CN117043382A

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

Info

Publication number: CN117043382A
Application number: CN202280017433.0A
Authority: CN
Inventors: 赵卿来; 金成圭; 朴俊澔; 韩箱浩
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2021-07-20
Filing date: 2022-06-22
Publication date: 2023-11-10
Also published as: KR20230014121A; WO2023003188A1

Abstract

The present invention relates to steel suitable for automotive materials, and more particularly, to a high-strength steel sheet excellent in hole expansibility and ductility and a method for manufacturing the same. The microstructure of the high strength steel sheet of the present invention is composed of a hard phase and a soft phase, and martensite phase as the hard phase is uniformly distributed on the recrystallized ferrite matrix by an optimized cold rolling and annealing process, and crack resistance during processing can be improved by introducing an unbalanced ferrite phase at the interface of the hard phase and the soft phase.

Description

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

Technical Field

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

Background

In recent years, in the field of automobile industry, CO has been used as a CO ₂ Emission-related environmental regulations and energy use regulations, high strength steel is required for the purpose of improving fuel efficiency or durability.

In particular, with the expansion of impact stability regulations of automobiles, high-strength steel excellent in strength is used as a material of structural members such as a frame (member), a seat rail (seat rail), a pillar (pilar) and the like for improving the impact resistance of a vehicle body.

Such automobile parts have a complicated shape in terms of stability and design, and are mainly molded and manufactured using a stamping die, and thus high strength and high level of moldability are required.

However, as the strength of steel increases, there is a feature that it is advantageous to absorb impact energy, but generally as the strength increases, elongation decreases, and thus there is a problem that formability decreases. In addition, when the yield strength is too high, there is a problem in that the formability is deteriorated and the manufacturing cost is increased because the introduction of the material into the mold is reduced during the forming.

Further, since the automobile parts have many molding sites that expand after the hole is formed, hole expansibility (Hole Expandability, HER) is required for smooth molding, but the hole expansibility of high-strength steel is low, and there is a problem that defects such as cracks (cracks) occur during the molding process. As described above, if the hole expansibility is poor, cracks are generated in the molded part of the member at the time of collision of the automobile, and the member is easily broken, so that the safety of passengers may be threatened.

In addition, typical high-strength steels used as automotive materials include Dual Phase Steel (DP Steel), transformation induced plasticity Steel (Transformation Induced Plasticity Steel, TRIP Steel), complex Phase Steel (Complex Phase Steel, CP Steel), ferrite-bainite Steel (Ferrite Bainite Steel, FB Steel), and the like.

DP steel, which is ultra-high tension steel, has a low yield ratio of about 0.5 to 0.6 level, and thus is easy to process, and has an advantage of high elongation inferior to TRIP steel. Therefore, it is mainly applied to outer doors, seat rails, seat belts, suspensions, arms, wheels, and the like.

The TRIP steel has a yield ratio in the range of 0.57 to 0.67, and thus has a characteristic of exhibiting excellent formability (high ductility), and is thus suitable for parts requiring high formability such as a skeleton member, a roof, a seat belt, a bumper, and the like.

CP steel has a low yield ratio, high elongation and bending workability, and thus is applied to side plates, underbody reinforcements, and the like, and FB steel has excellent hole expansibility, and thus is mainly applied to suspension lower arms, wheel discs, and the like.

Among them, DP steel mainly consists of ferrite excellent in ductility and a hard phase (martensite phase and bainite phase) high in strength, and a trace amount of retained austenite may be present. Such DP steel has low Yield strength and high tensile strength, and thus has excellent characteristics such as low Yield Ratio (YR), high work hardening rate, high ductility, continuous Yield behavior, aging resistance at ordinary temperature, bake hardenability, and the like. Further, by controlling the fraction, recrystallization degree, distribution uniformity, and the like of each phase (phase), high strength steel having high hole expansibility can be produced.

However, in order to secure an ultra-high strength of 1100MPa or more, it is necessary to increase the fraction of a hard phase (hardp phase) such as a martensite phase, which is advantageous for improving the strength, and in this case, the yield strength is increased, so that there is a problem that a defect such as a crack (ack) occurs during press forming.

In general, DP steel for automobiles is manufactured by manufacturing a slab through a steelmaking and continuous casting process, subjecting the slab to [ heating-rough rolling-hot finish rolling ] to obtain a hot rolled coil, and then annealing the hot rolled coil to manufacture a final product.

Wherein, the annealing process is mainly the process performed when manufacturing the cold-rolled steel sheet, and the cold-rolled steel sheet is manufactured as follows: the hot rolled coil is manufactured by pickling to remove surface scale, cold rolling at a prescribed rolling reduction at normal temperature, and then annealing and further temper rolling 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 producing a member requiring workability, and thus can be softened by heat treatment in a continuous annealing furnace as a subsequent process to improve workability.

As an example, the annealing process is to heat a steel sheet (cold rolled material) to about 650-850 ℃ in a heating furnace for a certain time, so that hardness can be reduced and workability can be improved through recrystallization and phase transformation phenomena.

The steel sheet not subjected to the annealing process has high hardness, particularly surface hardness, and insufficient workability, but the steel sheet subjected to the annealing process has a recrystallized structure, and thus the hardness, yield point, and tensile strength are reduced, and thus it is possible to contribute to improvement of workability.

As a representative method of reducing the yield strength of DP steel, in a heating process at the time of continuous annealing, an equiaxed crystal form is made by completely recrystallizing ferrite to be an equiaxed crystal form when austenite is formed and grown in a subsequent process, thus facilitating formation of an austenite phase having a small and uniform grain size.

Further, as a conventional technique for improving workability of high-strength steel, patent document 1 proposes a method according to structure refinement, and specifically discloses a method of dispersing fine precipitated copper particles having a particle size of 1 to 100nm in a structure of a complex-phase steel sheet mainly composed of a martensite phase. However, this technique requires addition of 2 to 5% of Cu to obtain good fine precipitated phase particles, and thus red hot shortness due to a large amount of Cu may occur, and there is a problem in that manufacturing cost excessively increases.

Patent document 2 discloses a steel sheet having a ferrite matrix structure, a structure containing 2 to 10 area% of pearlite (pearlite), and enhanced strength by precipitation strengthening and grain refinement by the addition of a carbon/nitride forming element (e.g., ti, etc.). The steel sheet is excellent in hole expansibility, but has a limitation in further improving tensile strength, and has a problem of generating cracks at the time of press forming due to high yield strength and low ductility.

Patent document 3 discloses a method for producing a cold-rolled steel sheet which has high strength and high ductility by using a tempered martensite phase and is excellent in shape even after continuous annealing, but since the content of carbon (C) in the steel is as high as 0.2% or more, there is a possibility that pit defects in the furnace may occur due to addition of a large amount of Si in addition to the problem of poor weldability.

In view of the above-described conventional techniques, in order to improve formability such as hole expansibility of high-strength steel satisfying physical properties such as weldability, it is necessary to develop a method of improving ductility while reducing yield strength while forming a uniform structure in the steel.

[ Prior Art literature ]

[ patent literature ]

(patent document 1) Japanese laid-open patent publication No. 2005-264176

(patent document 2) Korean laid-open patent publication No. 2015-0071874

(patent document 3) Japanese laid-open patent publication No. 2010-090432

Disclosure of Invention

Technical problem to be solved

An aspect of the present invention provides a high-strength steel sheet having a low yield ratio and high strength, which is suitable as a material for automobile structural members and the like, and having excellent formability such as hole expansibility by improving ductility, and a method for manufacturing the same.

The technical problem to be solved by the present invention is not limited to the above. Technical problem to be solved by the present invention will be understood from the entire contents of the present specification, and additional technical problems of the present invention will be easily understood by those skilled in the art to which the present invention pertains.

Technical proposal

An aspect of the present invention provides a high-strength steel sheet excellent in hole expansibility, comprising, in weight percent: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): less than 1.2% (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% (excluding 0%), iron (Fe) and other unavoidable impurities, and the microstructure contains 20-30% ferrite in area fraction, 5-15% unbalanced ferrite and the balance martensite.

Another aspect of the present invention provides a method for manufacturing a high-strength steel sheet excellent in hole expansibility, comprising the steps of: preparing a billet comprising the alloy composition; 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, performing secondary cooling at an average cooling rate 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, and the heating zone and the soaking zone are controlled in the temperature range of 810-850 ℃.

Advantageous effects

According to the present invention, it is possible to provide a steel sheet having excellent hole expansibility even with high strength, whereby formability and collision resistance are improved.

As described above, the steel sheet of the present invention having improved formability can prevent processing defects such as cracks and wrinkles during press forming, and therefore has an effect of being suitably applied to members for structures and the like which need to be processed into complex shapes. Furthermore, a material having improved collision resistance is also produced efficiently, so that defects such as cracks, which are likely to form, are avoided when an automobile using such a member inevitably collides.

Drawings

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

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

Fig. 3 shows a photograph of a microstructure of an invention example and a comparative example according to an embodiment of the present invention.

Best mode for carrying out the invention

The present inventors have conducted intensive studies to develop a material having a level of moldability that can be suitably used for an automobile material, such as a member that needs to be processed into a complicated shape.

In particular, the present inventors have found out a structure that can eliminate a difference in hardness between a soft phase and a hard phase affecting crack resistance of steel, and have confirmed that the object can be achieved by controlling the refinement of the hard phase and the shape of crystal grains that are advantageous in preventing the generation and propagation of pores, and completed the present invention.

In particular, in the present invention, an intermediate phase, preferably an unbalanced ferrite phase, is introduced to eliminate the difference in hardness between the soft phase and the hard phase, and in terms of forming such a structure, the technical significance is to optimize the alloy composition and the manufacturing conditions.

The present invention will be described in detail below.

The high-strength steel sheet excellent in hole expansibility and ductility according to an aspect of the present invention may include, in weight%: carbon (C): 0.05-0.12%, manganese (Mn): 2.5-3.0%, silicon (Si): less than 1.2% (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% (excluding 0%).

The reason why the alloy composition of the steel sheet provided in the present invention is limited as described above will be described in detail below.

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

Carbon (C): 0.05 to 0.12 percent

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

When the content of C exceeds 0.12%, martensite is formed during cooling when manufacturing steel due to an increase in hardenability, and thus strength excessively increases, and there is a problem of causing a decrease in elongation. Further, there is a possibility that welding defects may occur when the component is processed due to poor weldability. In addition, when the content of C is less than 0.05%, it is difficult to secure the strength of the target level.

Thus, the content of C may be 0.05 to 0.12%. The content of C may be more preferably 0.06% or more, and may be 0.10% or less.

Manganese (Mn): 2.5-3.0%

Manganese (Mn) is an element that precipitates sulfur (S) in steel into MnS, thereby preventing hot shortness caused by formation of FeS, and facilitating solid solution strengthening of steel.

When the content of such Mn is less than 2.5%, not only the above-described effects cannot be obtained, but also it is difficult to secure the strength of the target level. On the other hand, when the Mn content exceeds 3.0%, there is a high possibility that problems such as weldability and hot-rolling property occur, and martensite is more easily formed due to an increase in hardenability, so that ductility may be lowered. Further, too many Mn bands (Mn oxide bands) are formed in the structure, and thus there is a problem in that the risk of occurrence of defects such as processing cracks becomes large. Further, mn oxide is eluted from the surface during annealing, which has a problem of greatly impeding the plating property.

Thus, the Mn content may be 2.5-3.0%.

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

Silicon (Si) is a ferrite stabilization element, and facilitates securing a ferrite fraction at a target level by promoting ferrite transformation. Further, since the solid solution strengthening ability is excellent, it is a useful element for improving the strength of ferrite and ensuring the strength without reducing the ductility of steel.

When the content of such Si exceeds 1.2%, the solid solution strengthening effect is excessive, but ductility is lowered, and surface oxide defects are induced, thereby adversely affecting the quality of the plated surface. In addition, there is a problem that chemical treatability is hindered.

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

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

Chromium (Cr) is an element contributing to the structure desired in the present invention, and suppresses formation of a martensite phase and a bainite phase during annealing heat treatment, while forming fine carbides, thereby contributing to improvement of strength. That is, the Cr has an effect of suppressing bainite competing with unbalanced ferrite to form, and thus is advantageous in forming unbalanced ferrite phase at high temperature when Cr is contained at an appropriate level.

When the content of such Cr exceeds 0.1%, unbalanced ferrite phase is not formed on the contrary, and therefore ductility and hole expansibility of the steel are lowered, and when carbide is formed at grain boundaries, strength and elongation may be deteriorated. Further, there is a problem in that the manufacturing cost increases.

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

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

Molybdenum (Mo) is an element that promotes the formation of unbalanced ferrite phase by suppressing pearlite transformation, suppresses the formation of martensite phase during annealing heat treatment, and forms fine carbide to contribute to the improvement of strength.

When the content of such Mo exceeds 0.1%, hardenability is excessive, but unbalanced ferrite phase cannot be formed, so that ductility and hole expansibility of steel may be lowered, and there is a problem in that manufacturing cost is increased.

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

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

Niobium (Nb) is an element that segregates at austenite grain boundaries to inhibit coarsening of austenite grains during annealing heat treatment, and forms fine carbides, thereby contributing to improvement of strength.

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

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

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

Titanium (Ti) is an element forming fine carbides, contributing to ensuring yield strength and tensile strength. Further, ti has an effect of reducing the possibility of occurrence of cracks in continuous casting because it has an effect of suppressing formation of AlN due to Al inevitably present in steel by precipitating N in steel as TiN.

When the content of such Ti exceeds 0.1%, coarse carbides are precipitated, and there is a possibility that strength and elongation are reduced due to a reduction in the carbon content in the steel. Further, there is a possibility that nozzle clogging is caused at the time of continuous casting, and there is a problem that manufacturing cost increases.

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

Boron (B): less than 0.002% (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%, too much B is concentrated on the surface, and thus may cause deterioration of 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 grain size refining effect and deoxidization of steel, and 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, and thus the effect of improving strength is obtained, but excessive inclusions are formed at the time of steelmaking continuous casting operation, and thus the possibility of occurrence of surface defects of the plated steel sheet increases.

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

Phosphorus (P): less than 0.05 percent (except 0 percent)

Phosphorus (P) is a substitutional element having the greatest effect of solid solution strengthening, and is an element that improves in-plane anisotropy without significantly reducing formability and is advantageous in ensuring strength. However, when such P is added excessively, the possibility of occurrence of brittle fracture increases greatly, leading to an increase in the possibility of occurrence of plate breakage of a slab during hot rolling, and there is a problem of impeding the plating surface characteristics.

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

Sulfur (S): less than 0.01 percent (except 0 percent)

Sulfur (S) is an impurity element in steel and is an element inevitably added, and it is preferable to control the sulfur content to be as low as possible because ductility is hindered. In particular, S has a problem of increasing the possibility of occurrence of red hot shortness, and therefore, the sulfur content is preferably controlled to 0.01% or less. But 0% may be excluded in view of the level that is inevitably added during the manufacturing process.

Nitrogen (N): less than 0.01 percent (except 0 percent)

Nitrogen (N) is a solid solution strengthening element, but when the content of nitrogen (N) exceeds 0.01%, the risk of brittleness increases, and excessive AlN is precipitated by bonding with Al in steel, and thus the continuous casting quality may be hindered.

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

The rest of the invention is iron (Fe). However, it is possible that undesirable impurities are inevitably mixed from the raw materials or the surrounding environment in the conventional manufacturing process, and thus the magazine cannot be excluded. These impurities are well known to those skilled in the conventional manufacturing process and therefore are not specifically described in this specification in their entirety.

The microstructure in the steel sheet of the present invention having the above alloy composition may be composed of ferrite as a soft phase (soft phase), martensite as a hard phase, and an unbalanced ferrite phase formed at the interface thereof.

Specifically, the steel sheet of the present invention contains 20 to 30% by area of ferrite phase, 5 to 15% by area of unbalanced ferrite phase, and may contain martensite phase as a balance structure. In addition, a trace amount of the retained austenite phase may be contained.

In the present invention, the unbalanced ferrite phase is a structure that is advantageous in minimizing the difference in hardness between the soft phase and the hard phase, and is a structure different from the existing balanced ferrite (polygonal ferrite). The unbalanced ferrite may be acicular ferrite or bainitic ferrite. Furthermore, according to the cooling conditions, weissella ferrite (Widmanstatten ferrite), bulk ferrite (Massive ferrite), or the like may be included. Specifically, unbalanced ferrite is affected by components constituting a parent phase (moter phase), and contains relatively high C and Mn as compared to balanced ferrite. For example, in the case of balanced ferrite, it is assumed that the C concentration is 0.02%, and unbalanced ferrite has a higher C content of 0.03 to 0.04%.

Therefore, the C concentration and the Mn concentration of the hard phase formed in the vicinity (periphery) of the unbalanced ferrite are relatively reduced, and therefore the hardness difference between the soft phase and the hard phase is reduced while the hole expansibility can be improved. Further, when the Si concentration in the unbalanced ferrite is as low as less than 1%, the stacking fault energy (stacking fault energy) increases, and the cross slip becomes difficult and prevents the formation of holes (void) due to deformation (fig. 2).

When the fraction of such unbalanced ferrite phase is excessively high, the fraction of the hard phase is relatively reduced, and thus the strength of the target level cannot be ensured. In view of this, the unbalanced ferrite phase may be contained by 15% or less. On the other hand, when the fraction of the unbalanced ferrite phase is less than 5%, the above-described effect (minimizing the hardness difference between the hard phase and the soft phase) cannot be sufficiently obtained, and thus hole expansibility becomes poor.

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

The fraction of the martensite phase is not particularly limited in the structure other than the ferrite phase and the unbalanced ferrite phase, but may include a martensite phase having an area fraction of 50% or more in order to secure an ultra-high strength of 1100MPa or more. However, when the fraction of the martensite phase exceeds 75%, the ductility decreases, and thus it is difficult to secure the formability at the target level.

In addition, it is advantageous that the fraction is not more than 3% in terms of the retained austenite phase, and even if the fraction is 0%, there is no problem in ensuring desired physical properties.

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

Further, the steel sheet has a hole expansion ratio (Hole Expansion Ratio, HER) of 25% or more, and thus has an effect of excellent resistance to cracks that may occur at the time of processing and collision fracture resistance.

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

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

[ heating of billet ]

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

The process is performed in order to smoothly perform the subsequent hot rolling process and sufficiently obtain the desired physical properties of the steel sheet. In the present invention, the conditions of such a heating process are not particularly limited as long as they are ordinary conditions. As an example, the heating process may be performed at a temperature in the range of 1100-1300 ℃.

[ Hot Rolling ]

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

When the outlet side temperature at the time of hot finish rolling is lower than Ar3, the heat distortion resistance increases sharply, and the upper (top), lower (tail) and edge (edge) portions of the hot rolled coil become single-phase regions, so that in-plane anisotropy increases, and there is a possibility that formability is deteriorated. In addition, when the outlet side temperature at the time of the hot finish rolling exceeds 1000 ℃, the rolling load is relatively reduced to be advantageous for productivity, but there is a possibility that thick scale is generated.

More specifically, the finish hot rolling may be performed at a temperature ranging from 760 to 940 ℃.

Winding up

The hot rolled steel sheet manufactured as described above may be wound into a coil shape.

The winding may be performed at a temperature ranging from 400 to 700 ℃. When the winding temperature is lower than 400 ℃, excessive martensite phase or unbalanced ferrite phase is formed, and the strength of the hot rolled steel sheet is excessively increased, so that problems such as shape defects due to load may occur at the time of subsequent cold rolling. On the other hand, when the winding temperature exceeds 700 ℃, there is a problem that the acid-washing property becomes poor due to the increase of the surface scale.

[ Cooling ]

Preferably, the rolled hot rolled steel sheet is cooled to normal temperature at an average cooling rate of 0.1 ℃/sec(s) or less (except 0 ℃/sec). In this case, the rolled hot rolled steel sheet may be cooled after being transferred, placed, etc., and the process before cooling is not limited thereto.

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

[ Cold Rolling ]

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

In the present invention, the cold rolling may be performed at a cold rolling reduction of 55 to 70%. When the cold rolling reduction is less than 55%, the recrystallization driving force is weakened, which is difficult in obtaining good recrystallized grains, on the other hand, when the cold rolling reduction exceeds 70%, the risk of cracking at the edge portion of the steel sheet increases, and there is a possibility that the rolling load increases sharply.

In the present invention, it is possible to further promote the recrystallization of ferrite in the heating section in the subsequent continuous annealing process in a state where a suitable level of cold rolling reduction is applied at the time of cold rolling, thereby inducing the formation of fine ferrite, and thus small and uniform austenite can be formed in the ferrite grain boundary. This affects the size or distribution of the unbalanced structure during cooling and is advantageous in improving workability such as elongation, hole expansibility, etc. while maintaining the strength of the final product.

In addition, the cold rolling reduction may be achieved by only 1 cold rolling, that is, only 1 stand (stand), and the reduction of steel may be performed as described above, thus having an advantageous economical effect.

However, in the case of a thick steel material having a thickness of 6mm or more as a hot rolled steel sheet before cold rolling, the target rolling reduction can be achieved by repeating rolling using a reversing mill (reversing mill). In this case, the number of all passes (pass) of the repeated rolling may be set to 1 stand. A reversible rolling mill is a rolling mill for rolling a thin sheet steel, and refers to a rolling mill that reciprocates and rolls a material between a pair of rolls (rolls), and can set a single pass of the material at reciprocation to 1 pass.

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

[ continuous annealing ]

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

In general, a continuous annealing furnace (CAL) may be constituted by [ heating zone-soaking zone-cooling zone (slow cooling zone and quench zone) - (overaging zone as needed) ] and subjected to the following process: the cold-rolled steel sheet is charged into the continuous annealing furnace as described above, and then heated at a specific temperature in a heating zone, and is maintained for a prescribed time in a soaking zone after reaching a target temperature.

In the present invention, the temperatures of the heating zone and the soaking zone at the time of the continuous annealing may be controlled to be the same, which means that the termination 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 belt and the soaking belt can be controlled between 810 ℃ and 850 ℃.

When the temperature of the heating belt is lower than 810 ℃, sufficient heat input for recrystallization cannot be applied, on the other hand, when the temperature of the heating belt exceeds 850 ℃, productivity is lowered and excessive austenite phase is formed, so that the fraction of the hard phase after the subsequent cooling is greatly increased, and thus there is a possibility that ductility of the steel is deteriorated.

In addition, when the temperature of the soaking belt is lower than 810 ℃, excessive cooling is required at the termination temperature of the heating belt, thus being disadvantageous in economy, and the heat for recrystallization may be insufficient. On the other hand, when the temperature of the soaking zone exceeds 850 ℃, the fraction of austenite is excessively large, and the hard phase increases during cooling, so that formability may be lowered.

When the temperature of the soaking zone is increased in the above temperature range, the stability of austenite can be reduced, whereby the generation of unbalanced ferrite phase in the subsequent cooling process can be promoted.

Although specific description will be given later, in the present invention, the cooling is performed in stages while the cooling is performed after passing through the heating belt and the soaking belt, the unbalanced ferrite phase is introduced after the primary cooling, and the final structure may be composed of a certain fraction of the soft phase, the hard phase, and the unbalanced ferrite phase. Therefore, in the steel sheet of the present invention, not only the strength and ductility can be improved, but also the effect of improving workability can be obtained simultaneously due to the interfacial strengthening of the unbalanced ferrite phase.

Therefore, in order to obtain the desired microstructure of the present invention, it is preferable to control the heat input applied to the steel sheet in the heating section constituted by the heating zone and the soaking zone at the time of the continuous annealing.

[ staged Cooling ]

As described above, the cold-rolled steel sheet subjected to the heat treatment as described above is cooled to form a desired structure, and in this case, it is preferable to perform stepwise cooling.

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

In this case, the primary cooling is performed more slowly than the secondary cooling, and thus plate shape defects due to a rapid temperature drop during the secondary cooling, which is a subsequent relative quenching section, can be suppressed.

When the termination temperature at the time of the primary cooling is lower than 570 ℃, the diffusion activity of carbon is low due to an excessively low temperature, and the carbon concentration in ferrite increases, while on the other hand, the fraction of the hard phase becomes excessively large due to a low carbon concentration in austenite, and the yield ratio increases, whereby the tendency of occurrence of cracks during processing increases. Further, since the cooling rate between the soaking belt and the cooling belt (slow cooling belt) is too high, a problem occurs in that the shape of the plate becomes uneven. When the termination temperature exceeds 630 ℃, an excessively high cooling rate is required at the time of subsequent cooling (secondary cooling), and it is difficult to introduce an unbalanced ferrite phase.

Further, when the average cooling rate at the time of the primary cooling exceeds 10 ℃/sec, the diffusion of carbon cannot be sufficiently performed. In addition, the primary cooling may be performed at an average cooling rate of 1 ℃/sec or more in view of productivity.

As described above, after the primary cooling is completed, the quenching (secondary cooling) may be performed at a cooling rate equal to or higher than a predetermined level. At this time, when the secondary cooling termination temperature is lower than 300 ℃, a cooling deviation occurs in the width direction and the length direction of the steel sheet, and there is a possibility that the shape of the sheet becomes poor. On the other hand, when the secondary cooling termination temperature exceeds 400 ℃, the hard phase cannot be sufficiently ensured, and the strength may be lowered, and an increase in yield ratio and a decrease in elongation may be induced due to the formation of bainite.

Further, when the average cooling rate at the time of the secondary cooling is less than 5 ℃/sec, the fraction of the hard phase may be excessively high, but when the average cooling rate at the time of the secondary cooling exceeds 50 ℃/sec, the hard phase may become insufficient instead.

In addition, if necessary, after the completion of the staged cooling, an overaging treatment may be performed.

The overaging treatment is a process of maintaining a certain time after the end temperature of the secondary cooling, and uniformly heat-treating along the width direction and the length direction of the coil, thereby having the effect of improving the shape quality. For this purpose, the overaging treatment may be carried out for 200-800 seconds.

The overaging treatment may be performed after the end of the secondary cooling, and thus the temperature of the overaging treatment may be the same as the end temperature of the secondary cooling, or the overaging treatment may be performed within the end temperature range of the secondary cooling.

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, and particularly, the recrystallization of ferrite is maximized by an optimized cold rolling and annealing process, so that it is possible to have a structure in which martensite as the hard phase is uniformly distributed on the finally recrystallized ferrite matrix. In addition, by introducing an unbalanced ferrite phase at the interface between the hard phase and the soft phase, the effect of improving the crack resistance during processing is obtained.

Therefore, even if the steel sheet of the present invention has a high tensile strength of 1100MPa or more, excellent formability such as hole expansibility can be ensured by ensuring a low yield ratio and high ductility.

Hereinafter, the present invention will be described in more detail with reference to examples. However, these examples are merely illustrative of the practice of the present invention, and the present invention is not limited to these examples. This is because the scope of the invention is determined by what is recited in the claims and what is reasonably derived therefrom.

Detailed Description

Example (example)

Billets having the alloy compositions shown in table 1 below were manufactured, and then each billet was heated at 1200 ℃ for 1 hour, and then subjected to finish hot rolling at a finish rolling temperature of 880 to 920 ℃ to manufacture a hot rolled steel sheet. Thereafter, each of the hot rolled steel sheets was wound at 650 ℃ and then cooled to normal temperature at a cooling rate of 0.1 ℃/sec. Thereafter, the rolled hot rolled steel sheet was subjected to cold rolling and continuous annealing treatment according to the conditions shown in the following table 2, and then subjected to staged cooling (primary cooling-secondary cooling), and then to overaging treatment at 360 ℃ for 520 seconds, thereby manufacturing a final steel sheet.

At this time, the primary cooling in the staged cooling was performed at an average cooling rate of 3 ℃/sec, and the secondary cooling was performed at an average cooling rate of 20 ℃/sec. Furthermore, cold rolling is performed by 1 stand.

The microstructure of each steel sheet manufactured as described above was observed, and physical property indexes used in processing such as elongation and processing characteristics and hole expansion ratio were evaluated, and the results are shown in table 3 below.

In this case, the tensile test was performed for each test piece by collecting a tensile test piece of JIS No. 5 size in a direction perpendicular to the rolling direction and then performing the tensile test at a strain rate (strain rate) of 0.01/sec.

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

In addition, the martensite phase and the unbalanced ferrite phase corresponding to the hard phase in the structure phase (phase) were observed by SEM at 2000 and 5000 times magnification after etching with nitric acid etching solution (nital). At this time, the observed size, fraction, etc. of each phase were measured. In addition, for phase (phase), the fractions were measured after nitric acid etching using SEM and image analysis program (Image analyzer program).

TABLE 1

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TABLE 2

TABLE 3

As shown in tables 1 to 3, it is known that the steel alloy composition and manufacturing conditions satisfy the conditions proposed by the present invention, particularly, the cold rolling and continuous annealing processes satisfy the invention examples 1 to 6 of the processes proposed by the present invention, achieve sufficient recrystallization of ferrite and formation of fine hard phases during annealing after cold rolling, and are connected by unbalanced ferrite structure at the interface, thereby having high strength while having appropriate yield strength for plate shape processing, and having excellent elongation. Further, since the hole expansibility was excellent, it was confirmed that the moldability at the target level was ensured.

On the other hand, in comparative examples 1 to 6, in which the soaking temperature at the time of continuous annealing in the manufacturing process of the steel sheet was low, recrystallization did not sufficiently occur, and the proper fraction of austenite formed in the soaking zone had high stability, so unbalanced ferrite was not sufficiently introduced during cooling. As a result, the ductility and/or hole expansibility were poor.

In addition, although comparative examples 7 to 10 were heated at an appropriate temperature during continuous annealing, since the termination temperature at the time of primary cooling was high, ductility and/or hole expansibility were poor as the time for introducing unbalanced ferrite during cooling was insufficient.

In comparative examples 11 to 14, which contain too much Cr as a hardenability element, there is a risk of cracking during processing due to too high yield strength, and unbalanced ferrite phase cannot be introduced due to low temperature of the soaking zone, so that the result of poor ductility is exhibited in some of the comparative examples.

Fig. 3 shows the microstructure photographs of comparative examples 4 to 7 and inventive example 1.

As shown in fig. 3, in invention example 1, a uniform and fine unbalanced ferrite phase was introduced into a recrystallized ferrite matrix of a sufficient fraction during primary cooling, and a martensite phase of a certain fraction was formed during secondary cooling.

On the other hand, in comparative examples 4 to 7, it was confirmed that a small amount of unbalanced ferrite was introduced due to the conditions of the temperature of the soaking zone at the time of continuous annealing or the primary cooling termination temperature. Among them, in comparative example 4 in which the soaking zone temperature was not 800 ℃ and the primary cooling termination temperature was quite high and comparative example 7 in which the primary cooling termination temperature was quite high, it was found that the unbalanced ferrite was 1% or less, and almost no observation was made.

Claims

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.