JP7789205B2 - High-yield-ratio high-strength steel with excellent impact resistance after cold forming and its manufacturing method - Google Patents

Description

The present invention relates to high-strength steel and a manufacturing method thereof, and more specifically to high-strength steel having excellent impact resistance after cold forming and a high yield ratio, and a manufacturing method thereof.

Due to vehicle characteristics, conventional commercial vehicle chassis parts are typically made of steel with a thickness of 8 mm or more. High-strength hot-rolled steel sheets with a tensile strength of 600 MPa are used for members, and hot-rolled steel sheets with a tensile strength of 440 MPa are used for wheel discs. However, in recent years, in pursuit of weight reduction, high-strength steel sheets with a tensile strength of 700 MPa or more are being used for members and 590 MPa or more for wheel discs, with a trend toward reducing steel thickness and changing part designs. Additionally, while wheels were traditionally manufactured using a press forming process, they are now increasingly being manufactured using spinning and flow forming. Because these forming processes subject hot-rolled steel sheets to greater deformation, they require hot-rolled steel sheets with higher elongation rates, and the formed parts must ensure durability and impact resistance during use.

However, when conventional high-strength steel is used in the spinning and flow forming processes described above, the durability of the resulting parts must be equal to or greater than conventional steels. However, this can be difficult to achieve due to the occurrence of fine cracks on shear surfaces and other areas where the forming volume is high, resulting in poor durability.

Conventional high-strength steels, as disclosed in Patent Documents 1 and 2, undergo normal hot rolling in the austenite region, followed by coiling at high temperatures to form a ferrite matrix and fine precipitates. Patent Document 3 also employs a coiling technique in which the steel is cooled to a temperature at which a bainite matrix structure is formed, preventing the formation of coarse pearlite. Patent Document 4 also proposes a technique for refining austenite grains by using Ti, Nb, etc. to reduce the roll width by a large reduction of 40% or more in the unrecrystallized region during hot rolling. More recently, Patent Document 5 proposes a technique for improving the uniformity of the microstructure between the surface and deeper layers of the steel plate and suppressing the formation of coarse carbides. Patent Document 6 proposes a technique for simultaneously suppressing the formation of pearlite, MA (Martensite and Austenite) phases, and martensite, which adversely affect durability.

However, Patent Documents 1 to 4 fail to consider the occurrence of cracks on and around the sheared surfaces during shear forming of high-strength thick steel sheets, and are configured with cooling rate and large reduction conditions that are difficult to achieve during the production of thick steel sheets with a thickness of 8 mm or more. When using precipitate-forming elements such as Ti, Nb, and V to refine the grain size of thick steel sheets while simultaneously ensuring strength, coiling at high temperatures of 500 to 700°C, where precipitates are likely to form, can lead to excessive ferrite growth, reduced yield strength, and the formation of coarse pearlite. Furthermore, even when manufacturing at a low coiling temperature to utilize a bainite matrix structure, if the cooling rate in the width direction of the steel sheet during cooling after hot rolling is not uniformly controlled, a non-uniform microstructure will form in the hot-rolled steel sheet, making it difficult to achieve high elongation and stable high yield ratio properties. This also increases sensitivity to forming cracks, such as cracks on the sheared surfaces, during forming. Furthermore, since this technique is intended for hot-rolled steel sheets with a thickness of less than 5 mm used in passenger cars, the required cooling rate is too high and is therefore unsuitable for manufacturing thick steel sheets. Furthermore, applying a large reduction of 40% in the unrecrystallized region can deteriorate the shape quality of the rolled sheet and place a strain on the equipment, making it difficult to apply to thick steel sheets with a thickness of 8 mm or more. Patent Documents 5 and 6 are inventions aimed at thick steel sheets. First, Patent Document 5 proposes a manufacturing technique for thick high-strength steel sheets, in which the crystal grain shape in the deep thickness portion (1/4t to 1/2t) is equiaxed and has fine crystal grains, thereby suppressing the formation of MA phases and martensite, in order to improve the durability of the thick high-strength steel sheets. Patent Document 6 proposes a manufacturing technique in which a hot-rolled coil is divided into three equal parts in the length direction using a relational expression derived for specific components, and the head, mid, and tail portions are cooled to different cooling end temperatures at constant cooling rates and then coiled. This technology produces a uniform microstructure by controlling the cooling rate after hot rolling through a relationship derived for specific components, taking into account the quality of the shear surface of the part. While this is effective in improving the durability of commercial vehicle wheels, which contain numerous punched holes and are subject to constant load, it does not take into account impact resistance after forming. Furthermore, it is difficult to uniformly control the cooling of the steel sheet after hot rolling across its entire width, and when the hot-rolled steel sheet is thick, with a thickness of 8 mm or more, it is difficult to control the actual cooling rate.

The cooling process after hot rolling is typically carried out within a few tens of seconds from a 100-120m long ROT (Run-Out Table), but it is difficult to produce hot-rolled steel sheets that are cooled to the cooling end temperature or coiling temperature while the cooling rate deep within the thickness range satisfies the proposed range. Therefore, with the above-mentioned conventional technology, it is difficult to achieve the effect of suppressing the formation of coarse carbides in hot-rolled steel sheets with a thickness of 8mm or more, and there are problems with not being able to ensure a high level of impact resistance.

Korean Patent Publication No. 10-2010-0029138 (Published on March 15, 2010) Japanese Patent Publication No. 2007-262487 (published October 11, 2007) Korean Patent Publication No. 10-1528084 (Published June 10, 2015) Japanese Patent Publication No. Hei 9-143570 (published June 3, 1997) Korean Patent Publication No. 10-2020-0062422 (Published on June 4, 2020) Korean Patent Publication No. 10-2021-0068808 (Published on June 10, 2021)

One aspect of the present invention is to provide a high-strength steel that has excellent impact resistance after cold forming and a high yield ratio, and a method for manufacturing the same.

The objectives of the present invention are not limited to the above. A person of ordinary skill in the art would have no difficulty in understanding the further objectives of the present invention from the overall content of this specification.

One aspect of the present invention is a steel sheet containing, by weight%, C: 0.05 to 0.15%, Si: 0.01 to 0.5%, Mn: 1.0 to 2.0%, Al: 0.01 to 0.1%, Cr: 0.001 to 1.0%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, Ti: 0.03 to 0.08%, Nb: 0.01 to 0.05%, the balance being Fe and other inevitable impurities, with the sum of Nb and Ti being 0.04 to 0.1%;
The microstructure of the surface layer portion within a thickness range of 50 μm from the surface contains, in area%, 95% or more of equiaxed ferrite and 3% or less of pearlite, and contains 5% or less in total of one or more of bainitic ferrite, bainite, MA (Martensite-Austenite constituent) phase, and martensite,
It is possible to provide a steel sheet in which the microstructure in the center portion within the range of ¼ to ¾ of the thickness contains, in area %, 80 to 95% bainitic ferrite, 10% or less bainite, 3% or less pearlite, 5 to 10% in total of one or two of an MA (Martensite-Austenite constituent) phase and martensite, and the remainder contains equiaxed ferrite.

The steel plate can be 8 to 25 mm thick.

The above steel plate has a tensile strength of 590 MPa or more, a fracture elongation of 25% or more, a yield ratio of 0.75 to 0.9, and an impact toughness of 70 J or more at -20°C after cold forming.

The above steel plate may have a ratio (E/YS) of impact toughness (E) at -20°C after cold forming to yield strength (YS) before cold forming of 0.15 or more.

The steel plate includes edge portions corresponding to 30% of the width at both ends and a central portion of the central 40% region excluding both edge portions, and the difference in tensile strength between the edge portions and the central portion is 10 MPa or less, the difference in fracture elongation is 8% or less, and the difference in impact toughness at -20°C after cold forming is 20 J or less.

Another aspect of the present invention is a method for producing a steel slab containing, by weight, 0.05-0.15% C, 0.01-0.5% Si, 1.0-2.0% Mn, 0.01-0.1% Al, 0.001-1.0% Cr, 0.001-0.05% P, 0.001-0.01% S, 0.001-0.01% N, 0.03-0.08% Ti, 0.01-0.05% Nb, the balance being Fe and other unavoidable impurities, the total of which is 0.04-0.1% Nb and Ti;
hot-rolling the reheated steel slab; and cooling the hot-rolled steel sheet to a temperature range of 500 to 650°C at an average cooling rate of 1 to 30°C/s, which is equal to or greater than the CR value defined by the following Relation 1, and coiling the steel sheet,
In the cooling and winding step, the edge portions corresponding to 30% of the area at both ends in the coil width direction are cooled at a temperature of 550 to 650 ° C. (TE), and the central portion of the central 40% area corresponding to the area excluding both edge portions in the width direction is cooled at a temperature of 500 to 550 ° C. (TC).
A method for manufacturing a steel plate in which the average temperature difference between the edge portion and the center portion is 50 to 150°C can be provided.
[Relationship 1]
CR=45-16.3×[C]-5.6×[Si]-16.3×[Mn]-2.9×[Cr]+15×[Ti]+23×[Nb]-0.9×(t-8)
(Here, [C], [Si], [Mn], [Cr], [Ti] and [Nb] are the weight percentages of each element, and t is the thickness of the steel plate (mm).)

The reheating temperature is 1100 to 1350°C,
The hot rolling temperature may be 800 to 1150°C.

The method may further include air-cooling the wound coil at a temperature range of 200°C or less.

The steel plate can be 8 to 25 mm thick.

One aspect of the present invention provides high-strength steel that has excellent impact resistance after cold forming and a high yield ratio, as well as a manufacturing method thereof.

One aspect of the present invention provides high-strength steel that can be used for chassis members and wheels of medium- to large-sized commercial vehicles, and a manufacturing method thereof.

1 is a photograph showing the microstructure of Example 2 according to the present invention, observed with a scanning electron microscope (×3,000). 1 is a photograph showing the microstructure of Comparative Example 2 according to an embodiment of the present invention, observed with a scanning electron microscope (×3,000).

The following describes preferred embodiments of the present invention. The embodiments of the present invention can be modified in various ways, and the scope of the present invention should not be construed as being limited to the embodiments described below. The present embodiments are provided to further explain the present invention to those skilled in the art to which the invention pertains.

In order to solve the above-mentioned conventional problems and ensure excellent formability and impact resistance, the inventors studied the changes in impact resistance after cold forming depending on the microstructural characteristics of the steel sheet. As a result, they confirmed that the desired physical properties can be achieved by optimizing the alloy composition and manufacturing conditions and controlling the microstructural characteristics according to the thickness and width directions of the steel sheet, leading to the completion of the present invention.

In hot-rolled steel sheets typically manufactured in coil form, coarse carbides and pearlite are likely to form when maintained at high temperatures of approximately 500-700°C for long periods of time. In particular, if the ferrite phase transformation that begins during the cooling process after hot rolling proceeds slowly, the untransformed phases tend to contain a high amount of carbon, creating conditions favorable for the formation of coarse carbides and pearlite. Furthermore, the cooling rate is slower in the center of the coil than in the edge regions, further promoting the development of these structures. Therefore, to prevent the formation of coarse carbides and pearlite in the center of the coil, the wound coil must be cooled to room temperature by forced cooling, such as by water cooling. However, this is undesirable because the edge regions, where the cooling rate is faster, experience excessive formation of martensite and MA (Martensite and Austenite) phases, resulting in a non-uniform microstructure. This makes it difficult to achieve high elongation and increases cracking on the shear surface. Therefore, this invention proposes a method that can suppress the formation of coarse carbides and pearlite without forcibly cooling the coil.

The present invention is described in detail below.

The steel composition of the present invention is described in detail below.

Unless otherwise specified, the percentages representing the content of each element in this invention are based on weight.

The steel according to one aspect of the present invention contains, by weight, C: 0.05-0.15%, Si: 0.01-0.5%, Mn: 1.0-2.0%, Al: 0.01-0.1%, Cr: 0.001-1.0%, P: 0.001-0.05%, S: 0.001-0.01%, N: 0.001-0.01%, Ti: 0.03-0.08%, Nb: 0.01-0.05%, the balance being Fe and other unavoidable impurities, and the sum of Ti and Nb can be 0.04-0.1%.

Carbon (C): 0.05-0.15%
Carbon (C) is the most economical and effective element for strengthening steel, and increasing its content increases the precipitation strengthening effect or the bainite phase fraction, making it easier to ensure strength. However, as the thickness of a hot-rolled steel sheet increases, the cooling rate at the center of the thickness during cooling after hot rolling slows, and a high carbon (C) content can easily lead to the formation of coarse carbides and pearlite. Therefore, if the carbon (C) content is less than 0.05%, it is difficult to achieve sufficient strengthening effects. If the carbon (C) content exceeds 0.15%, there are problems such as reduced impact resistance due to the formation of pearlite and coarse carbides at the center of the thickness, and weldability may also be deteriorated. The preferred lower limit is 0.055%, and the more preferred upper limit is 0.12%.

Silicon (Si): 0.01 to 0.5%
Silicon (Si) is an element that is effective in deoxidizing molten steel and solid-solution strengthening steel, and also in delaying the formation of coarse carbides to improve formability. However, if its content is less than 0.01%, the effects of solid-solution strengthening and delaying carbide formation cannot be maximized, while if its content exceeds 0.5%, red scale is formed on the steel sheet surface during hot rolling, resulting in significantly inferior quality, as well as reduced ductility and weldability. Preferably, Si may be contained in an amount of 0.05% or more, and more preferably 0.3% or less.

Manganese (Mn): 1.0 to 2.0%
Manganese (Mn), like Si, is an effective element for solid-solution strengthening of steel, increasing the hardening ability of steel and facilitating the formation of bainite during cooling after hot rolling. However, if the Mn content is less than 1.0%, the above-mentioned effect of addition cannot be obtained, and if the Mn content exceeds 2.0%, the hardening ability increases significantly, making martensitic phase transformation more likely to occur, and there is a risk of large segregation developing from the thickness center during slab casting in the continuous casting process. Preferably, Mn can be contained in an amount of 1.3% or more, and more preferably 1.8% or less.

Aluminum (Al): 0.01 to 0.1%
Aluminum (Al) is a component added primarily for deoxidation, and if its content is less than 0.01%, its effect may be insufficient. On the other hand, if its content exceeds 0.1%, it combines with nitrogen to form AlN, which is likely to cause corner cracks in the slab during continuous casting and defects due to the formation of inclusions. Aluminum may be added preferably at 0.015% or more, more preferably at 0.06% or less.

Chromium (Cr): 0.001 to 1.0%
Chromium (Cr), like Mn, strengthens steel through solid solution and retards the ferrite phase transformation during cooling, promoting the formation of bainite. However, if its content is less than 0.001%, the above effects cannot be achieved. If its content exceeds 1.0%, the ferrite phase transformation may be excessively retarded, resulting in excessive martensite formation and reduced elongation. Furthermore, excessive Cr addition can lead to the development of large segregations in the center of the steel, resulting in uneven microstructures across the thickness and reduced impact resistance. The preferred lower limit is 0.01%, and the more preferred upper limit is 0.5%.

Phosphorus (P): 0.001 to 0.05%
Like Si, phosphorus (P) simultaneously strengthens the solid solution and promotes ferrite phase transformation. However, manufacturing with a phosphorus (P) content of less than 0.001% is economically disadvantageous due to the high manufacturing costs and is insufficient to obtain sufficient strength. On the other hand, if the P content exceeds 0.05%, embrittlement occurs due to grain boundary segregation, which easily causes microcracks during molding and significantly reduces impact resistance.

Sulfur (S): 0.001 to 0.01%
S is an impurity present in steel, and if its content exceeds 0.01%, it combines with Mn and other elements to form non-metallic inclusions, which can cause microcracks during forming and significantly reduce impact resistance. However, reducing the content to less than 0.001% requires a long time during steelmaking operations, resulting in poor productivity.

Nitrogen (N): 0.001 to 0.01%
Nitrogen (N), along with carbon, is a representative solid solution strengthening element and forms coarse precipitates with titanium, aluminum, and other elements. While nitrogen (N) generally has a stronger solid solution strengthening effect than carbon, the toughness of steel decreases significantly as the nitrogen (N) content increases. Therefore, the upper limit of nitrogen (N) content can be limited to 0.01%. However, reducing the nitrogen content to less than 0.001% can result in a longer steelmaking operation time and reduced productivity.

Titanium (Ti): 0.03 to 0.08%
Titanium (Ti) is a typical precipitation strengthening element and forms coarse TiN in steel due to its strong affinity with N. TiN has the effect of suppressing grain growth during the heating process for hot rolling. Furthermore, TiC precipitates, formed when the remaining titanium (Ti) reacts with N and dissolves in the steel and combines with C, are a useful component for improving the strength of steel. If the titanium (Ti) content is less than 0.03%, the above effect cannot be achieved. However, if the titanium (Ti) content exceeds 0.08%, the generation of coarse TiN and the coarsening of the precipitates can lead to problems such as degraded impact resistance during forming. Preferably, the titanium content is 0.04% or more, and more preferably 0.075% or less.

Niobium (Nb): 0.01 to 0.05%
Niobium (Nb), along with Ti, is a representative precipitation strengthening element that precipitates during hot rolling and effectively improves the strength and impact toughness of steel by refining grains through delayed recrystallization. If the niobium (Nb) content is less than 0.01%, the above effect cannot be obtained, while if the Nb content exceeds 0.05%, excessive delayed recrystallization during hot rolling can lead to the formation of elongated grains and coarse complex precipitates, deteriorating formability. The preferred lower limit is 0.015%, and the more preferred upper limit is 0.04%.

In addition to the above-mentioned composition, the steel of the present invention may contain the remaining iron (Fe) and unavoidable impurities. Unavoidable impurities cannot be excluded because they may be unintentionally mixed in during normal manufacturing processes. Since such impurities are known to any engineer in the field of normal steel manufacturing, their full content will not be specifically mentioned in this specification.

The steel according to one aspect of the present invention may have a total of niobium (Nb) and titanium (Ti) of 0.04 to 0.1%.

Niobium (Nb) and titanium (Ti) precipitate as (Ti, Nb) (C, N) composite precipitates, and when they precipitate during hot rolling, they significantly enhance the grain refinement effect by delaying recrystallization. However, excessive formation of these composite precipitates results in an increase in coarse composite precipitates, resulting in little strength improvement and poor formability. If the total content of niobium (Nb) and titanium (Ti) is less than 0.04%, the grain refinement and strength improvement effects may be limited. On the other hand, if the total content exceeds 0.1%, formability will deteriorate and this is economically disadvantageous. The preferred lower limit is 0.045%, and the more preferred upper limit is 0.09%.

The steel microstructure of the present invention is described in detail below.

Unless otherwise specified in this invention, the percentages indicating the fraction of microstructures are based on area.

In one aspect of the steel of the present invention, the microstructure of the surface layer within a 50 μm thick range from the surface contains, by area percentage, 95% or more equiaxed ferrite, 3% or less pearlite, and a total of 5% or less of one or more of bainitic ferrite, bainite, MA (Martensite-Austenite constituent) phase, and martensite; the microstructure of the center portion within a 1/4 to 3/4 thickness range contains, by area percentage, 80 to 95% bainitic ferrite, 10% or less bainite, 3% or less pearlite, and a total of 5 to 10% of one or two of the MA (Martensite-Austenite constituent) phase and martensite, with the remainder being equiaxed ferrite.

In the present invention, if the surface layer contains less than 95% equiaxed ferrite, the ductility may be insufficient during spinning and flow forming, which are used in the manufacture of commercial vehicle wheels. This may result in severe work hardening in the surface layer, leading to the risk of fine cracks occurring during forming. In particular, if more than 3% of highly brittle pearlite is formed, or if more than 5% of one or more of the highly hard bainitic ferrite, bainite, MA phase, and martensite is included, cracks may easily propagate along the interface with the base phase. Therefore, to suppress the formation of fine cracks in the surface layer during forming and prevent crack propagation, it is preferable to include a total of 5% or less of one or more of bainitic ferrite, bainite, MA phase, and martensite. In the present invention, the microstructure in the surface layer may be 100% equiaxed ferrite, and the total of pearlite, bainitic ferrite, bainite, MA phase, and martensite may be 0%.

Furthermore, if the bainitic ferrite content in the center is less than 80%, cracks easily occur on the sheared surfaces during punching and shear forming during wheel manufacturing, resulting in poor impact resistance after forming. Furthermore, during the cooling process of a steel plate after hot rolling, bainitic ferrite, the matrix structure, is formed in the center of the thickness of the rolled plate. A high concentration of residual carbon remains in the untransformed austenite, facilitating the formation of pearlite. However, if the pearlite content exceeds 3%, severe cracking occurs on the sheared surfaces during forming, resulting in poor impact resistance after forming. When the pearlite fraction is 3% or less, cracks do not occur during forming, such as shearing, and excellent impact resistance at low temperatures can be achieved. In the present invention, pearlite can contain carbides and nitrides with a diameter of 1 μm or more.

In contrast, containing 5-10% MA phase or martensite can be advantageous for ensuring impact resistance and high strength after forming without affecting crack generation. The MA phase is advantageous for creating potential density around it to ensure high strength, and when formed together with a matrix structure composed of ferrite and bainite, it can provide excellent impact resistance even when the potential density increases after cold forming. However, if the MA phase or martensite is less than 5%, the yield strength and tensile strength will be insufficient, and if it is contained in excess of 10%, there will be problems with insufficient ductility and poor formability. Bainite, too, may also be problematic in that it lacks ductility if it exceeds 10%.

In the present invention, the microstructure in the center can contain 0% bainite and 0% pearlite, and in addition to bainitic ferrite, bainite, pearlite, MA phase, and martensite, equiaxed ferrite can be unavoidably included.

In the present invention, the area fraction of the microstructure can be analyzed using an optical microscope and a scanning electron microscope (SEM), and the area fraction of the phases can be measured from an image observed at 3,000x magnification at a position corresponding to the center of the thickness of the rolled cross section.

The steel manufacturing method of the present invention is described in detail below.

Steel according to one aspect of the present invention can be produced by reheating, rolling, cooling, and coiling a steel slab that satisfies the alloy composition described above.

Reheating Steel slabs meeting the alloy composition of the present invention can be reheated at temperatures ranging from 1100 to 1350°C.

If the reheating temperature is below 1100°C, the precipitates will not be fully redissolved, which may result in reduced precipitate formation in processes after hot rolling, leaving coarse TiN remaining. On the other hand, if the temperature exceeds 1350°C, abnormal grain growth of austenite grains may result in a decrease in strength.

Hot Rolling The reheated steel slab can be hot rolled in the temperature range of 800 to 1150°C.

If the hot rolling temperature exceeds 1150°C, the temperature of the hot-rolled steel sheet will become too high, the grain size will become coarse, and the surface quality of the hot-rolled steel sheet may deteriorate. On the other hand, if the temperature is below 800°C, recrystallization will be delayed, causing elongated grains to develop, which may result in severe anisotropy and poor formability. Rolling at a temperature below the austenite temperature range may further worsen the development of a non-uniform microstructure.

Cooling and Coiling The hot-rolled steel sheet may be cooled to a temperature range of 500 to 650°C at an average cooling rate within a range of 1 to 30°C/s, which is equal to or greater than the CR value defined by the following Relation 1, and then coiled. During the cooling, the edge portions of 30% of the coil at both ends in the width direction may be cooled to a temperature of 550 to 650°C (TE), and the central portion of the central 40% of the coil excluding both edges in the width direction may be cooled to a temperature of 500 to 550°C (TC). At this time, the average temperature difference between the edge portions and the central portion may be 50 to 150°C.

In the present invention, Relational Formula 1 was derived to induce an appropriate level of ferrite phase transformation during cooling of a steel sheet, form a fine and uniform MA phase, and suppress excessive pearlite formation. If the cooling rate is less than the CR value of Relational Formula 1, ferrite in the center of the thickness becomes coarse and pearlite is excessively formed, resulting in severe cracking on the shear surface and deterioration of impact resistance after forming. In addition, if the cooling rate exceeds 30°C/s, bainite, MA phase, and martensite are excessively formed, resulting in insufficient ductility and poor shear surface quality.
[Relationship 1]
CR=45-16.3×[C]-5.6×[Si]-16.3×[Mn]-2.9×[Cr]+15×[Ti]+23×[Nb]-0.9×(t-8)
(Here, [C], [Si], [Mn], [Cr], [Ti] and [Nb] are the weight percentages of each element, and t is the thickness of the steel plate (mm).)

In order to prevent excessive carbide and pearlite formation, the cooling temperature after hot rolling must be below the end of cooling temperature. However, excessive bainite formation can result in a reduction in ferrite or excessive formation of MA phases and martensite, making it difficult to achieve the target elongation.

Therefore, in the present invention, in order to increase the cooling rate at the center of the coil width direction and reduce the time the coil is maintained at high temperature after coiling, the cooling end temperatures at the center and edge portions of the coil width direction can be set differently during cooling after hot rolling. However, in this case, the average temperature difference between the edge and center portions can be 50 to 150°C. If the average temperature difference is less than 50°C, it may be difficult to achieve the above effect. On the other hand, if the temperature exceeds 150°C, the above effect will not be further increased, but it may be difficult to control the temperature of each section of the coil.

The present invention does not particularly limit the method for controlling the end cooling temperatures of the edge and center during coiling. As an example, when cooling a hot-rolled steel sheet, a method can be applied in which the cooling water injected into the edge is stopped before it reaches the steel sheet, or the amount of cooling water injected is adjusted by varying the amount. Alternatively, the two methods can be used in parallel.

In the present invention, in order to ensure the desired strength, formability, and impact resistance, it is preferable to satisfy all of the conditions for the above relational expression 1 and the cooling end temperature. When all of the above cooling conditions are satisfied, the center in the thickness direction has a uniform and fine microstructure with bainitic ferrite as the base structure, and coarse carbides and pearlite are reduced in the inner winding part of the coil and the center in the thickness direction, where the cooling rate is slow, thereby eliminating the non-uniform structure of the hot-rolled steel sheet. Furthermore, the non-uniform formation of MA phases and the formation of coarse martensite can be suppressed in the outer winding part and edge part of the coil, where the cooling rate is relatively fast.

Cooling The wound coil can be air cooled to a temperature range up to 200°C.

In the present invention, the wound coil can be air-cooled at a temperature range of 200°C or less. Air-cooling the coil means cooling it in air at room temperature at a cooling rate of 0.001 to 10°C/h. If the cooling rate exceeds 10°C/h, some untransformed phases in the steel at the outer winding portion of the coil are likely to transform into the MA phase, which may result in poor shear formability, punching formability, and durability of the steel. On the other hand, controlling the cooling rate to less than 0.001°C/h requires separate heating and heat retention equipment, which can be economically disadvantageous. A preferred lower limit is 0.01°C/h, and a more preferred upper limit is 1°C/h.

The steel of the present invention manufactured in this manner has a thickness of 8 to 25 mm, a tensile strength of 590 MPa or more, a fracture elongation of 25% or more, a yield ratio of 0.75 to 0.9, an impact toughness at -20°C after cold forming of 70 J or more, and a ratio of the impact toughness after cold forming to the yield strength before cold forming of 0.15 or more, thereby providing excellent impact toughness properties while maintaining a high yield ratio. Furthermore, the steel includes edge portions corresponding to 30% of the width at both ends and a central portion of the central 40% region corresponding to the region excluding both edge portions, and the difference in tensile strength between the edge portions and the central portion may be 10 MPa or less, the difference in fracture elongation of 8% or less, and the difference in impact toughness at -20°C after cold forming of 20 J or less.

The present invention will now be described in more detail through examples. However, it should be noted that the following examples are intended to illustrate and explain the present invention in more detail, and are not intended to limit the scope of the present invention.

(Example)
Hot-rolled steel sheets were manufactured from steel slabs having the alloy compositions shown in Table 1 below under the conditions shown in Table 2 below. The steel slabs were reheated at temperatures of 1100 to 1350°C and then hot-rolled. Table 2 shows the cooling rates applied during manufacturing and the CR values of Relation 1. The cooling end temperatures are the temperature (TC) within 40% of the width at the center and the temperature (TE) within 30% of the width at both ends. The difference between the temperature at the center and the temperature at the edge is also shown.

[Relational expression 1] CR=45-16.3×[C]-5.6×[Si]-16.3×[Mn]-2.9×[Cr]+15×[Ti]+23×[Nb]-0.9×(t-8)
(Here, [C], [Si], [Mn], [Cr], [Ti] and [Nb] are the weight percentages of each element, and t is the thickness of the steel plate (mm).)

Table 3 below lists the microstructure measurements of the manufactured steel sheets. The microstructure was measured at the surface and center in the thickness direction, and the fractions were also measured at the edge and center in the width direction. For the surface, the microstructure was observed from the surface to a thickness of 50 μm, while for the center, the microstructure was observed from the surface to 1/4 to 3/4 of the thickness (25-75% interval, where t is thickness (mm)) in the thickness direction. For the edge, the microstructure was observed in the 30% region at both ends in the width direction, while for the center, the microstructure was observed based on the central 40% region excluding the edge. The area fractions of the MA phase and martensite were measured using an optical microscope and an image analyzer at 1,000x magnification after etching using the Lepera etching method. Additionally, the fractions of equiaxed ferrite (PF), bainitic ferrite (BF), bainite (B), and pearlite (P) were measured using a scanning electron microscope (SEM) at 3,000x and 5,000x magnifications. Here, PF represents polygonal ferrite with an equiaxed crystal shape, and BF includes ferrites observed in low-temperature regions, such as acicular ferrite and bainitic ferrite. Furthermore, P includes pearlite and coarse carbides and nitrides with diameters of 1 μm or more.

Table 4 below shows the physical properties of each manufactured test piece, measured at the center and edge in the width direction. Yield strength (YS), tensile strength (TS), yield ratio (YR), and fracture elongation (T-El) were evaluated by conducting tensile tests on JIS No. 5 standard test pieces taken perpendicular to the rolling direction. In addition, the impact absorption energy (E) at -20°C after cold forming was measured, and the ratio of impact absorption energy to yield strength (E/YS) at -20°C after cold forming is shown. Impact absorption energy was measured using Charpy V-notch test pieces manufactured in accordance with ASTM standards (ASTM A370) taken perpendicular to the rolling direction.

As shown in Tables 3 and 4, inventive examples that satisfy the alloy composition and manufacturing conditions of the present invention, the microstructural characteristics proposed in the present invention were satisfied and the physical properties targeted by the present invention were secured. Figure 1 is a photograph showing the microstructure of the hot-rolled steel sheet of invention example 2 according to one aspect of the present invention, observed with a scanning electron microscope (x3,000).

On the other hand, in Comparative Example 1, the total Ti and Nb content exceeded the range specified by the present invention, and impact resistance deteriorated due to the formation of coarse precipitates and TiN caused by excessive precipitates within the ferrite grains.

In Comparative Example 2, the cooling rate did not meet the cooling rate standard suggested by Relation 1, resulting in excessive pearlite formation within the microstructure, as shown in Figure 2.

As shown in Figure 2, this did not significantly reduce strength, but it did not ensure the desired impact resistance properties.

In Comparative Example 3, the winding temperature in the widthwise center portion exceeded the range proposed by the present invention. In particular, excessive pearlite formed in the thicknesswise center portion in both the center and edge portions, making it impossible to achieve the target impact resistance properties.

In Comparative Example 4, the winding temperature of the widthwise edge portion exceeded the range proposed by the present invention. As a result, it was shown that impact resistance deteriorated due to excessive pearlite formation in the center of the thickness direction of the edge portion. This was because the edge portion temperature was high, and heat transfer from the winding coil proceeded slowly at the edge portion.

In Comparative Example 5, the coiling temperature in the widthwise center portion did not reach the range of the present invention. As a result, excessive bainite was formed in the center portion of the thickness direction of the center portion, and pearlite, MA phase, martensite, etc. were formed in the surface layer portion in excess of the levels proposed by the present invention, resulting in a deterioration in elongation. Meanwhile, the edge portion satisfied the coiling temperature range and had relatively good elongation and impact resistance properties, but ferrite in the surface layer portion did not reach the range proposed by the present invention. This is believed to be due to the fact that the low cooling end temperature in the widthwise center portion caused the temperature at the edge portion of the coil to rapidly decrease after coiling.

In Comparative Example 6, it was confirmed that when the coiling temperature of the widthwise edge portion did not reach the range of the present invention, excessive bainite was formed in the center portion in the thickness direction, resulting in a deterioration in elongation. Furthermore, there was a lack of ferrite in the surface layer of the center portion, and there was a lack of bainitic ferrite in the center portion, which resulted in the ratio of impact resistance to yield strength not meeting the level proposed by the present invention.

In Comparative Example 7, the steel thickness was less than 8 mm. Due to the excessive cooling rate applied for the given steel composition, ferrite was insufficient in the surface layer, bainitic ferrite decreased in the center of the thickness direction, and excessive pearlite formed. This is believed to be due to an increase in untransformed phases during the initial cooling process, resulting in the formation of pearlite in areas with a relatively high C content. As a result, the desired level of elongation could not be achieved.

Comparative Example 8 is a case in which the total Ti and Nb content did not reach the range of the present invention. The phase fractions of the microstructure measured at each position satisfied the range proposed by the present invention, but the microstructure became coarse due to the reduction in precipitates during hot rolling, and the fine precipitates also decreased after cooling and coiling, resulting in insufficient strength and a significant decrease in impact resistance.

In Comparative Example 9, the cooling rate was equal to or greater than the CR value in Relational Formula 1, satisfying the range proposed by the present invention, but exceeding 30°C/s. As a result, there was a shortage of polygonal ferrite in the surface layer, and excessive bainite was formed in the center of the thickness direction, making it impossible to achieve the desired level of elongation.

The present invention has been described in detail above through examples, but other embodiments are possible. Therefore, the technical spirit and scope of the claims set forth below are not limited to the examples.

Claims (9)

The alloy contains, by weight, C: 0.05 to 0.15%, Si: 0.01 to 0.5%, Mn: 1.0 to 2.0%, Al: 0.01 to 0.1%, Cr: 0.001 to 1.0%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, Ti: 0.03 to 0.08%, and Nb: 0.01 to 0.05%, with the balance being Fe and other unavoidable impurities, and the total of Nb and Ti being 0.04 to 0.1%;
The microstructure of the surface layer portion within a thickness range of 50 μm from the surface contains, in area%, 95% or more of equiaxed ferrite and 3% or less of pearlite, and contains 5% or less in total of one or more of bainitic ferrite, bainite, MA (Martensite-Austenite constituent) phase, and martensite,
The microstructure of the center portion within the range of 1/4 to 3/4 of the thickness contains, in area percentages, 80 to 95% bainitic ferrite, 10% or less bainite, 3% or less pearlite, 5 to 10% in total of one or two of an MA (Martensite-Austenite constituent) phase and martensite, and the remainder containing equiaxed ferrite. The steel plate according to claim 1, wherein the steel plate has a thickness of 8 to 25 mm. The steel plate according to claim 1, wherein the steel plate has a tensile strength of 590 MPa or more, a breaking elongation of 25% or more, and a yield ratio of 0.75 to 0.9. The steel plate according to claim 1, wherein the ratio (E/YS) of the impact toughness (E) at -20°C after cold forming to the yield strength (YS) before cold forming is 0.15 or greater. The steel sheet includes edge portions corresponding to 30% regions at both ends in the width direction and a central portion of a central 40% region corresponding to a region excluding both edge portions,
The difference in tensile strength between the edge portion and the center portion is 10 MPa or less, the difference in fracture elongation is 8% or less, and the difference in impact toughness at −20 ° C. after cold forming is 20 J or less. Steel plate according to claim 1. reheating a steel slab containing, by weight, C: 0.05 to 0.15%, Si: 0.01 to 0.5%, Mn: 1.0 to 2.0%, Al: 0.01 to 0.1%, Cr: 0.001 to 1.0%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, Ti: 0.03 to 0.08%, and Nb: 0.01 to 0.05%, with the balance being Fe and other unavoidable impurities, and the sum of Nb and Ti being 0.04 to 0.1%;
hot-rolling the reheated steel slab; and cooling the hot-rolled steel sheet to a temperature range of 500 to 650°C at an average cooling rate of 1 to 30°C/s, the average cooling rate being equal to or greater than the CR (°C/s) value defined by the following Relation 1, and coiling the steel sheet,
In the cooling and winding step, the edge portions corresponding to 30% of the area at both ends in the coil width direction are cooled at a temperature of 550 to 650 ° C. (TE), and the central portion of the central 40% area corresponding to the area excluding both edge portions in the width direction is cooled at a temperature of 500 to 550 ° C. (TC),
The method for producing a steel plate according to claim 1 , wherein the average temperature difference between the edge portion and the center portion is 50 to 150°C.
[Relationship 1]
CR=45-16.3×[C]-5.6×[Si]-16.3×[Mn]-2.9×[Cr]+15×[Ti]+23×[Nb]-0.9×(t-8)
(Here, [C], [Si], [Mn], [Cr], [Ti] and [Nb] are the weight percentages of each element, and t is the thickness of the steel plate (mm).) The reheating temperature is 1100 to 1350°C,
The method for producing a steel sheet according to claim 6, wherein the hot rolling temperature is 800 to 1150°C. The method for manufacturing steel sheet according to claim 6, further comprising the step of air-cooling the wound coil at a temperature range of 200°C or less. The method for manufacturing steel plates described in claim 6, wherein the steel plates have a thickness of 8 to 25 mm.