Hot-rolled steel sheet and method for producing the same

Abstract

A hot-rolled steel sheet and a method for manufacturing the same are provided. The hot-rolled steel sheet of the present invention contains, by weight%, C: 0.03 to 0.08%, Si: 0.01 to 1.0%, Mn: 1.0 to 2.0%, Sol.Al: 0.01 to 0.1%, Cr: 0.005 to 0.5%, Mo: 0.005 to 0.3%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, Ti: 0.005 to 0.12%, Nb: 0.005 to 0.06%, V: 0.005 to 0.2%, B: 0.0003 to 0.003%, with the balance being Fe and inevitable impurities, and satisfies relational expression 1-2.

Description

Technical Field

[0001] The present invention mainly relates to the manufacture of high-strength hot-rolled steel sheets used for members of automobile chassis parts, lower arms, reinforcing materials, and connecting materials. More specifically, the tensile strength is 590 MPa or more, the HER0 (Hole Expansion Ratio of specimen with machined hole) is excellent at 100% or more, and the HER value, which is the actual flanging formability, satisfies that the product of the tensile strength and HER is 45,000 (MPa%) or more even in a wide range where the punching clearance is 5 to 20%. The present invention relates to a high-strength hot-rolled steel sheet and a method for manufacturing the same.

Background Art

[0002] Conventional high-strength hot-rolled steel sheets for chassis parts are manufactured into two-phase composite structure steels with a ferrite-bainite mixed structure as the basic matrix structure to improve formability and improve elongation flangibility, or high-strength high-burring steels with a ferrite phase or bainite phase as the basic matrix structure have been proposed.

[0003] In Patent Document 1, immediately after hot rolling the steel, it is maintained for several seconds in the ferrite transformation range according to specific cooling conditions, and then wound up at the bainite formation temperature so that bainite is formed. By forming the metal structure into a mixed structure of polygonal ferrite and bainite, it was attempted to ensure both strength and elongation flangibility at the same time. In Patent Document 2, a high-burring steel with bainitic ferrite and granular bainitic ferrite as the matrix structure based on the C-Si-Mn component system was proposed. In Patent Document 3, a technique for improving elongation flangibility by manufacturing a steel having 95% or more of the bainite phase and having few crystal grains stretched in the rolling direction was proposed. However, Patent Document 1 is a technique that excessively uses Si so as to delay the formation of pearlite and make the formation of the bainite phase advantageous. The surface quality such as red scale defects deteriorates, and since the ferrite phase fraction, which is a soft phase, is 90% or more, it is not suitable for the manufacture of high-strength steel.

[0004] Patent Document 2 discloses a technology based on bainitic ferrite, which is a low-temperature type ferrite phase, and granular bainitic ferrite. However, in order to ensure further strength, Cu must be utilized, which may cause surface defects and high-temperature brittleness. Therefore, Ni must be added to prevent this, which is economically disadvantageous.

[0005] Moreover, Patent Document 3 states that in order to form fewer elongated crystal grains, hot rolling must be carried out, and in order to manufacture with a bainite phase as the base structure, cooling must be performed at an excessively high cooling rate. As a result, local differences in cooling rate are likely to occur, and there is a problem that the shape quality of the rolled plate is likely to deteriorate.

[0006] In addition, alloying elements such as Si, Mn, Mo, Cr, Cu, and Ni, which are mainly utilized for manufacturing high-strength steel as described above, are effective in improving the strength and elongation flange formability of the hot-rolled steel sheet. However, if a large amount of alloying elements are added to improve such physical properties, segregation of alloying elements and non-uniformity of the microstructure will occur, resulting in a decrease in elongation flange formability. In particular, steels with high hardenability are sensitive to changes in the microstructure during cooling, and non-uniform low-temperature transformation tissue phases are formed, making it difficult to obtain higher elongation flange formability.

Prior Art Documents

Patent Documents

[0007]

Patent Document 1

Patent Document 2

Patent Document 3

Summary of the Invention

Problems to be Solved by the Invention

[0008] The present invention provides a hot-rolled steel sheet having a tensile strength of 590 MPa or more, excellent HER0 of 100% or more, and an HER value which is the actual flanging formability, and the product of the tensile strength × HER satisfies 45,000 (MPa%) or more even in a wide range of punching clearances of 5 to 20%, and a method for manufacturing the same.

[0009] Further, the technical problems to be achieved in the present invention are not limited to the above-described technical problems, and other technical problems not mentioned can be clearly understood by those having ordinary knowledge in the technical field to which the present invention pertains from the following description.

Means for Solving the Problems

[0010] One aspect of the present invention is By weight, C: 0.03 to 0.08%, Si: 0.01 to 1.0%, Mn: 1.0 to 2.0%, Sol.Al: 0.01 to 0.1%, Cr: 0.005 to 0.5%, Mo: 0.005 to 0.3%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, Ti: 0.005 to 0.12%, Nb: 0.005 to 0.06%, V: 0.005 to 0.2%, B: 0.0003 to 0.003%, the balance being Fe and unavoidable impurities, satisfying the following relational expressions 1 and 2, The microstructure is mainly composed of a ferrite-based low-temperature transformation phase composed of one or more of acicular ferrite, granular bainitic ferrite, and bainitic ferrite, the total of the bainite phase and the polygonal ferrite phase is less than 40%, and the total of the remaining pearlite, martensite, retained austenite, and MA phase is less than 3%, and the average dislocation density (Geometrically Necessary Dislocation) of the above microstructure is 1.0×10 14 ~2.5×10 14 m -2 relates to a hot-rolled steel sheet satisfying the above.

[0011] [Relational Expression 1] 0.2 ≤ X ≤ 0.6 X = (Nb / 93 + Ti* / 48 + V / 51) / (C / 12 + N / 14) Ti* = Ti - 3.42N - 1.5S

[0012] In the above relational expression 1, Nb, Ti, C, N, and S are the weight percentages of the corresponding alloying elements, and 0 is substituted when they are not added.

[0013] [Relational Expression 2] 1.5 ≤ T ≤ 3.5 T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B]

[0014] In the above relational expression 2, Mn, Mo, Cr, and B are the weight percentages of the corresponding alloying elements, and 0 is substituted when they are not added.

[0015] The above hot-rolled steel sheet has a tensile strength of 590 MPa or more, and the HER value can satisfy that the product of the tensile strength × HER is 45,000 MPa% or more in the range where the punching clearance is 5 - 20%.

[0016] Also, another aspect of the present invention is By weight, C: 0.03 - 0.08%, Si: 0.01 - 1.0%, Mn: 1.0 - 2.0%, Sol.Al: 0.01 - 0.1%, Cr: 0.005 - 0.5%, Mo: 0.005 - 0.3%, P: 0.001 - 0.05%, S: 0.001 - 0.01%, N: 0.001 - 0.01%, Ti: 0.005 - 0.12%, Nb: 0.005 - 0.06%, V: 0.005 - 0.2%, B: 0.0003 - 0.003%, the balance being Fe and unavoidable impurities, reheating a steel slab satisfying the following relational expressions 1 and 2 to 1150 - 1350°C; Hot-rolling the reheated steel slab at a temperature in the range of 850 - 1150°C; First cooling the hot-rolled steel sheet at an average cooling rate of 50 - 100°C / sec to a temperature in the range of 400 - 500°C; Reheating the steel plate maintained for 0.5 to 3 seconds after the above primary cooling to a temperature in the range of 450 to 550 °C; After secondary cooling the reheated steel plate to a temperature in the range of 400 to 500 °C at an average cooling rate of 1 to 30 °C / sec and then winding it up; and Cooling the wound coil at an average cooling rate of 0.1 to 25 °C / hour to a temperature in the range of normal temperature to 200 °C, In one or more of the above primary cooling step and the above secondary cooling step, the temperature (T E ) of both edge portions (W×60%) with respect to the entire width (W) of the steel plate is 450 to 550 °C, and the temperature (T C ) of the center portion of the width (W×40%) is 400 to 500 °C, and by finishing the cooling, after winding in the above winding step, the average temperature (T A ) of the coil can be maintained in the range of 400 to 500 °C, relating to a method for manufacturing a hot-rolled steel plate.

[0017] [Relational Expression 1] 0.2 ≦ X ≦ 0.6 X = (Nb / 93 + Ti* / 48 + V / 51) / (C / 12 + N / 14) Ti* = Ti - 3.42N - 1.5S

[0018] In the above Relational Expression 1, Nb, Ti, C, N, and S are the weight % of the corresponding alloying elements, and when not added, 0 is substituted.

[0019] [Relational Expression 2] 1.5 ≦ T ≦ 3.5 T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B]

[0020] In the above Relational Expression 2, Mn, Mo, Cr, and B are the weight % of the corresponding alloying elements, and when not added, 0 is substituted.

Advantages of the Invention

[0021] The present invention having the above-described configuration can effectively provide a high-strength steel sheet having a tensile strength of 590 MPa or more, excellent HER0 of 100% or more, and an HER value, which is an actual flanging formability, satisfying a product of tensile strength × HER of 45,000 MPa% or more even in a wide punching clearance range of 5 to 20% by controlling the microstructure and average dislocation density of steel.

Brief Description of the Drawings

[0022]

Figure 1

Modes for Carrying Out the Invention

[0023] Hereinafter, the present invention will be described.

[0024] In order to solve the above-described problems of the prior art, the inventors of the present invention examined the changes in local formability and flanging property (HER, Hole Expansion ratio) due to the characteristics of components, manufacturing processes, and microstructures for steels having different components and microstructures from each other. As a result, it was confirmed that each detailed structural feature has a correlation with local formability and flanging property, and thereby the following relational expressions 1 and 2 were derived. Based on the derived relational expressions, the microstructure of the steel has a ferrite-based low-temperature transformation phase as the main phase, a total of less than 40% of bainite phase and polygonal ferrite phase, 10% or more of the above polygonal ferrite, and the remaining pearlite, martensite, and MA phase in a total of less than 3%, and the average dislocation density (Geometrically Necessary Dislocation) value is 1.0×10 14 ~2.5×10 14 m -2When the range is satisfied, the tensile strength is 590 MPa or more, HER0 is excellent at 100% or more, and the HER value, which is the actual flanging formability, satisfies that the product of the tensile strength and HER is 45,000 MPa% or more even in a wide range where the punching clearance is 5 to 20%. The present invention is presented by confirming this.

[0025] Hereinafter, the reasons for restricting the composition components and their contents of the hot-rolled steel sheet of the present invention will be described. Here, “%” means weight % unless otherwise specified.

[0026] C: 0.03 to 0.08% The above C is the most economical and effective element in steel strengthening and greatly affects the hardness values of each structure. When its addition amount increases, the hardening ability increases, the phase fraction of hard phases such as bainite phase and martensite phase in the fine structure increases, and the tensile strength increases. In addition, fine precipitates are formed together with Ti and Nb having a high affinity for C, and both the yield strength and the tensile strength increase due to precipitation strengthening. However, if the content is less than 0.03%, it is difficult to obtain a sufficient strengthening effect. If it exceeds 0.08%, there are problems that the hardness of each phase including the bainite phase and the martensite phase increases, excessive strength increase and formability decrease, and weldability also decreases. Therefore, the above C content is preferably included in the range of 0.03 to 0.08%. In order to stably ensure the strength and formability at the target level in the present invention, it is more preferably limited to 0.04 to 0.07%.

[0027] Si: 0.01 to 1.0%, The above Si deoxidizes molten steel, has a solid solution strengthening effect, and is advantageous in delaying the formation of coarse carbides and improving formability. In the present invention, in order to obtain such an effect, it is preferably contained in an amount of 0.01% or more. However, if it exceeds 1.0%, red scale due to Si is formed on the steel sheet surface during hot rolling, not only the quality of the steel sheet surface becomes very poor, but also there are problems that ductility and weldability decrease. Therefore, it is preferable to limit its content to 1.0% or less.

[0028] Mn: 1.0 to 2.0% The above-mentioned Mn is an element effective for solid solution strengthening of steel, similar to Si. It increases the hardenability of steel and facilitates the formation of bainite phase and martensite phase, which are hard phases, during cooling after hot rolling. However, if its content is less than 1.0%, the above-mentioned effects due to addition cannot be obtained. If it exceeds 2.0%, the hardenability increases significantly, and the hardness of each phase, including bainite phase and martensite phase, increases, resulting in problems such as excessive increase in strength and decrease in formability. During the continuous casting process, segregation zones develop significantly in the thickness center during slab casting. During cooling after hot rolling, non-uniform fine structures are formed in the thickness direction, and the elongation flange property decreases. In particular, it becomes difficult to produce uniform fine structures during cooling over the entire length and width of the hot rolled sheet. More preferably, the upper limit of the above-mentioned Mn content is restricted to 1.8%.

[0029] Sol.Al: 0.01 - 0.1% The above-mentioned Sol.Al is mainly a component added for deoxidation. If its content is less than 0.01%, the effects due to addition are insufficient. If it exceeds 0.1%, it combines with nitrogen to form AlN, and corner cracks are likely to occur in the slab during continuous casting, and defects due to the formation of inclusions are likely to occur. Therefore, it is preferably restricted to 0.01 - 0.1%, and more preferably, the upper limit of its addition amount is limited to 0.06%.

[0030] Cr: 0.005 - 0.5% The above-mentioned Cr plays a role in solid solution strengthening of steel and delaying the ferrite phase transformation during cooling, contributing to the formation of bainite. However, if its addition amount is less than 0.005%, the above-mentioned effects due to addition cannot be obtained. If it exceeds 0.5%, the ferrite transformation is overly delayed, and the elongation rate decreases due to the formation of martensite phase. Also, similar to Mn, segregation zones develop significantly in the thickness center, making the fine structures in the thickness direction non-uniform and reducing the elongation flange property. More preferably, the upper limit of the above-mentioned Cr content is restricted to 0.3%.

[0031] Mo: 0.005 - 0.3% Mo increases the hardenability of steel and facilitates the formation of a bainite structure. However, if the addition amount is less than 0.005%, the above effects due to the addition cannot be obtained. If it exceeds 0.3%, a martensite phase is formed due to excessive hardenability increase, and the formability decreases rapidly. Also, it is economically disadvantageous and harmful to weldability. Therefore, it is preferable to limit the content of the above Mo to 0.005 - 0.3%. More preferably, the Mo content is limited to 0.05 - 0.2%.

[0032] P: 0.001 - 0.05% The above P, like Si, has both a solid solution strengthening and a ferrite transformation promoting effect at the same time. However, in order to control its content to less than 0.001%, it requires a high manufacturing cost, so it is economically disadvantageous and insufficient to obtain strength. On the other hand, if its content exceeds 0.05%, brittleness due to grain boundary segregation occurs, and fine cracks are likely to occur during forming, greatly deteriorating ductility, elongation flangeability, and impact resistance. Therefore, it is preferable to limit the content of the above P to 0.001 - 0.05%.

[0033] S: 0.001 - 0.01% The above S is an impurity present in steel. When its content exceeds 0.01%, it combines with Mn etc. to form non-metallic inclusions, which causes a problem that fine cracks are likely to occur during cutting processing of steel, greatly reducing elongation flangeability and impact resistance. In the present invention, the lower limit of the S content is not particularly limited. However, in order to manufacture at less than 0.001%, it takes a long time during steelmaking operation and productivity decreases, so considering this, the lower limit of its content can be limited to 0.001%. Therefore, in the present invention, it is preferable to limit the content of S to 0.001 - 0.01%, and it is more preferable to limit its upper limit to 0.005%.

[0034] N: 0.001 - 0.01% The above-mentioned N is a typical solid solution strengthening element together with C, and forms coarse precipitates together with Ti, Al, etc. Generally, the solid solution strengthening effect of N is superior to that of carbon, but there is a problem that the toughness decreases significantly as the amount of N in the steel increases. On the other hand, in order to control its content to less than 0.001%, it takes a long time during the steelmaking operation, resulting in a decrease in productivity. Therefore, in the present invention, it is preferable to limit its content to 0.001 to 0.01%.

[0035] Ti: 0.005 to 0.12% The above-mentioned Ti is a typical precipitation strengthening element together with Nb and V, and forms coarse TiN in the steel with a strong affinity for N. TiN has the effect of suppressing the growth of crystal grains during the heating process for hot rolling. In addition, the remaining Ti after reacting with nitrogen dissolves in the steel and combines with carbon to form TiC precipitates, which are useful components for improving the strength of the steel. However, if the content of Ti is less than 0.005%, the above effects cannot be obtained, and if the Ti content exceeds 0.12%, there is a problem that the elongation flange property is reduced during forming due to the generation of coarse TiN and the coarsening of precipitates. Therefore, in the present invention, it is preferable to limit the Ti content to 0.005 to 0.12%. More preferably, the above Ti content is limited to 0.045 to 0.11%.

[0036] Nb: 0.005 to 0.06% The above-mentioned Nb is a typical precipitation strengthening element together with Ti and V, and is effective in improving the strength and impact toughness of the steel due to the effect of precipitation during hot rolling and the refinement of crystal grains by recrystallization delay. However, if the content of Nb is less than 0.005%, the above effects cannot be obtained, and if the Nb content exceeds 0.06%, there is a problem that the elongation flange property is reduced due to the formation of elongated crystal grains and coarse composite precipitates caused by excessive recrystallization delay during hot rolling. Therefore, in the present invention, it is preferable to limit the Nb content to 0.005 to 0.06%.

[0037] V: 0.005 to 0.2% The above V is a typical precipitation strengthening element together with Nb and Ti. It hardly precipitates during hot rolling and plays a role in forming precipitates after coiling to improve the strength of the steel. Therefore, it is effective for further strength improvement without increasing the deformation resistance and rolling load due to recrystallization delay during hot rolling. In the present invention, in order to obtain such an effect, the content of V is preferably 0.005% or more. However, when its content is excessive, there is a problem that the elongation flange property is reduced due to the formation of coarse precipitates, which is also economically disadvantageous. Therefore, in the present invention, it is preferably limited to 0.2% at the upper limit, and more preferably limited to 0.15%.

[0038] B: 0.0003 - 0.003% When the above B exists in a solid solution state in the steel, it mainly segregates at the grain boundaries, has the effect of stabilizing the grain boundaries and improving the brittleness of the steel, and plays a role in stabilizing the solid solution N and suppressing the formation of coarse AlN nitrides. In addition, it delays the ferrite phase transformation and is effective for the formation of the hard phases, bainite and martensite. In the present invention, the above effects can only be obtained when the B content is 0.0003% or more. When its content exceeds 0.003%, there is a problem that the effect no longer increases and the ductility decreases, resulting in a reduction in formability. Therefore, in the present invention, it is preferably limited to 0.003% at the upper limit, and more preferably limited to 0.002%.

[0039] Except for the above composition, the rest is Fe. However, in the normal manufacturing process, unavoidable impurities may be mixed in from the raw materials or the surrounding environment, so these cannot be excluded. Since these impurities are known to anyone with ordinary knowledge in this technical field, they are not specifically mentioned in this specification in all their details. On the other hand, the addition of effective components other than the above composition is not excluded.

[0040] Relationship 1 On the one hand, in the present invention, in the steel material having the component ranges as described above, it is preferable to control the value of the following relational expression 1 composed of the contents of C, N, S, Nb, Ti and V to be 0.2 or more and 0.6 or less, and more preferably to be 0.3 or more and 0.5 or less.

[0041] When the X value of the following relational expression 1 exceeds 0.6, the formation of precipitates increases and the strength increases somewhat. However, due to recrystallization delay during hot rolling, a fine structure stretched in the rolling direction is likely to be formed, and the elongation rate in the direction perpendicular to rolling decreases. Also, when the hot-rolled steel sheet is cooled, the solid-solution C and solid-solution N atoms in the untransformed phase are insufficient, making it difficult to stably form a hard phase, the grain boundaries become fragile, and the shear surface quality deteriorates. On the other hand, when the X value is less than 0.2, the growth of crystal grains becomes easy during reheating, recrystallization becomes non-uniform during hot rolling, locally coarse crystal grains are formed, the solid-solution C and solid-solution N become excessive more than necessary, and the hardness value of the hard phase tends to be high. Eventually, there is a possibility that the elongation flange property deteriorates.

[0042] [Relational Expression 1] 0.2 ≦ X ≦ 0.6 X = (Nb / 93 + Ti* / 48 + V / 51) / (C / 12 + N / 14) Ti* = Ti - 3.42N - 1.5S

[0043] In the above relational expression 1, Nb, Ti, C, N, and S are the weight percentages of the corresponding alloying elements, and 0 is substituted when not added.

[0044] Relational Expression 2 In the present invention, in the steel material having the component ranges as described above, it is preferable to control the value of the following relational expression 2 composed of Mn, Mo, Cr, and B to be 1.5 or more and 3.5 or less. The following relational expression 2 factors out the combination of alloying elements that can maintain the formation of bainite, martensite, and MA phase, which are hard phases, in the microstructure of the steel of the present invention at an appropriate level. The larger the "T" value of the following relational expression 2, the more the formation of bainite, martensite, and MA phase, which are hard phases, increases, and the hardness values of the respective hard phases also increase. Therefore, the larger this value is, the more advantageous it is for ensuring strength. However, if it becomes excessive, the ductility of the steel decreases, the hardness difference between the soft phase and the hard phase increases, the flange formability of elongation decreases, and there is a problem that the material variation increases also in the entire length and width of the hot-rolled steel sheet.

[0045] [Relational Expression 2] 1.5 ≤ T ≤ 3.5 T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B]

[0046] In the above relational expression 2, Mn, Mo, Cr, and B are the weight percentages of the corresponding alloying elements, and when not added, 0 is substituted.

[0047] Hereinafter, the microstructure of the high-strength hot-rolled steel sheet of the present invention will be described.

[0048] The steel proposed in the present invention has, in its microstructure, a ferrite-based low-temperature transformation phase as the main phase, with a total of less than 40% of bainite phase and polygonal ferrite phase, and a total of less than 3% of the remaining pearlite, martensite, and MA phase, and the average dislocation density (Geometrically Necessary Dislocation) value of the above microstructure is 1.0×10 14 ~2.5×10 14 m -2 satisfies the range.

[0049] When the microstructure of the steel sheet is controlled within the above-mentioned range proposed by the present invention, the target tensile strength of the steel of 590 MPa or more is ensured, and the HER0 (Hole Expansion Ratio of specimen with machined hole) is excellent at 100% or more. The HER value, which is the actual flanging formability, can satisfy that the product of the tensile strength × HER is 45,000 MPa% or more even in a wide range where the punching clearance is 5 - 20%.

[0050] However, among the microstructures composed of the above-mentioned polygonal ferrite phase, bainite phase, and ferrite-based low-temperature transformation phase, when the phase fraction of bainite with a high dislocation density increases and the total with the polygonal ferrite phase exceeds 40%, the elongation rate significantly decreases. Even if HER0 satisfies 100% or more, the dependence of the sheared surface quality on the change in punching clearance increases, and it becomes difficult for the product of the actual tensile strength × HER to satisfy 45,000 MPa% or more. Also, when the phase fraction of polygonal ferrite in the above microstructure excessively increases and the total with the bainite phase exceeds 40%, the strength decreases and it becomes difficult to satisfy the target strength. When the punching clearance is large, the generation of burr becomes intense, HER decreases, and it also becomes difficult for the product of the tensile strength × HER to satisfy 45,000 MPa% or more. More preferably, the total of the above bainite and polygonal ferrite phase fractions is limited to 10 - 40%.

[0051] On the other hand, in the present invention, the ferrite-based low-temperature transformation phase refers to ferrite-based phases and microstructures generated by low-temperature transformation such as acicular ferrite, granular bainitic ferrite, and bainitic ferrite. These ferrite-based low-temperature transformation phases are distinguished from polygonal ferrite in that the shape of the crystal grains is not equiaxed, the grain boundaries are irregular, and the dislocation density is high. They are distinguished from bainite in that no iron carbide precipitates inside or at the boundaries. Also, the above low-temperature transformation phase can include two or more phases and microstructures.

[0052] The present invention does not particularly limit the area ratio of each of the ferrite phase, bainite phase, and ferrite-based low-temperature transformation phase in the microstructure of steel. However, since polygonal ferrite contributes to the ductility of steel and the formation of fine precipitates, it is preferable to limit its area ratio to 10% or more.

[0053] In addition, since bainite is a harder phase than the ferrite-based low-temperature transformation phase, an increase in the bainite phase results in a decrease in ductility. Considering this, it is preferable to limit the area ratio of the bainite phase to 20% or less.

[0054] Moreover, the steel of the present invention can also contain pearlite, martensite, retained austenite, and MA (Martensite and Austenite) in an amount of less than 3 area% as the remaining microstructure other than the above-described microstructure. If the above phases contain 3 area% or more alone or in total, the elongation flange formability of the steel will significantly decrease, and it may be difficult to satisfy the product of tensile strength × HER proposed in the present invention being 45,000 MPa% or more.

[0055] In the present invention, the classification and area fraction of the polygonal ferrite phase, bainite phase, ferrite-based low-temperature transformation phase, martensite phase, retained austenite phase, etc. formed in the steel can be measured using electron backscattered diffraction (EBSD, (JEOL JSM-7001F)), and the area fraction can be measured from the results of analysis at a magnification of 1000 to 3000.

[0056] On the one hand, the above-mentioned ferrite low-temperature transformation phase exhibits a higher hardness value compared to polygonal ferrite. Further, the above bainite phase is a phase containing fine iron carbides and is a hard phase having a higher hardness value compared to the ferrite-based low-temperature transformation product phase. When such a phase is formed in admixture with the soft polygonal ferrite phase, the steel is characterized by exhibiting a stable HER value along with high strength. However, since various physical characteristics including the hardness value in terms of composition vary greatly depending on the components constituting the steel, it has been difficult to clearly distinguish how excellent the original HER value of the steel material is only based on the compositional ratio and whether it exhibits a stable HER value even with fluctuations in the punching clearance.

[0057] Therefore, in the present invention, among the microstructural characteristics of the steel material, it was confirmed that the average dislocation density (Geometrically Necessary Dislocation) of the microstructure is an important influencing factor for the balance between the strength and HER value of the steel material.

[0058] That is, in the present invention, a high-strength hot-rolled steel sheet is provided in which the average dislocation density (Geometrically Necessary Dislocation) of the microstructure satisfies 1.0×10 14 ~2.5×10 14 m -2 . Maintaining the matrix structure of the ferrite-based low-temperature transformation phase and the bainite phase having such a certain dislocation density at a certain level is advantageous for maintaining both high strength, high ductility, and high HER.

[0059] The above average dislocation density (Geometrically Necessary Dislocation) can be calculated using the kernel average misorientation (KAM) data measured by EBSD as shown in the following mathematical formula.

[0060]

Equation

[0061] Here, θ is the average misorientation (KAM values), u is the unit length (step size in the EBSD measurement), b is the Burgers vector.

[0062] Such calculations can be conveniently performed using software such as OIM analysis TM (EDAX) for analyzing the above EBSD measurement results. On the other hand, the EBSD measurement was evaluated based on a cross-section parallel to the rolling direction at the 1 / 4 thickness position of the hot-rolled steel sheet.

[0063] Next, the manufacturing method of the high-strength hot-rolled steel sheet of the present invention will be described in detail.

[0064] The manufacturing method of the hot-rolled steel sheet of the present invention includes the steps of reheating a steel slab satisfying the above-described compositional components and the above relational expressions 1 and 2 to 1150 to 1350 °C; hot-rolling the reheated steel slab at a temperature in the range of 850 to 1150 °C; performing primary cooling on the hot-rolled steel sheet to a temperature in the range of 400 to 500 °C at an average cooling rate of 50 to 100 °C / sec; reheating the primarily cooled steel sheet to a temperature in the range of 450 to 550 °C by maintaining it for 0.5 to 3 seconds; performing secondary cooling on the reheated steel sheet to a temperature in the range of 400 to 500 °C at an average cooling rate of 1 to 30 °C / sec and then winding it up; and cooling the wound coil at an average cooling rate of 0.1 to 25 °C / hour to a temperature in the range of normal temperature to 200 °C. In one or more of the above primary cooling step and the above secondary cooling step, the temperature (T E ) of both edge portions (W×60%) with respect to the entire width (W) of the steel sheet is 450 to 550 °C, and the temperature (T C ) of the central portion of the width (W×40%) is 400 to 500 °C. By completing the cooling in this way, after winding in the above winding step, the average temperature (T A ) of the coil can be maintained in the range of 400 to 500 °C.

[0065] First, in the present invention, a steel slab satisfying the above-described composition components and the above relational expressions 1 and 2 is reheated to 1150 to 1350°C. At this time, if the reheating temperature is less than 1150°C, the precipitates do not fully redissolve, the formation of precipitates decreases in the process after hot rolling, coarse TiN remains, and the slab is not sufficiently heat-matured, making it difficult to control the temperature of the steel sheet constant during hot rolling. On the other hand, if it exceeds 1350°C, the strength decreases due to abnormal grain growth of austenite grains, so the reheating temperature is preferably limited to 1150 to 1350°C.

[0066] Next, in the present invention, the reheated steel slab is hot-rolled at a temperature in the range of 850 to 1150°C.

[0067] The above hot rolling is performed at a temperature in the range of 850 to 1150°C. If hot rolling starts at a temperature higher than 1150°C, the temperature of the hot-rolled steel sheet becomes high, the grain size becomes coarse, and the surface quality of the hot-rolled steel sheet deteriorates. On the other hand, if the above hot rolling ends at a temperature lower than 850°C, the elongated grains due to excessive recrystallization delay develop and the anisotropy becomes intense, and the formability also deteriorates.

[0068] Then, in the present invention, the hot-rolled steel sheet is first cooled at an average cooling rate of 50 to 100°C / s to a temperature in the range of 400 to 500°C.

[0069] As will be described later, in the present invention, after the first cooling is completed, after a short cooling pause period, it will be coiled after the second cooling. Therefore, during the first cooling, if the cooling end temperature becomes too low, there is a possibility that more hard phases than necessary will be formed in the subsequent process. In particular, if it is cooled to less than 400°C, a polygonal ferrite phase is not formed in the microstructure or its area fraction is formed to be less than 10%, and the elongation of the steel is significantly insufficient. On the other hand, if it is cooled above 500°C, fine precipitates are excessively formed, the yield strength increases, the solid solution C and solid solution N necessary for the formation of the hard phase decrease, and the area fraction of the bainite phase decreases, making it difficult to ensure the target strength. Therefore, it is preferable to perform the first cooling to a temperature in the range of 400 to 500°C.

[0070] Also, it is preferable that the average cooling rate during the primary cooling is 50 to 100 °C / s. If the above cooling rate is less than 50 °C / sec, there is a possibility that an excessive amount of polygonal ferrite phase fraction is formed, which is disadvantageous for ensuring strength. If the above cooling rate exceeds 100 °C / sec, the polygonal ferrite phase fraction significantly decreases in the region where the primary cooling end temperature is low, and there is a possibility that the elongation rate is insufficient.

[0071] Next, in the present invention, the steel sheet cooled in the primary cooling is maintained for 0.5 to 3 seconds to reheat the steel sheet to a temperature in the range of 450 to 550 °C.

[0072] That is, the cooling of the steel sheet cooled in the primary cooling is interrupted for 0.5 to 3 seconds so that the temperature of the steel sheet satisfies a temperature in the range of 450 to 550 °C due to internal latent heat and transformation heat generation. At this time, if the cooling interruption time is less than 0.5 seconds, there is no reheating effect. If it exceeds 3 seconds, the polygonal ferrite phase fraction in the microstructure greatly increases, and the bainite and bainitic ferrite phases, which are hard phases, decrease.

[0073] Also, as the temperature of the steel sheet approaches 600 °C, pearlite structure and coarse carbides may be formed in the thickness center of the high-temperature region, and the flange formability of elongation decreases. Conversely, in the region where the temperature of the steel sheet is low and close to 400 °C, excessive cooling occurs and the elongation rate is insufficient.

[0074] Next, in the present invention, after the reheated steel sheet is secondarily cooled at an average cooling rate of 1 to 30 °C / s to a temperature in the range of 400 to 500 °C, it is coiled.

[0075] The end temperature during the secondary cooling is preferably in the range of 400 to 500 °C, and more preferably 430 to 480 °C. If the secondary cooling end temperature is too high, bainite may not be sufficiently formed, and it may be difficult to ensure strength. On the other hand, if it is too low, excessive amounts of bainite phase, martensite phase, and MA phase may be formed, and both the ductility and flange formability of elongation of the steel may decrease.

[0076] Also, the average cooling rate during secondary cooling is preferably 1 to 30 °C / s. If the cooling rate is too high, the MA phase is likely to be formed, and excessive bainite phase is formed, resulting in a decrease in elongation. Although the lower limit of the cooling rate is not particularly limited, in order to slowly control the cooling rate below 1 °C / s, separate cooling and heat preservation equipment, etc. are required, which may be economically disadvantageous. Therefore, considering this, the lower limit can be limited to 1 °C / s.

[0077] On the other hand, in the present invention, in one or more of the above primary cooling step and the above secondary cooling step, the temperature (T E ) of both side edge portions (W×60% of the total width (W) of the steel plate) becomes 450 to 550 °C, and the temperature (T C ) of the width center portion (W×40%) is cooled to 400 to 500 °C at the end of manufacturing, and the average temperature (T A ) of the coil after winding is controlled to be 400 to 500 °C. Here, both side edge portions (W×60%) are the total value of the width W×30% of one side edge portion and the width W×30% of the other side edge portion with respect to the total width (W) of the steel plate in the width direction of the steel plate, and the width center portion (W×40%) indicates the width size of the center portion of the steel plate excluding the width sizes of the above both edge portions.

[0078] At this time, (T E ) and (T C ) mean the position-specific temperatures of the steel plate immediately after the above primary cooling and / or secondary cooling are completed. When the temperature T E of the edge portion becomes excessively low, the MA phase and the martensite phase are likely to be formed in the edge portion of the wound coil with a high cooling rate after winding, and both the elongation and the flangeability of elongation decrease. Conversely, when the temperature T C of the width center portion becomes excessively high, pearlite and coarse carbides are formed in the center portion where the cooling rate of the coil after winding is very slow, and the shearing quality and the flangeability of elongation decrease. Here, the upper limit of T E and the lower limit of T C are not particularly limited, but if T E becomes excessively high or T CIf the temperature is too low, the above-mentioned standard for the secondary cooling end temperature will not be met, so the upper and lower limits can be set to 550°C and 400°C, respectively.

[0079] In the present invention, the steel sheet after the secondary cooling is wound up, and the average temperature of the wound coil (T A ) can be controlled to be 400 to 500°C.

[0080] Thereafter, in the present invention, the wound coil is cooled to a temperature ranging from room temperature to 200°C at an average cooling rate of 0.1 to 25°C / hour.

[0081] In this case, if the cooling rate exceeds 25°C / hour, an MA phase is likely to form in the steel, which reduces the stretch flangeability of the steel, and in order to control the cooling rate to less than 0.1°C / hour, separate heating equipment or the like is required, which is economically disadvantageous. It is preferable to cool at a rate of 1 to 10°C / hour.

[0082] Alternatively, the present invention may further include the steps of pickling and oiling the coiled steel sheet after the secondary cooling.

[0083] The method may further include a step of heating the pickled and oiled steel sheet to a temperature range of 450 to 740° C. and then hot-dip galvanizing the steel sheet.

[0084] In the present invention, the hot dip galvanizing can use a plating bath containing 0.01 to 30% by weight of magnesium (Mg), 0.01 to 50% by weight of aluminum (Al), the balance being Zn and unavoidable impurities. [Example]

[0085] The present invention will be described in more detail below with reference to examples. However, the description of such examples is intended to illustrate the implementation of the present invention, and the present invention is not limited by the description of such examples.

[0086] (Example) Steel slabs having the component compositions as shown in Table 1 below were each prepared. Next, hot-rolled steel sheets were manufactured using the manufacturing conditions shown in Table 2 below for each of the prepared steel slabs. In Table 2 below, FDT represents the hot rolling temperature, CR1st represents the primary cooling rate, Hold Time represents the cooling interruption time after the primary cooling, MT represents the temperature of the steel sheet before the start of the secondary cooling, and CR2nd represents the secondary cooling rate. After the completion of such primary cooling, the average temperature of both edge portions of the steel sheet is T E and the average temperature of the center portion of the width after the completion of cooling is T C and the average temperature of the coiled coil after the completion of the secondary cooling is T A respectively. On the other hand, the coiled coil was cooled at an average cooling rate of 0.1 to 25 °C / hour to a temperature in the range of normal temperature to 200 °C.

[0087]

Table 1

[0088]

Table 2

[0089] The microstructure of each hot-rolled steel sheet manufactured as described above was analyzed in detail and shown in Table 3 below. At this time, the microstructure was analyzed at the position of 1 / 4 to 1 / 2 of the thickness of the cross-section in the rolling direction of the steel sheet. Then, the classification and area fraction measurement of the polygonal ferrite phase (PF), bainite phase (B), ferrite-based low-temperature transformation phase (BF), martensite phase (M), and MA phase formed in the steel sheet were analyzed at a magnification of 3000 to 5000 using Electron Back Scattered Diffraction (EBSD, (JEOL JSM-7001F)). Also, the average dislocation density (Geometrically Necessary Dislocation, GND) was calculated using OIM analysis TM (EDAX) after EBSD measurement based on a cross-section parallel to the rolling direction at the 1 / 4 position of the thickness of the hot-rolled steel sheet. On the other hand, as shown in Table 3, in this example, as a result of observing the microstructure of the hot-rolled steel sheet, no retained austenite was observed at all.

[0090]

Table 3

[0091] On the other hand, Table 4 below shows the evaluation of the mechanical properties of each of the above-manufactured hot-rolled steel sheets and the results thereof.

[0092] The mechanical properties showed the HER results evaluated with different punching clearances of 5%, 10%, and 20% in the tensile test for each steel plate. Here, YS, TS, and T-El represent the 0.2% offset yield strength, tensile strength, and elongation at fracture, respectively, which are the results of testing specimens taken perpendicular to the rolling direction using JIS No. 5 standard test pieces. Also, the HER evaluation was conducted based on the JFST 1001-1996 standard. The results of the HER test in Table 4 below are the averaged values after being carried out three times.

[0093]

Table 4

[0094] As shown in Tables 1-4 above, it can be confirmed that Invention Examples 1-10, where the composition components and manufacturing process conditions of the steel satisfy the scope of the present invention, all have excellent mechanical properties.

[0095] On the other hand, Comparative Examples 1-2 are cases where Relational Expression 1 is not satisfied. In Comparative Example 1, the component configuration exceeded the proposed range of Relational Expression 1, and the yield strength and tensile strength were high due to the precipitation strengthening effect. Also, among the fine structures of the steel, the average dislocation density was 0.9× 14 m -2 which is smaller than the proposed range of the present invention, excessive formation of polygonal ferrite phase occurred, and the HER value further decreased when the punching clearance was large.

[0096] Comparative Example 2 is a case where the component configuration of the steel did not reach the proposed range of Relational Expression 1, excessive formation of martensite phase and MA phase occurred, the average dislocation density was 2.6× 14 m -2 which is high, and the HER value decreased compared to the material strength.

[0097] Comparative Examples 3-4 are cases where the composition of the steel component did not satisfy Relational Expression 2. Comparative Example 3 is a case where the compositional components of the steel exceeded the proposed range of Relational Expression 1. At this time, it excessively contained alloy components with high hardening ability, ensuring stable strength but having insufficient elongation and a decreased HER value. And Comparative Example 4 is a case where the compositional components of the steel did not reach the proposed range of Relational Expression 1, lacking alloy components with high hardening ability and not forming a bainite phase in the microstructure, and unable to ensure the target strength.

[0098] On the other hand, Comparative Examples 5-14 are cases where the compositional components of the steel are within the scope of the present invention, but the conditions of the manufacturing process deviate from the scope of the present invention.

[0099] Specifically, Comparative Examples 5-6 are cases where the finishing temperature during the first cooling after hot rolling deviated from the proposed range of the present invention. In Comparative Example 5, the finishing temperature during the first cooling exceeded the range of the present invention, forming unnecessary pearlite structures and resulting in a decreased HER value. And in Comparative Example 6, the finishing temperature during the first cooling did not reach the range of the present invention, lacking polygonal ferrite phase, and unnecessary MA phases were formed within the crystal grains of the ferrite-based low-temperature transformation phase, and martensite structures were also formed at some grain boundaries, resulting in a decreased HER value.

[0100] Comparative Examples 7-8 are cases where the cooling rate during the first cooling deviated from the proposed range of the present invention. Specifically, in Comparative Example 7, the cooling rate exceeded the proposed range of the present invention, lacking polygonal ferrite phase, and most of the microstructure was composed of ferrite-based low-temperature transformation phases, resulting in a significant decrease in elongation. And in Comparative Steel 8, the cooling rate did not reach the proposed range, the fraction of the polygonal ferrite phase increased significantly, the elongation was good, but the HER decreased when the punching clearance was large.

[0101] Comparative Example 9 is a case where the cooling stop time after the first cooling is longer than necessary. The fraction of the polygonal ferrite phase in the microstructure of the steel plate increased significantly, and unnecessary pearlite structures were also formed, and the HER value decreased both when the punching clearance was small and large.

[0102] Comparative Examples 10 to 11 are cases where the cooling end temperature during secondary cooling is outside the proposed range of the present invention. In Comparative Example 10, the secondary cooling end temperature exceeded the temperature range of the present invention, the polygonal ferrite phase fraction increased significantly, and an unnecessary pearlite structure was also formed, resulting in a decrease in the HER value. And in Comparative Example 11, the temperature of the steel sheet immediately after secondary cooling did not reach the range of the present invention, and martensite and MA phases were unnecessarily formed in the microstructure, resulting in a decrease in the HER value.

[0103] Comparative Example 12 is a case where the cooling rate during secondary cooling exceeded the proposed range of the present invention, and an excessive amount of MA phase was formed in the microstructure, resulting in a decrease in the HER value.

[0104] Comparative Examples 13 - 14 are cases where the temperature in the width direction of the steel sheet did not meet the criteria of the present invention immediately before coiling. Specifically, in Comparative Example 13, the temperature of the edge portion of the steel sheet did not reach the criteria of the present invention, and a low-temperature transformation phase including a martensite phase was excessively formed at the edge portion. Although the strength was high, the elongation rate was insufficient and the HER also decreased. And in Comparative Example 14, the temperature at the center of the steel sheet exceeded the criteria of the present invention, and a pearlite structure was formed at the center of the steel sheet, and the bainite structure also deteriorated and it was difficult to distinguish it from the pearlite structure, resulting in a decrease in the HER characteristics.

[0105] On the other hand, FIG. 1 is a graph showing the relationship between HER and the punching clearance of Invention Examples 1 - 10 and Comparative Examples 5 - 14 in the examples of the present invention. As shown in FIG. 1, in the case of Invention Examples 1 - 10 of the present invention, the tensile strength is 590 MPa or more, and HER0 is excellent at 100% or more compared with Comparative Examples 5 - 14. It can be confirmed that the HER value, which is the actual flanging formability, satisfies that the product of the tensile strength × HER is 45,000 MPa% or more even in a wide range where the punching clearance is 5 - 20%.

[0106] As described above, in the detailed description of the present invention, preferred embodiments of the present invention have been described. However, it goes without saying that those with ordinary knowledge in the technical field to which the present invention pertains can make various modifications without departing from the scope of the present invention. Therefore, the scope of the rights of the present invention should not be defined by being limited to the described embodiments, but should be defined by not only the scope of the claims described later but also equivalents thereof.

Claims

1. By weight, C: 0.03 to 0.08%, Si: 0.01 to 1.0%, Mn: 1.0 to 2.0%, Sol.Al: 0.01 to 0.1%, Cr: 0.005 to 0.5%, Mo: 0.005 to 0.3%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, Ti: 0.005 to 0.12%, Nb: 0.005 to 0.06%, V: 0.005 to 0.2%, B: 0.0003 to 0.003%, the balance being Fe and unavoidable impurities, satisfying the following relational expressions 1 and 2, The microstructure has a ferrite-based low-temperature transformation phase consisting of one or more of acicular ferrite, granular bainitic ferrite, and bainitic ferrite as the main phase, with the total of the bainite phase and the polygonal ferrite phase being less than 40%, and the remaining pearlite, martensite, retained austenite, and MA phase being less than 3% in total, and the average dislocation density (Geometrically Necessary Dislocation) of the microstructure being 1.0×10 １４ ~2.5×10 １４ m －２ A hot-rolled steel sheet satisfying the above requirements. 【Relational expression 1】 0.2 ≤ X ≤ 0.6 X = (Nb / 93 + Ti* / 48 + V / 51) / (C / 12 + N / 14) Ti* = Ti - 3.42N - 1.5S In the above relational expression 1, Nb, Ti, C, N, and S are the weight percentages of the corresponding alloying elements, and 0 is substituted when not added. 【Relational expression 2】 1.5 ≤ T ≤ 3.5 T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] In the above relational expression 2, Mn, Mo, Cr, and B are the weight percentages of the corresponding alloying elements, and 0 is substituted when not added.

2. The hot-rolled steel sheet according to Claim 1, wherein the total area fraction of the bainite phase and the polygonal ferrite phase is 10 to 40%.

3. The hot-rolled steel sheet according to Claim 1, wherein the area fraction of the polygonal ferrite phase is 10% or more.

4. The hot-rolled steel sheet according to Claim 1, wherein the area fraction of the bainite phase is 20% or less.

5. The hot-rolled steel sheet according to Claim 1, having a tensile strength of 590 MPa or more, and the HER value satisfies that the product of the tensile strength × HER is 45,000 MPa% or more in the range where the punching clearance is 5 to 20%.

6. Reheating a steel slab containing, by weight, C: 0.03 to 0.08%, Si: 0.01 to 1.0%, Mn: 1.0 to 2.0%, Sol.Al: 0.01 to 0.1%, Cr: 0.005 to 0.5%, Mo: 0.005 to 0.3%, P: 0.001 to 0.05%, S: 0.001 to 0.01%, N: 0.001 to 0.01%, Ti: 0.005 to 0.12%, Nb: 0.005 to 0.06%, V: 0.005 to 0.2%, B: 0.0003 to 0.003%, the balance being Fe and unavoidable impurities, and satisfying the following relational expressions 1 and 2, to 1150 to 1350 °C; Hot-rolling the reheated steel slab at a temperature in the range of 850 to 1150 °C; The step of performing primary cooling on the hot-rolled steel sheet at an average cooling rate of 50 to 100 °C / sec until the temperature reaches a range of 400 to 500 °C; The step of reheating the steel sheet to a temperature in the range of 450 to 550 °C by maintaining the once-cooled steel sheet for 0.5 to 3 seconds; The step of performing secondary cooling on the reheated steel sheet at an average cooling rate of 1 to 30 °C / sec until the temperature reaches a range of 400 to 500 °C and then winding it up; and The step of cooling the wound coil at an average cooling rate of 0.1 to 25 °C / hour until the temperature reaches a range of normal temperature to 200 °C, In one or more of the primary cooling step and the secondary cooling step, the temperature (T Ｅ ) of both edge portions (W×60%) with respect to the entire width (W) of the steel sheet is 450 to 550°C, and the temperature (T Ｃ ) of the center portion of the width (W×40%) is 400 to 500°C. By completing the cooling in this way, after winding at the winding stage, the average temperature (T Ａ ) of the coil can be maintained in the range of 400 to 500°C. A method for manufacturing a hot-rolled steel sheet. [Relational Expression 1] 0.2 ≤ X ≤ 0.6 X = (Nb / 93 + Ti* / 48 + V / 51) / (C / 12 + N / 14) Ti* = Ti - 3.42N - 1.5S In the above Relational Expression 1, Nb, Ti, C, N, and S are the weight percentages of the corresponding alloying elements, and 0 is substituted when not added. [Relational Expression 2] 1.5 ≤ T ≤ 3.5 T = [Mn] + 2.8[Mo] + 1.5[Cr] + 500[B] In the above Relational Expression 2, Mn, Mo, Cr, and B are the weight percentages of the corresponding alloying elements, and 0 is substituted when not added.

7. The method for manufacturing a hot-rolled steel sheet according to Claim 6, further comprising the step of pickling and oiling the steel sheet wound up after the secondary cooling.

8. The method for manufacturing a hot-rolled steel sheet according to Claim 6, further comprising the step of hot-dip galvanizing after heating the pickled and oiled steel sheet to a temperature range of 450 to 740 °C.

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