JP7730903B2 - Ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties and its manufacturing method - Google Patents

Description

The present invention relates to an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties and a manufacturing method thereof. More specifically, the present invention relates to an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties that can be used for automotive structural components such as members, seat rails, pillars, etc., and a manufacturing method thereof.

Recently, stricter safety regulations for automobile passengers and pedestrians have mandated the construction of safety devices, countering efforts to reduce vehicle weight for improved fuel efficiency and resulting in increased vehicle weight. Consumers are increasingly interested in environmentally friendly, fuel-efficient hybrid and electric vehicles. To produce these environmentally friendly and safe vehicles, however, the body structure must be lightweight and the stability of the body materials must be ensured. However, hybrid vehicles incorporate various additional components, such as electric engines, electric batteries, and secondary fuel storage tanks in addition to the conventional gasoline engine. Furthermore, the continued addition of driver convenience facilities increases the vehicle's weight. Therefore, to achieve vehicle weight reduction, it is essential to develop a thin material with excellent strength, ductility, and bending properties. Therefore, to solve this problem, it is necessary to develop giga-grade steel sheets that can ensure high strength and ductility, with a tensile strength of over 980 MPa.

Meanwhile, in recent years, as regulations regarding the impact stability of automobiles have expanded, high-strength steel with excellent yield strength has been adopted for structural components such as members, seat rails, and pillars to improve the impact resistance of vehicle bodies. The higher the yield strength relative to tensile strength, i.e., the higher the yield ratio (yield strength/tensile strength), the better the structural components absorb impact energy. However, as the strength of steel plate generally increases, the elongation decreases, resulting in poor formability. Therefore, there is a need for the development of materials that simultaneously offer improved yield ratios, formability, and bending properties, which are key physical properties when processing parts.

A typical manufacturing method for increasing yield strength is to use water cooling during continuous annealing. Specifically, after soaking in water during the annealing process, the steel sheet is immersed in water and tempered to transform the microstructure from martensite to tempered martensite. Patent Document 1 is a representative example of prior art for this method. Patent Document 1 describes a technique for continuously annealing steel material containing 0.18-0.3% carbon, water cooling it to room temperature, and then overaging it at a temperature of 120-300°C for 1-15 minutes to produce steel material with a martensite volume fraction of 80-97% and the remainder ferrite. When ultra-high-strength steel is manufactured using this water-cooling and tempering method, although the yield ratio is very high, temperature deviations in the width and length directions can cause problems with the coil's shape quality. This can lead to problems during roll forming, such as poor material quality in certain areas and reduced workability.

Patent Document 2 is a prior art technique that improves the workability of the above-mentioned high-tensile steel sheet. Patent Document 2 relates to a steel sheet made of a composite structure primarily composed of tempered martensite, and is characterized by dispersing fine precipitated Cu particles with a particle size of 1 to 100 nm within the structure to improve workability. However, Patent Document 1 has the problem that excessive Cu content of 2 to 5% is added to precipitate good fine Cu particles, which can cause red shortness due to Cu and excessively increase manufacturing costs.

Meanwhile, Patent Document 3 proposes a steel sheet with a ferrite matrix and a microstructure containing 2-10 area percent pearlite, with improved strength achieved through precipitation strengthening and grain refinement, primarily through the addition of carbonitride-forming elements such as Ti. Patent Document 3 has the advantage of easily achieving high strength at low manufacturing costs, but has the disadvantage that the recrystallization temperature rises sharply due to the fine precipitates, necessitating high-temperature annealing to ensure sufficient recrystallization and ductility. Furthermore, existing precipitation-strengthened steels, which are strengthened by precipitating carbonitrides in a ferrite matrix, have the problem that it is difficult to obtain high-strength steels of 600 MPa or higher.

Therefore, there is a need to develop a steel material that can solve the above-mentioned problems, has high yield strength and bending properties, and yet has ultra-high strength that allows for cold forming.

Japanese Patent Registration No. 2528387 Japanese Patent Publication No. 2005-264176 Korean Patent Publication No. 2015-0073844

One aspect of the present invention is to provide an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties, and a method for manufacturing the same.

One embodiment of the present invention contains, in weight percent, C: 0.03-0.12%, Si: 0.03-0.30%, Mn: 2.1-2.9%, Al: 0.005-0.07%, Nb: 0.01-0.08%, Ti: 0.005-0.08%, B: 0.0005-0.005%, Cr: 0.7-1.4%, Mo: 0.005-0.10%, N: 0.008% or less (excluding 0%), balance Fe and other unavoidable impurities, and satisfies the following relations 1 to 3. The microstructure contains, in area percentages, 4 to 19% fresh martensite, 78 to 95% total of tempered martensite and bainite, and 0.2 to 2.0% retained austenite, and the average crystal grain size of the microstructure is 0.5 to 6 μm. This ultra-high strength cold-rolled steel sheet has excellent yield strength and bending properties.

[Relationship 1]
0.18≦C+Si/30+Mn/20+2P+4S≦0.30
[Relationship 2]
180≦48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb≦270
[Relationship 3]
700≦ ( 48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb ) / ( C+Si/30+Mn/20+2P+4S ) ≦1200
(However, the contents of the alloying elements in the above Relational Formulas 1 to 3 are expressed in weight percent.)

Another embodiment of the present invention has, in weight percent, C: 0.03-0.12%, Si: 0.03-0.30%, Mn: 2.1-2.9%, Al: 0.005-0.07%, Nb: 0.01-0.08%, Ti: 0.005-0.08%, B: 0.0005-0.005%, Cr: 0.7-1.4%, Mo: 0.005-0.10%, N: 0.008% or less. a step of heating a slab containing Ar3 (excluding 0%), the balance Fe and other inevitable impurities, and satisfying the following relational expressions 1 to 3; a step of finish-rolling the heated slab so that the delivery temperature of finish rolling is Ar3+50°C to Ar3+150°C to obtain a hot-rolled steel sheet; a step of cooling the hot-rolled steel sheet to Ms+50°C to Ms+300°C and then coiling the hot-rolled steel sheet; The method includes the steps of cold-rolling a coiled hot-rolled steel sheet to obtain a cold-rolled steel sheet, continuously annealing the cold-rolled steel sheet in a temperature range of Ar3 + 10°C to Ar3 + 70°C, soaking the continuously annealed cold-rolled steel sheet for 50 to 200 seconds, primarily cooling the soaked cold-rolled steel sheet to 620 to 700°C at a cooling rate of 1 to 10°C/s, secondary cooling the primarily cooled cold-rolled steel sheet to 360 to 420°C at a cooling rate of 5 to 50°C/s, and overaging the secondarily cooled cold-rolled steel sheet for 250 to 650 seconds, followed by a cooling period of 320 to 400°C. The method satisfies the following relations 4 to 6 during the secondary cooling and overaging treatment, and provides an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties.

[Relationship 1]
0.18≦C+Si/30+Mn/20+2P+4S≦0.30
[Relationship 2]
180≦48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb≦270
[Relationship 3]
700≦ ( 48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb ) / ( C+Si/30+Mn/20+2P+4S ) ≦1200
[Relationship 4]
10≦A≦70
[Relationship 5]
30≦B≦100
[Relationship 6]
2.5≦overaging treatment time/B≦14
(Note that the contents of the alloying elements in the above Relational Formulas 1 to 3 are expressed in weight percent, and in the above Relational Formulas 4 to 6, A is Ms-secondary cooling end temperature (°C), and B is Ms-overaging treatment end temperature (°C).)

One aspect of the present invention provides an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties, and a method for manufacturing the same.

10 is a photograph of the microstructure of Example 6 according to an embodiment of the present invention, observed by SEM.

The following describes an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties according to one embodiment of the present invention. First, the alloy composition of the present invention will be described. The contents of the alloy composition described below are in weight percent unless otherwise specified.

C: 0.03-0.12%
Carbon (C) is a very important element added for solid solution strengthening. Furthermore, carbon contributes to improving strength by combining with precipitated elements to form fine carbides. If the C content is less than 0.03%, it is difficult to achieve the desired strength. On the other hand, if the C content exceeds 0.12%, the increased hardenability leads to excessive martensite formation during cooling, resulting in a rapid increase in strength and reduced bending properties. This may make it difficult to achieve the HER, R/t, and maximum three-point bending angle desired by the present invention. Furthermore, weldability may be reduced, increasing the likelihood of welding defects occurring during component processing at the customer's facility. Therefore, the C content is preferably in the range of 0.03 to 0.12%. The lower limit of the C content is more preferably 0.04%, and even more preferably 0.05%. The upper limit of the C content is more preferably 0.10%, and even more preferably 0.09%.

Si: 0.03~0.30%
Silicon (Si) is one of the five major elements of steel, and a small amount is naturally added during the manufacturing process. Si contributes to increased strength and inhibits carbide formation, preventing carbon from forming as carbides during annealing, soaking, and cooling. Furthermore, this carbon distributes and accumulates in retained austenite, allowing the austenite phase to remain at room temperature, which is beneficial for ensuring elongation. If the Si content is less than 0.03%, it may be difficult to fully achieve the above-mentioned effects. On the other hand, if the Si content exceeds 0.30%, the effect of solid solution strengthening may be enhanced, resulting in reduced elongation and the occurrence of surface scale defects, which may degrade the plating surface quality and reduce chemical conversion treatability. Therefore, the Si content is preferably in the range of 0.03 to 0.30%. The lower limit of the Si content is more preferably 0.04%, and even more preferably 0.05%. The upper limit of the Si content is more preferably 0.25%, and even more preferably 0.20%.

Mn: 2.1-2.9%
Manganese (Mn) completely precipitates sulfur in steel as MnS, preventing hot embrittlement due to the formation of FeS, and solid-solution strengthening the steel. If the Mn content is less than 2.1%, it is difficult to achieve the strength target of the present invention. On the other hand, if the Mn content exceeds 2.9%, problems such as weldability and hot rollability are likely to occur. At the same time, the Mn content increases the hardenability, which can lead to excessive martensite formation and reduced elongation. Furthermore, there are problems with the formation of Mn-bands (Mn oxide bands) in the microstructure, increasing the risk of processing cracks and sheet breakage. Furthermore, there are problems with Mn oxides dissolving on the surface during annealing, significantly impairing galvanic properties. Therefore, the Mn content is preferably in the range of 2.1 to 2.9%. The lower limit of the Mn content is more preferably 2.2%, and even more preferably 2.3%. The upper limit of the Mn content is more preferably 2.8%, and even more preferably 2.7%.

Al: 0.005-0.07%
Aluminum (Al) is an element added for deoxidation during steelmaking. If the Al content is less than 0.005%, it is difficult to obtain a sufficient deoxidizing effect. If it exceeds 0.07%, it reacts with oxygen (O) in the molten steel to form high-melting-point oxides (inclusions), which may cause nozzle clogging. Furthermore, the inclusions formed in this way have sharp shapes, which may result in reduced bending properties. Therefore, the Al content is preferably 0.005 to 0.07%. The lower limit of the Al content is more preferably 0.010%, and even more preferably 0.020%. The upper limit of the Al content is more preferably 0.06%, and even more preferably 0.05%.

Nb: 0.01-0.08%
Niobium (Nb) is an element that segregates at austenite grain boundaries, suppresses coarsening of austenite grains during annealing, and forms fine carbides, contributing to increased strength. If the Nb content is less than 0.01%, the above-mentioned effects are insufficient. On the other hand, if the Nb content exceeds 0.08%, coarse carbides precipitate, reducing the amount of solute carbon in the steel, which can decrease strength and elongation and increase manufacturing costs. Therefore, the Nb content is preferably in the range of 0.01 to 0.08%. The lower limit of the Nb content is more preferably 0.02%, and even more preferably 0.03%. The upper limit of the Nb content is more preferably 0.07%, and even more preferably 0.06%.

Ti: 0.005-0.08%
Titanium (Ti) is a fine carbide-forming element that contributes to ensuring yield strength and tensile strength. Ti is also a nitride-forming element that precipitates N in steel as TiN, suppressing AlN precipitation and reducing the risk of cracking during continuous casting. If the Ti content is less than 0.005%, it may be difficult to achieve the above-mentioned effects. On the other hand, if the Ti content exceeds 0.08%, coarse carbides precipitate, reducing the amount of solute carbon in the steel, which can reduce strength and elongation and cause nozzle clogging during continuous casting. Therefore, the Ti content is preferably in the range of 0.005 to 0.08%. The lower limit of the Ti content is more preferably 0.007%, and even more preferably 0.01%. The upper limit of the Ti content is more preferably 0.07%, and even more preferably 0.06%.

B: 0.0005-0.005%
Boron (B) is an element that significantly contributes to ensuring the hardenability of steel, and to achieve this effect, it is preferable to add 0.0005% or more. However, if the B content exceeds 0.005%, boron carbide is formed at grain boundaries, providing a site for ferrite nucleation, which may actually deteriorate the hardenability. Therefore, the B content is preferably in the range of 0.0005 to 0.005%. The lower limit of the B content is more preferably 0.0010%, and even more preferably 0.0015%. The upper limit of the B content is more preferably 0.0045%, and even more preferably 0.004%.

Cr: 0.7-1.4%
Chromium (Cr) is an element that improves hardenability and increases the strength of steel. If the Cr content is less than 0.7%, it may be difficult to achieve the target strength. On the other hand, if the Cr content exceeds 1.4%, the ductility of the steel sheet may decrease. Therefore, the Cr content is preferably in the range of 0.7 to 1.4%. The lower limit of the Cr content is more preferably 0.75%, and even more preferably 0.8%. The upper limit of the Cr content is more preferably 1.3%, and even more preferably 1.2%.

Mo: 0.005-0.10%
Molybdenum (Mo) is a carbide-forming element. When added in combination with carbonitride-forming elements such as Ti, Nb, and V, it maintains the size of precipitates at a fine size, thereby improving yield strength and tensile strength. Mo also improves the hardenability of steel and finely forms martensite at grain boundaries, allowing for control of the yield ratio. To achieve these effects, Mo is preferably added in an amount of 0.005 % or more. However, since Mo is an expensive element, a higher Mo content can be disadvantageous in manufacturing, so it is preferable to appropriately control its content. A Mo content exceeding 0.10% not only significantly increases manufacturing costs and reduces economic viability, but also reduces the ductility of the steel due to excessive grain refinement and solid solution strengthening. Therefore, the Mo content is preferably in the range of 0.005% to 0.10%. The lower limit of the Mo content is more preferably 0.007%, and even more preferably 0.01%, and the upper limit of the Mo content is more preferably 0.08%, and even more preferably 0.06%.

N: 0.008% or less (excluding 0%)
Nitrogen (N) is an element that is inevitably contained in steel during the manufacturing process, but it contributes to improving the strength of steel by forming carbonitrides. However, if the N content exceeds 0.008%, not only does the risk of brittleness increase significantly, but the excess N remaining after forming TiN may cause B, which contributes to hardenability, to be consumed in the form of BN. Therefore, the N content is preferably 0.008% or less. The N content is more preferably 0.007% or less, and even more preferably 0.006% or less.

On the other hand, the cold-rolled steel sheet of the present invention preferably satisfies the alloy composition described above and also satisfies the following relations 1 to 3. This makes it possible to produce the ultra-high-strength steel sheet with excellent bending workability and a tensile strength of 980 MPa or more, which is the target of the present invention.

[Relationship 1]
0.18≦C+Si/30+Mn/20+2P+4S≦0.30

The above-mentioned Relational Formula 1 is a chemical relationship for ensuring strength and weldability. If the value of the above-mentioned Relational Formula 1 is less than 0.18, it will be difficult to ensure the strength targeted by the present invention, and if it exceeds 0.30, weldability may be reduced. Therefore, the value of the above-mentioned Relational Formula 1 is preferably in the range of 0.18 to 0.30. The lower limit of the value of the above-mentioned Relational Formula 1 is more preferably 0.19, and even more preferably 0.20. The upper limit of the value of the above-mentioned Relational Formula 1 is more preferably 0.28, and even more preferably 0.26.

[Relationship 2]
180≦48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb≦270

The above-mentioned relational expression 2 relates to the hardening ability index for ensuring hardening ability. If the value of the above-mentioned relational expression 2 is less than 180, it will be difficult to ensure the strength targeted by the present invention due to insufficient hardening ability, and if it exceeds 270, the hardening ability will be excessively high, which may result in reduced bending properties and formability. Therefore, the value of the above-mentioned relational expression 2 is preferably in the range of 180 to 270. The lower limit of the value of the above-mentioned relational expression 2 is more preferably 190, and even more preferably 200. The upper limit of the value of the above-mentioned relational expression 2 is more preferably 260, and even more preferably 250.

[Relationship 3]
700≦ ( 48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb ) / ( C+Si/30+Mn/20+2P+4S ) ≦1200

The above-mentioned Relational Formula 3 is a component relationship that simultaneously ensures the strength, hardenability, and weldability targeted by the present invention. If the value of the above-mentioned Relational Formula 3 is less than 700, not only may weldability be reduced, but insufficient hardenability may make it difficult to achieve the strength targeted by the present invention. If the value exceeds 1200, the hardenability may become excessively high, resulting in reduced bending properties and formability. Therefore, the value of the above-mentioned Relational Formula 3 is preferably in the range of 700 to 1200. The lower limit of the value of the above-mentioned Relational Formula 3 is more preferably 700, and even more preferably 800. The upper limit of the value of the above-mentioned Relational Formula 3 is more preferably 1150, and even more preferably 1100.

The remaining component of this invention is iron (Fe). However, during normal manufacturing processes, unintended impurities may inevitably be present from the raw materials or the surrounding environment, and this cannot be ruled out. Since these impurities are known to any engineer working in normal manufacturing processes, this specification will not specifically address all of them.

However, phosphorus and sulfur are the most commonly mentioned impurities, so a brief explanation of these is as follows:

P: 0.04% or less (excluding 0%)
Phosphorus (P) is an element that can cause embrittlement by segregating at grain boundaries and/or interphase boundaries. Therefore, its content must be controlled as low as possible, preferably 0.04% or less. The P content is more preferably 0.03% or less, and even more preferably 0.02% or less.

S: 0.005% or less (excluding 0%)
Sulfur (S) is an impurity that can cause high-temperature cracks by forming MnS non-metallic inclusions in steel and by segregating during continuous casting solidification. Therefore, its content must be controlled as low as possible, preferably to 0.005% or less. The S content is more preferably limited to 0.004% or less, and even more preferably to 0.003% or less.

The above impurities may include one or more of the following tramp elements: Sb, Mg, Sn, Sb, Zn, and Pb, the total of which may be 0.1% by weight or less. Tramp elements are impurity elements derived from scrap and other materials used as raw materials in the steelmaking process, and if the total exceeds 0.1%, it may cause surface cracks in the slab and reduce the surface quality of the steel plate.

The following describes the microstructure and other aspects of an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties according to one embodiment of the present invention.

The microstructure of the cold-rolled steel sheet of the present invention preferably contains, by area percentage, 4-19% fresh martensite, 78-95% total of tempered martensite and bainite, and 0.2-2.0% retained austenite. The microstructure of the cold-rolled steel sheet of the present invention primarily comprises tempered martensite (hereinafter also referred to as "TM") and bainite (hereinafter also referred to as "B"). Because the tempered martensite and bainite are difficult to distinguish from each other in terms of microstructure, the present invention controls the total fraction of the tempered martensite and bainite. If the total fraction of the tempered martensite and bainite is less than 78%, it is difficult to achieve the target strength, and if it exceeds 95%, bending properties and elongation may be reduced. The fresh martensite (hereinafter also referred to as "FM") is a structure advantageous for ensuring strength. If the fraction of the fresh martensite is less than 4%, it is difficult to achieve the target strength, and if it exceeds 19%, bending properties and elongation may be reduced. The retained austenite (hereinafter also referred to as "RA") is a structure advantageous for ensuring elongation. If the fraction of the retained austenite is less than 0.2%, it is difficult to fully achieve the above effect. If it exceeds 2.0%, it may transform into martensite during processing, resulting in a deterioration in HER or bending properties. On the other hand, the microstructure may further include 10% or less of ferrite. Although the ferrite structure is a structure that is inevitably formed in the manufacturing process, it can also fulfill positive functions. For example, the ferrite can contribute to ensuring elongation. However, if the fraction exceeds 10%, it may be difficult to ensure the strength desired by the present invention. The ferrite fraction is more preferably 7% or less, and even more preferably 5% or less.

Meanwhile, the average crystal grain size of the microstructure is preferably 0.5 to 6 μm. The finer the average crystal grain size, the more advantageous it is for ensuring physical properties such as strength and HER. However, controlling the average crystal grain size to less than 0.5 μm may require excessively large amounts of Nb, Ti, Mo, V, and other elements that are effective in refining crystal grains, which may increase manufacturing costs. If the average crystal grain size exceeds 6 μm, it may be difficult to achieve the strength targeted by the present invention, and the HER and bending properties may be significantly reduced. Therefore, the average grain size is preferably in the range of 0.5 to 6.0 μm. The lower limit of the average crystal grain size is more preferably 1.0 μm, and even more preferably 1.5 μm. The upper limit of the average crystal grain size is more preferably 5.5 μm, and even more preferably 5.0 μm.

The cold-rolled steel sheet of the present invention provided as described above may have a yield strength (YS): 800 to 980 MPa, a tensile strength (TS): 980 to 1180 MPa, an elongation (EL): 4 to 12%, a yield ratio (YS/TS): 0.70 to 0.95, a hole expansion ratio (HER): 35 to 80%, an R/t of 0.8 or less, and a maximum three-point bending angle: 90 to 140°. The yield strength is more preferably 820 to 960 MPa, and even more preferably 850 to 950 MPa. The tensile strength is more preferably 1000 to 1170 MPa, and even more preferably 1020 to 1160 MPa. The elongation is more preferably 5 to 11%, and even more preferably 6 to 10%. The yield ratio is more preferably 0.72 to 0.92, and even more preferably 0.75 to 0.90. The hole expansion ratio is more preferably 40 to 75%, and even more preferably 45 to 70%. The R/t is more preferably 0.15 to 0.70, and even more preferably 0.20 to 0.60. The maximum three-point bending angle is more preferably 95 to 135°, and even more preferably 100 to 130°.

The cold-rolled steel sheet of the present invention may also have a hardness (HvBM) of 300 to 400 Hv. The base material hardness is more preferably 310 to 390 Hv, and even more preferably 320 to 380 Hv. Furthermore, the fusion zone (HvFZ) hardness of the weld formed after welding may be 350 to 450 Hv. If the fusion zone hardness of the weld is less than 350 Hv, sufficient fusion zone hardness may not be ensured, potentially resulting in low weld strength. On the other hand, if it exceeds 450 Hv, the fusion zone hardness may be too high, increasing the susceptibility to cracking and reducing the strength and impact absorption energy of the weld. The hardness of the cold-rolled steel sheet, i.e., the hardness (HvBM) corresponding to the base material after welding, should preferably be as similar to the fusion zone hardness (HvFZ), and therefore the ratio (HvFZ/HvBM) is preferably 1.30 or less. The above HvFZ/HvBM ratio is more preferably 1.25 or less, and even more preferably 1.20 or less.

The following describes an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties according to one embodiment of the present invention.

First, a slab satisfying the above-mentioned alloy composition is heated. In the present invention, the slab heating temperature is not particularly limited, but for example, the slab can be heated at 1100 to 1300°C. If the slab heating temperature is less than 1100°C, the slab temperature will be low and rolling loads may occur during rough rolling. If the slab heating temperature exceeds 1300°C, the structure may become coarse, resulting in disadvantages such as increased power costs. The lower limit of the slab heating temperature is more preferably 1125°C, and even more preferably 1150°C. The upper limit of the slab heating temperature is more preferably 1275°C, and even more preferably 1250°C. Meanwhile, the slab can have a thickness of 230 to 270 mm.

The heated slab is then finish-rolled so that the exit temperature of the finish rolling is Ar3+50°C to Ar3+150°C to obtain a hot-rolled steel sheet. If the exit temperature of the finish rolling is less than Ar3+50°C, there is a high possibility that the hot deformation resistance will increase rapidly. If the exit temperature of the finish rolling exceeds Ar3+150°C, there is a high possibility that not only will an excessively thick oxide scale be formed but also the microstructure of the steel sheet will become coarse. Therefore, the exit temperature of the finish rolling is preferably in the range of Ar3+50°C to Ar3+150°C. The lower limit of the exit temperature of the finish rolling is more preferably Ar3+60°C, and even more preferably Ar3+70°C. The upper limit of the exit temperature of the finish rolling is more preferably Ar3+140°C, and even more preferably Ar3+130°C. On the other hand, the above Ar3 means the temperature at which the alloy transforms into austenite when heated, and its value can be determined by the formula, for example, 910-203C ^1/2 +44.7Si+31.5Mo-30Mn-11Cr+700P+400Al+400Ti.

The hot-rolled steel sheet is then cooled to Ms + 50°C to Ms + 300°C and then coiled. If the coiling temperature is less than Ms + 50°C, excess martensite or bainite is generated, resulting in an excessive increase in the strength of the hot-rolled steel sheet, which can lead to problems such as poor shape due to the load during cold rolling. On the other hand, if the coiling temperature exceeds Ms + 300°C, an increase in surface scale can lead to deterioration of pickling properties. Therefore, the coiling temperature is preferably in the range of Ms + 50°C to Ms + 300°C. The lower limit of the coiling temperature is more preferably Ms + 60°C, and even more preferably Ms + 70°C. The upper limit of the coiling temperature is more preferably Ms + 290°C, and even more preferably Ms + 270°C. After coiling, the coiled hot-rolled steel sheet can be cooled to room temperature at a cooling rate of 0.1°C/s or less. The above Ms refers to the temperature at which martensite begins to transform upon cooling, and its value can be calculated using a formula such as 539-423C-30.4Mn-7.5Si+30Al.

The coiled and cooled hot-rolled steel sheet is then cold-rolled to obtain a cold-rolled steel sheet. The cold rolling can be performed at a reduction ratio of 40 to 70%. If the cold rolling reduction ratio is less than 40%, the driving force for recrystallization is weakened, which can make it difficult to obtain good recrystallized grains and makes shape control very difficult. If the cold rolling reduction ratio exceeds 70%, cracks are likely to occur at the edges of the steel sheet and the rolling load may increase rapidly. Therefore, it is preferable to perform the cold rolling at a reduction ratio of 40 to 70%. Meanwhile, pickling can be performed before the cold rolling to remove scale and impurities adhering to the surface.

The cold-rolled steel sheet is then continuously annealed in a temperature range of Ar3 + 10°C to Ar3 + 70°C. If the continuous annealing temperature is below Ar3 + 10°C, the austenite will not be sufficiently transformed, making it difficult to obtain the martensite or bainite fraction desired by the present invention in subsequent processes and potentially making it difficult to ensure strength. On the other hand, if the temperature exceeds Ar3 + 70°C, the austenite grain size will become coarse, making it difficult to achieve the target bending properties. Therefore, the continuous annealing temperature is preferably in the range of Ar3 + 10°C to Ar3 + 70°C. The lower limit of the continuous annealing temperature is more preferably Ar3 + 20°C, and even more preferably Ar3 + 30°C. The upper limit of the continuous annealing temperature is more preferably Ar3 + 60°C, and even more preferably Ar3 + 50°C.

The continuously annealed cold-rolled steel sheet is then soaked for 50 to 200 seconds. This is to ensure a sufficient austenite fraction at the annealing temperature proposed by the present invention, along with recrystallization and grain growth of the cold-rolled structure. If the soaking time is less than 50 seconds, reverse transformation to austenite will not occur sufficiently, increasing the ferrite fraction in the final structure and potentially making it difficult to achieve the target strength. On the other hand, if the soaking time exceeds 200 seconds, the austenite grain size will coarsen, potentially reducing the bending properties of the final product. The lower limit of the soaking time is more preferably 55 seconds, and even more preferably 60 seconds. The upper limit of the soaking time is more preferably 190 seconds, and even more preferably 180 seconds.

The soaked cold-rolled steel sheet is then subjected to primary cooling at a cooling rate of 1 to 10°C/s to 620 to 700°C. The primary cooling stage is intended to ensure the equilibrium carbon concentration of ferrite and austenite and increase the ductility and strength of the steel sheet. If the primary cooling end temperature is less than 630°C or exceeds 700°C, it becomes difficult to achieve the ductility and strength targeted by the present invention. If the cooling rate is less than 1°C/s, the ferrite transformation is accelerated, making it difficult to achieve the targeted microstructure fraction. If the cooling rate exceeds 10°C/s, excessive martensite transformation makes it difficult to achieve the desired elongation.

The primarily cooled cold-rolled steel sheet is then subjected to secondary cooling to 360-420°C at a cooling rate of 5-50°C/s. This secondary cooling is one of the key control factors in the present invention, and the secondary cooling end temperature is an extremely important condition for simultaneously ensuring strength, ductility, and bending properties. If the secondary cooling end temperature is below 360°C, it becomes difficult to ensure ductility due to an excessive increase in martensite fraction. If it exceeds 420°C, it becomes difficult to ensure sufficient martensite, making it difficult to achieve the target strength. Therefore, the secondary cooling end temperature, which is one of the key control factors for ensuring the target physical properties in the present invention, is preferably in the range of 360-420°C. The lower limit of the secondary cooling end temperature is more preferably 365°C, and even more preferably 370°C. The upper limit of the secondary cooling end temperature is more preferably 410°C, and even more preferably 405°C. If the secondary cooling rate is less than 5°C/s, ferrite transformation occurs preferentially before martensite and bainite transformation due to the slow cooling rate, resulting in the failure to obtain the appropriate amount of fine structure fraction desired by the present invention. If the rate exceeds 50°C/s, the excessive cooling rate can cause shape deterioration, resulting in reduced sheet threadability and sheet breakage. The lower limit of the secondary cooling rate is more preferably 7.5°C/s, and even more preferably 10°C/s. The upper limit of the secondary cooling rate is more preferably 47.5°C/s, and even more preferably 45°C/s.

On the other hand, in order to ensure the target levels of tempered martensite and bainite fractions, which are important microstructures in this invention, it is important to precisely control the difference between the Ms temperature and the secondary cooling end temperature. More specifically, it is preferable to satisfy the following relational expression 4. If the difference between Ms and the secondary cooling end temperature, i.e., the value of A, is less than 10°C, there will be little martensite or bainite transformation, making it difficult to achieve the target strength. If the value of A exceeds 70°C, the time spent in the martensite region will be long, resulting in an excessive increase in the martensite fraction and making it difficult to achieve ductility. Therefore, it is preferable that the difference between Ms and the secondary cooling end temperature, i.e., the value of A, be 10 to 70°C. The lower limit of the above value of A is more preferably 15°C, and even more preferably 20°C. The upper limit of the above value of A is more preferably 65°C, and even more preferably 60°C. On the other hand, Ms refers to the temperature at which martensitic transformation begins, and its value can be calculated using the following formula 1.

[Relationship 4]
10≦A≦70
(However, in the above relational expression 4, A is Ms-secondary cooling end temperature (°C).)

The secondary-cooled cold-rolled steel sheet is then overaged for 250 to 650 seconds, and then terminated at 320 to 400°C. The overaging is preferably performed at a temperature equal to or higher than the temperature at the end of the secondary cooling. The overaging promotes the transformation of fresh martensite formed at the end of the secondary cooling into tempered martensite, thereby stably ensuring high yield strength and bending properties. Therefore, overaging is a crucial factor in ensuring the high bending workability desired in the present invention. In the present invention, the overaging time is precisely controlled within the range of 250 to 650 seconds. If the overaging time is less than 250 seconds, slight transformation from fresh martensite to tempered martensite may occur, potentially reducing bending workability. On the other hand, if the overaging time exceeds 650 seconds, productivity may decrease and excessive tempered martensite transformation may make it difficult to achieve the tensile strength desired in the present invention. The lower limit of the overaging treatment time is more preferably 260 seconds, and even more preferably 270 seconds. The upper limit of the overaging treatment time is more preferably 600 seconds, and even more preferably 550 seconds. If the overaging treatment end temperature is less than 320°C, excessive fresh martensite transformation may make it difficult to ensure elongation, and bending properties may be reduced. If the overaging treatment end temperature is more than 400°C, a slight transformation from fresh martensite to tempered martensite may occur, and bending properties may be reduced. The lower limit of the overaging treatment end temperature is more preferably 325°C, and even more preferably 330°C. The upper limit of the overaging treatment end temperature is more preferably 395°C, and even more preferably 380°C. On the other hand, to further improve the HER and bending properties, after the overaging treatment, the steel may be cooled to the secondary cooling end temperature and then reheated to perform further overaging treatment.

On the other hand, in order to ensure the target level of tempered martensite fraction, which is an important microstructure in the present invention, it is important to precisely control the difference between the Ms temperature and the overaging treatment end temperature. More specifically, it is preferable to satisfy the following relational expression 5. If the difference between the Ms temperature and the overaging treatment end temperature, i.e., the value of B, is less than 30°C, martensitic transformation may be insufficient, making it difficult to achieve the target strength. If the value of B exceeds 100°C, excessive fresh martensitic transformation may occur, making it difficult to achieve the target elongation and bending properties. Therefore, it is preferable that the difference between the Ms temperature and the overaging treatment end temperature, i.e., the value of B, be 30 to 100°C. The lower limit of the B value is more preferably 35°C, and even more preferably 40°C. The upper limit of the B value is more preferably 95°C, and even more preferably 90°C.

[Relationship 5]
30≦B≦100
(In the above relational expression 5, B is Ms - overaging treatment end temperature (°C).)

In the present invention, in order to achieve the target microstructure fraction and mechanical properties, it is preferable that the following relational expression 6 be satisfied during the secondary cooling and overaging treatment.

[Relationship 6]
2.5≦overaging treatment time/B≦14

The above-mentioned relational expression 6 is used in the present invention to precisely control the target microstructure and ensure the target physical properties. If the value of the above-mentioned relational expression 6 is less than 2.5, the overaging holding time is short or the overaging end temperature is low, and excessive fresh martensitic transformation may occur, making it difficult to achieve the target elongation and bending properties. On the other hand, if the value of the above-mentioned relational expression 6 exceeds 14, the overaging holding time may be long or the overaging end temperature may be high, making it difficult to achieve the target microstructure fraction and the target physical properties. Therefore, the value of the above-mentioned relational expression 6 is preferably in the range of 2.5 to 14. The lower limit of the value of the above-mentioned relational expression 6 is more preferably 3.0, and even more preferably 3.5. The upper limit of the value of the above-mentioned relational expression 6 is more preferably 12, and even more preferably 10.

Meanwhile, the present invention may further include a step of temper rolling the overaged cold-rolled steel sheet at an elongation rate of 0.1 to 2.0% after the overaging treatment. Typically, temper rolling results in little increase in tensile strength, but an increase in yield strength of at least 50 MPa or more. If the elongation rate is less than 0.1%, it may be difficult to control the shape, and if it exceeds 2.0%, high elongation may result in significant instability in operability.

The present invention will be described in more detail through the following 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. The scope of the present invention is determined by the matters set forth in the claims and matters that can be reasonably inferred therefrom.

Example 1
Molten steel having the alloy composition shown in Table 1 below was prepared, formed into a 250 mm thick slab, heated at 1200°C for 12 hours, and then finish-rolled and coiled under the conditions shown in Table 2 below to produce a hot-rolled steel sheet. The hot-rolled steel sheet thus produced was pickled and then cold-rolled at a cold reduction of 50% to produce a cold-rolled steel sheet. This cold-rolled steel sheet was subjected to continuous annealing, soaking, primary and secondary cooling, and overaging under the conditions shown in Table 3 below to produce a final cold-rolled steel sheet.

The microstructure, average grain size, and mechanical properties of the final cold-rolled steel sheet produced in this manner were measured, and the results are shown in Table 4 below.

Microstructure and average grain size were measured using electron backscatter diffraction (EBSD) equipment.

Among the mechanical properties, tensile strength (TS), yield strength (YS), and elongation (EL) were measured by tensile testing after tensile test specimens were taken in the horizontal direction of rolling. The gauge length and width of the tensile test specimens used were 80 mm and 20 mm, respectively.

The fusion zone hardness (HvFZ) and base material hardness (HvBS) were measured _five times at ¼ t (t = thickness point) using a Vickers hardness tester under a load of 500 gf after BOP (Bead On Plate) welding was performed on the cold rolled steel sheets using a CO2 laser welder under the conditions of 6 kW-3 min, and then the hardness was averaged.

As can be seen from Tables 1 to 4 above, Examples 1 to 5, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, are able to achieve the microstructure sought by the present invention and have excellent mechanical properties.

On the other hand, Comparative Examples 1 to 8 do not meet the alloy composition proposed by the present invention and do not meet some of the manufacturing conditions, so the microstructure sought by the present invention cannot be achieved and mechanical properties are reduced.

Example 2
A molten steel having the alloy composition of Inventive Steel 1 described in Example 1 was prepared, and then formed into a 250 mm thick slab. The slab was heated at 1200°C for 12 hours, and then finish-rolled under the conditions shown in Table 5 below, followed by coiling to produce a hot-rolled steel sheet. The hot-rolled steel sheet thus produced was pickled, and then cold-rolled at a cold reduction of 50% to produce a cold-rolled steel sheet. This cold-rolled steel sheet was subjected to continuous annealing, soaking, primary and secondary cooling, and overaging under the conditions shown in Table 6 below, to produce a final cold-rolled steel sheet.

The microstructure, average grain size, and mechanical properties of the final cold-rolled steel sheet produced in this manner were measured, and the results are shown in Table 7 below.

Microstructure and average grain size were measured using electron backscatter diffraction (EBSD) equipment.

Among the mechanical properties, tensile strength (TS), yield strength (YS), and elongation (EL) were measured by tensile testing after tensile test specimens were taken in the horizontal direction of rolling. The gauge length and width of the tensile test specimens used were 80 mm and 20 mm, respectively.

Of the mechanical properties, HER was measured according to the ISO 16330 standard, and the holes were sheared using a 10 mm diameter punch with a 12% clearance.

Among the mechanical properties, R/t is the value obtained by dividing R (limit bending radius) by the thickness of the steel plate. Here, the above R was calculated by taking a test piece 30 mm wide x 35 mm long in the direction horizontal to the rolling direction (long axis), grinding one side by 0.2 mm, and then conducting a bending test using the V-block method in accordance with JIS Z 2248 so that the ground surface did not come into contact with the punch. The bending radius was then varied from 0 to 5 mm, and the minimum bending radius at which the material could be bent without breaking was determined.

The maximum three-point bending angle, one of the mechanical properties, was measured three times for each test piece in accordance with VDA (Verband Der Automobilindustrie) standards, and the average value was calculated.

As can be seen from Tables 5 to 7 above, in the case of Examples 6 to 10, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, the type and fraction of microstructure and average grain size targeted by the present invention are secured, thereby ensuring the mechanical properties (tensile properties, HER, bending properties) that the present invention aims to achieve.

On the other hand, while Comparative Examples 9 to 16 satisfy the alloy composition proposed by the present invention, they do not meet the manufacturing conditions, and therefore the fine structure or average crystal grain size targeted by the present invention cannot be achieved, resulting in reduced mechanical properties. In particular, in the cases of Comparative Examples 9 and 10, the finish rolling temperature and coiling temperature did not satisfy the conditions of the present invention, resulting in sheet breakage.

Figure 1 is a photograph of the microstructure of Invention Example 6 observed with an SEM. As can be seen from Figure 1, Invention Example 6 confirms that the microstructure targeted by the present invention is properly formed.

Example 3
A molten steel having the alloy composition of Inventive Steel 2 described in Example 1 was prepared, and then formed into a 250 mm thick slab. The slab was heated at 1200°C for 12 hours, and then finish-rolled under the conditions shown in Table 8 below, followed by coiling to produce a hot-rolled steel sheet. The hot-rolled steel sheet produced in this manner was pickled, and then cold-rolled at a cold reduction of 50% to produce a cold-rolled steel sheet. This cold-rolled steel sheet was subjected to continuous annealing, soaking, primary and secondary cooling, and overaging under the conditions shown in Table 9 below, to produce a final cold-rolled steel sheet.

The microstructure, average grain size, and mechanical properties of the final cold-rolled steel sheet produced in this manner were measured, and the results are shown in Table 10 below.

Microstructure and average grain size were measured using electron backscatter diffraction (EBSD) equipment.

The maximum three-point bending angle, one of the mechanical properties, was measured three times for each test piece in accordance with VDA (Verband Der Automobilindustrie) standards, and the average value was calculated.

As can be seen from Tables 8 to 10 above, in the case of Examples 11 to 15, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, the microstructure targeted by the present invention is maintained, thereby ensuring the bending properties that the present invention aims to achieve.

On the other hand, while Comparative Examples 17 to 20 satisfy the alloy composition proposed by the present invention, they do not satisfy the overaging treatment conditions and relational expressions 5 and 6 among the manufacturing conditions. As a result, the microstructure targeted by the present invention cannot be achieved, and bending properties are reduced.

Claims (12)

The alloy contains, by weight, C: 0.03 to 0.12%, Si: 0.03 to 0.30%, Mn: 2.1 to 2.9%, Al: 0.005 to 0.07%, Nb: 0.01 to 0.08%, Ti: 0.005 to 0.08%, B: 0.0005 to 0.005%, Cr: 0.7 to 1.4%, Mo: 0.005 to 0.10%, N: 0.008% or less (excluding 0%), with the balance being Fe and other unavoidable impurities;
The following relations 1 to 3 are satisfied:
The microstructure contains, in area percentage, fresh martensite: 4 to 19%, the sum of tempered martensite and bainite: 78 to 95%, and retained austenite: 0.2 to 2.0%.
The average crystal grain size of the microstructure is 0.5 to 6 μm, and the ultra-high strength cold-rolled steel sheet has excellent yield strength and bending properties.
[Relationship 1]
0.18≦C+Si/30+Mn/20+2P+4S≦0.30
[Relationship 2]
180≦48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb≦270
[Relationship 3]
700≦ ( 48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb ) / ( C+Si/30+Mn/20+2P+4S ) ≦1200
(However, the contents of the alloying elements in the above Relational Formulas 1 to 3 are expressed in weight percent.) An ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties as described in claim 1, wherein the impurities further include P: 0.04% or less (excluding 0%) and S: 0.005% or less (excluding 0%). An ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties as described in claim 1, wherein the impurities include one or more of Sb, Mg, Sn, Sb, Zn, and Pb, the total of which is 0.1 wt% or less. An ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties as described in claim 1, wherein the microstructure further contains 10% or less ferrite. The cold-rolled steel sheet according to claim 1 has an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties, which has a yield strength (YS): 800 to 980 MPa, a tensile strength (TS): 980 to 1180 MPa, an elongation (EL): 4 to 12%, a yield ratio (YS/TS): 0.70 to 0.95, a hole expansion ratio (HER): 35 to 80%, an R/t: 0.8 or less, and a maximum three-point bending angle: 90 to 140°. The cold-rolled steel sheet has a hardness (HvBM) of 300 to 400 Hv, a fusion zone hardness (HvFZ) of the weld formed after welding of 350 to 450 Hv, and HvFZ/HvBM is 1.30 or less. An ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties as described in claim 1. A method for producing an ultra-high strength cold-rolled steel sheet having excellent yield strength and bending properties according to claim 1,
heating a slab containing, by weight, C: 0.03 to 0.12%, Si: 0.03 to 0.30%, Mn: 2.1 to 2.9%, Al: 0.005 to 0.07%, Nb: 0.01 to 0.08%, Ti: 0.005 to 0.08%, B: 0.0005 to 0.005%, Cr: 0.7 to 1.4%, Mo: 0.005 to 0.10%, N: 0.008% or less ( excluding 0%), with the balance being Fe and other unavoidable impurities, and satisfying the following relations 1 to 3;
Finish rolling the heated slab so that the delivery temperature of the finish rolling is Ar3+50°C to Ar3+150°C to obtain a hot-rolled steel sheet;
cooling the hot-rolled steel sheet to Ms+50°C to Ms+300°C and then coiling it;
cold rolling the coiled hot-rolled steel sheet to obtain a cold-rolled steel sheet;
continuous annealing the cold-rolled steel sheet in a temperature range of Ar3+10°C to Ar3+70°C;
soaking the continuously annealed cold-rolled steel sheet for 50 to 200 seconds;
subjecting the soaked cold-rolled steel sheet to primary cooling at a cooling rate of 1 to 10°C/s to 620 to 700°C;
Secondarily cooling the primarily cooled cold-rolled steel sheet to 360 to 420°C at a cooling rate of 5 to 50°C/s;
and overaging the second-cooled cold-rolled steel sheet for 250 to 650 seconds, and then finishing at 320 to 400°C.
A method for producing an ultra-high strength cold-rolled steel sheet having excellent yield strength and bending properties, which satisfies the following Relational Expressions 4 to 6 during the secondary cooling and overaging treatment:
[Relationship 1]
0.18≦C+Si/30+Mn/20+2P+4S≦0.30
[Relationship 2]
180≦48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb≦270
[Relationship 3]
700≦ ( 48.8+49logC+35.1Mn+25.9Si+76.5Cr+105.9Mo+1325Nb ) / ( C+Si/30+Mn/20+2P+4S ) ≦1200
[Relationship 4]
10≦A≦70
[Relationship 5]
30≦B≦100
[Relationship 6]
2.5≦overaging treatment time/B≦14
(Note that the contents of the alloy elements in the above Relational Formulas 1 to 3 represent weight percent, and in the above Relational Formulas 4 to 6, A is Ms-secondary cooling end temperature (°C), and B is Ms-overaging treatment end temperature (°C).) The method for manufacturing ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties described in claim 7, wherein the slab heating is performed at 1100 to 1300°C. The method for manufacturing ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties described in claim 7, wherein the slab has a thickness of 230 to 270 mm. The method for manufacturing an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties described in claim 7 further comprises, after the coiling, a step of cooling the coiled hot-rolled steel sheet to room temperature at a cooling rate of 0.1°C/s or less. The method for manufacturing ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties described in claim 7, wherein the cold rolling is performed at a reduction ratio of 40 to 70%. The method for manufacturing an ultra-high strength cold-rolled steel sheet with excellent yield strength and bending properties described in claim 7 further comprises, after the overaging treatment, temper rolling the overaged cold-rolled steel sheet at an elongation rate of 0.1 to 2.0%.