a non-equilibrium ferrite structure at the interface, so that they have high strength and appropriate yield strength for plate processing, and have excellent elongation. In addition, due to the excellent hole expansion property, it can be confirmed that the target level of formability can be ensured.

On the other hand, in Comparative Examples 1 to 6 where the soaking temperature during continuous annealing in the manufacturing process of the steel sheet was low, recrystallization did not sufficiently occur, and since the appropriate fraction of austenite formed in the soaking zone had high stability, non-equilibrium ferrite was not sufficiently introduced during the cooling process. As a result, the ductility and/or hole expandability were poor.

In addition, although Comparative Examples 7 to 10 were heated at an appropriate temperature during continuous annealing, the termination temperature during the primary cooling was high, and the introduction time of non-equilibrium ferrite during the cooling process was insufficient, resulting in poor ductility and/or hole expandability.

In addition, in Comparative Examples 11 to 14 containing excessive amounts of Cr as a hardenability element, there is a risk of cracking during processing due to excessively high yield strength, and the non-equilibrium ferrite phase cannot be introduced due to the low temperature in the soaking zone, resulting in poor ductility in some comparative examples.

FIG. 3 shows microstructure photographs of Comparative Examples 4 to 7 and Inventive Example 1. In FIG.

As shown in FIG. 3 , in Inventive Example 1, a uniform and fine non-equilibrium ferrite phase is introduced into a sufficient fraction of the recrystallized ferrite matrix during the primary cooling process, and a certain fraction of the martensite phase is formed during the secondary cooling process.

On the other hand, in Comparative Examples 4 to 7, it was confirmed that a small amount of non-equilibrium ferrite was introduced due to the conditions of the soaking zone temperature or the primary cooling end temperature during continuous annealing. Among them, in Comparative Example 4 where the soaking zone temperature did not reach 800°C and the primary cooling end temperature was quite high, and in Comparative Example 7 where the primary cooling end temperature was quite high, it was found that the non-equilibrium ferrite was less than 1% and was hardly observed.

Claims (10)

1. A high-strength steel sheet excellent in hole expansibility, comprising, in weight%: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): 1.2% or less, excluding 0%, chromium (Cr): 0.1% or less except 0% molybdenum (Mo): 0.1% or less except 0% niobium (Nb): 0.1% or less and 0% excluded, titanium (Ti): 0.1% or less except 0% boron (B): 0.002% or less except 0% aluminum (sol.al): 0.02-0.05%, phosphorus (P): 0.05% or less, excluding 0%, sulfur (S): 0.01% or less, excluding 0%, nitrogen (N): the microstructure contains 20-30% ferrite by area fraction, 5-15% unbalanced ferrite and the balance martensite, except for 0.01% and 0% of iron (Fe) and other unavoidable impurities.

2. The high-strength steel sheet excellent in hole expansibility according to claim 1, wherein the steel sheet contains a martensite phase having an area fraction of 50% or more.

3. The high-strength steel sheet excellent in hole expansibility according to claim 1, wherein the steel sheet further comprises a residual austenite phase having an area fraction of 3% or less and including 0%.

4. The high-strength steel sheet excellent in hole expansibility according to claim 1, wherein the steel sheet has a tensile strength of 1100MPa or more, a yield strength of 550 to 700MPa, and a total elongation of 12% or more.

5. The high-strength steel sheet excellent in hole expansibility according to claim 1, wherein the steel sheet has a Hole Expansibility (HER) of 25% or more.

6. A method of manufacturing a high-strength steel sheet excellent in hole expansibility, the method comprising the steps of:

preparing a billet comprising, in weight percent: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): 1.2% or less, excluding 0%, chromium (Cr): 0.1% or less except 0% molybdenum (Mo): 0.1% or less except 0% niobium (Nb): 0.1% or less and 0% excluded, titanium (Ti): 0.1% or less except 0% boron (B): 0.002% or less except 0% aluminum (sol.al): 0.02-0.05%, phosphorus (P): 0.05% or less, excluding 0%, sulfur (S): 0.01% or less, excluding 0%, nitrogen (N): 0.01% or less and 0% excluding iron (Fe) and other unavoidable impurities;

heating the steel billet at the temperature of 1100-1300 ℃;

hot-rolling the heated steel slab to manufacture a hot-rolled steel sheet;

rolling the hot rolled steel plate at the temperature of 400-700 ℃;

cooling the rolled hot rolled steel plate to normal temperature;

cold rolling the cooled hot rolled steel sheet to manufacture a cold rolled steel sheet;

continuously annealing the cold-rolled steel plate;

after the continuous annealing, cooling at an average cooling rate of 1-10 ℃/sec for the first time to a temperature range of 570-630 ℃; and

after the primary cooling, carrying out secondary cooling at an average cooling speed of 5-50 ℃ to a temperature range of 300-400 ℃,

wherein the continuous annealing is performed in an apparatus provided with a heating zone, a soaking zone and a cooling zone, the heating zone and the soaking zone being controlled in a temperature range of 810-850 ℃.

7. The method for producing a high-strength steel sheet excellent in hole expansibility according to claim 6, wherein the hot rolling is hot finish rolling at an outlet side temperature Ar3 or more and 1000 ℃ or less.

8. The method for producing a high-strength steel sheet excellent in hole expansibility according to claim 6, wherein the cooling after rolling is performed at a cooling rate of 0.1 ℃/sec or less except for 0 ℃.

9. The method for producing a high-strength steel sheet excellent in hole expansibility according to claim 6, wherein the cold rolling is performed in one stand and the total reduction is 55 to 70%.

10. The method for producing a high-strength steel sheet excellent in hole expansibility according to claim 6, wherein after said secondary cooling, further comprising a step of performing an overaging treatment for 200 to 800 seconds.