Hot-rolled steel sheet and method for producing same

Abstract

The invention provides a hot-rolled steel sheet and a manufacturing method thereof. The hot-rolled steel sheet according to the present invention contains, in wt%, 0.03 to 0.08% of C, 0.01 to 1.0% of Si, 1.0 to 2.0% of Mn, 0.01 to 0.1% of acid-soluble aluminum, 0.005 to 0.5% of Cr, 0.005 to 0.3% of Mo, 0.001 to 0.05% of P, 0.001 to 0.01% of S, 0.001 to 0.01% of N, 0.005 to 0.12% of Ti, 0.005 to 0.06% of Nb, 0.005 to 0.2% of V, 0.0003 to 0.003% of B, and the balance Fe and unavoidable impurities, and satisfies relational expressions 1 to 2.

Description

Technical Field

[0001] The present invention mainly relates to the manufacture of a high-strength hot-rolled steel plate for components, lower arms, reinforcements and connectors of automobile chassis parts, and more specifically, to a high-strength hot-rolled steel plate and a manufacturing method thereof, wherein the tensile strength of the hot-rolled steel plate is above 590 MPa, and the hole expansion ratio of specimen with machined hole (Hole Expanding Ratio of specimen with machined hole, HER 0 ) is excellent and is more than 100%, so that the HER value as the actual hole expansion and flanging formability satisfies that the product of tensile strength × HER is more than 45000 (MPa%) in a wide blanking gap range of 5-20%. Background Art

[0002] In order to improve the formability of existing high-strength hot-rolled steel sheets for chassis parts, it is proposed to make them into dual-phase composite steel with a ferrite-bainite mixed structure as the basic matrix structure to improve the stretch flangeability, or to make them into high-strength high-expansion flanging steel with ferrite phase or bainite phase as the basic matrix structure.

[0003] In Patent Document 1, the steel is held in the ferrite transformation zone for several seconds according to specific cooling conditions immediately after hot rolling, and then coiled at a bainite forming temperature to form bainite. Thus, an attempt is made to simultaneously ensure strength and stretch flangeability by forming the metal structure into a mixed structure of polygonal ferrite and bainite. In addition, Patent Document 2 proposes a high hole expansion flanging steel based on a C-Si-Mn component system with bainitic ferrite and granular bainitic ferrite as the matrix structure, and Patent Document 3 proposes a technology for improving stretch flangeability by manufacturing a steel having a bainite phase of more than 95% and fewer grains extending in the rolling direction. However, Patent Document 1 is a technology that delays the formation of pearlite by using excessive Si to facilitate the formation of a bainite phase, wherein the surface quality is poor, such as red scale defects, and the ferrite phase fraction as a soft phase is more than 90%, so it is not suitable for manufacturing high-strength steel.

[0004] In addition, Patent Document 2 is a technology that uses bainitic ferrite and granular bainitic ferrite, which are low-temperature ferrite phases, as the matrix structure, but in order to ensure further strength, Cu needs to be used, so surface defects and high-temperature brittleness may occur. Ni needs to be added to prevent this problem, so it is economically disadvantageous.

[0005] Furthermore, Patent Document 3 requires high temperature rolling to form a small amount of elongated grains, and requires cooling at an excessively high cooling rate to convert the bainite phase into a matrix structure, which easily results in local cooling rate differences and deterioration in the shape quality of the rolled plate.

[0006] In addition, although alloy components such as Si, Mn, Mo, Cr, Cu, and Ni, which are mainly used to produce the above-mentioned high-strength steel, are effective in improving the strength and stretch flangeability of the above-mentioned hot-rolled steel sheet, if a large amount of alloy components are added to improve the above-mentioned physical properties, the segregation of alloy components and the inhomogeneity of microstructures will occur, resulting in deterioration of stretch flangeability. In particular, high hardenability steel is sensitive to changes in microstructure during cooling, forming an inhomogeneous low-temperature phase transformation structure phase, and therefore there is a problem that it is difficult to obtain higher stretch flangeability.

[0007] [Prior art literature]

[0008] [Patent Document]

[0009] (Patent Document 1) Japanese Patent Publication No. 1994-293910

[0010] (Patent Document 2) Korean Patent No. 10-1114672

[0011] (Patent Document 3) Korean Patent Publication No. 2013-0080038 Summary of the invention

[0012] 1. Technical issues to be resolved

[0013] The object of the present invention is to provide a hot-rolled steel plate and a method for manufacturing the same, wherein the hot-rolled steel plate has a tensile strength of 590 MPa or more, and a HER 0 The HER value is excellent and is 100% or more, so that the product of tensile strength×HER is 45000 (MPa%) or more in a wide blanking gap range of 5-20% as the actual hole expansion and flanging formability.

[0014] In addition, the technical problems to be solved by the present invention are not limited to the technical problems mentioned above, and those skilled in the art in the technical field to which the present invention belongs can clearly understand other technical problems not mentioned through the following description.

[0015] (II) Technical solution

[0016] One aspect of the present invention relates to a hot-rolled steel plate, which comprises, by weight, 0.03-0.08% C, 0.01-1.0% Si, 1.0-2.0% Mn, 0.01-0.1% acid-soluble aluminum (Sol.Al), 0.005-0.5% Cr, 0.005-0.3% Mo, 0.001-0.05% P, 0.001-0.01% S, 0.001-0.01% N. 01-0.01%, Ti: 0.005-0.12%, Nb: 0.005-0.06%, V: 0.005-0.2%, B: 0.0003-0.003%, the balance is Fe and inevitable impurities, and the hot rolled steel sheet satisfies the following equations 1 and 2, the main phase of the microstructure is a ferrite system low temperature phase transformation phase composed of one or more of acicular ferrite, granular bainitic ferrite and bainitic ferrite, and the microstructure contains less than 40% of bainite phase and polygonal ferrite phase in total and less than 3% of retained pearlite, martensite, retained austenite and MA phase in total, and the average dislocation density (geometric necessary dislocation) of the microstructure satisfies 1.0×10 14 -2.5×10 14 m -2 .

[0017] [Equation 1]

[0018] 0.2≤X≤0.6

[0019] X=(Nb / 93+Ti* / 48+V / 51) / (C / 12+N / 14)

[0020] Ti*=Ti-3.42N-1.5S

[0021] In the above-mentioned relational formula 1, Nb, Ti, C, N and S are the weight % of the corresponding alloying elements, and are entered as 0 when not added.

[0022] [Equation 2]

[0023] 1.5≤T≤3.5

[0024] T=[Mn]+2.8[Mo]+1.5[Cr]+500[B]

[0025] In the above relational expression 2, Mn, Mo, Cr, and B are the weight % of the corresponding alloying elements, and are entered as 0 when not added.

[0026] The tensile strength of the hot-rolled steel sheet may be 590 MPa or more, and the HER value may satisfy that the product of tensile strength×HER is 45000 MPa% or more within a range of a blanking gap of 5-20%.

[0027] In addition, another aspect of the present invention relates to a method for manufacturing a hot-rolled steel plate, the method comprising the following steps: reheating a steel billet to 1150-1350° C., wherein the steel billet comprises, by weight%, C: 0.03-0.08%, Si: 0.01-1.0%, Mn: 1.0-2.0%, acid-soluble aluminum: 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 of Fe and unavoidable impurities, and the steel billet satisfies the following equations 1 and 2;

[0028] hot rolling the reheated steel billet at a temperature in the range of 850-1150° C.;

[0029] Cooling the hot-rolled steel plate at an average cooling rate of 50-100° C. / s to a temperature in the range of 400-500° C.;

[0030] Reheating the steel plate to a temperature in the range of 450-550° C. by holding the steel plate after the primary cooling for 0.5-3 seconds;

[0031] Secondarily cooling the reheated steel sheet at an average cooling rate of 1-30° C. / s to a temperature in the range of 400-500° C., and then coiling; and

[0032] The coiled sheet is cooled to a temperature ranging from room temperature to 200° C. at an average cooling rate of 0.1-25° C. / hour,

[0033] In one or more of the primary cooling process and the secondary cooling process, the temperature (T E ) becomes 450-550℃ and the temperature (T C ) becomes 400-500°C, so that the average temperature (T A ) can be maintained in the range of 400-500°C.

[0034] [Equation 1]

[0035] 0.2≤X≤0.6

[0036] X=(Nb / 93+Ti* / 48+V / 51) / (C / 12+N / 14)

[0037] Ti*=Ti-3.42N-1.5S

[0038] In the above-mentioned relational formula 1, Nb, Ti, C, N and S are the weight % of the corresponding alloying elements, and are entered as 0 when not added.

[0039] [Equation 2]

[0040] 1.5≤T≤3.5

[0041] T=[Mn]+2.8[Mo]+1.5[Cr]+500[B]

[0042] In the above relational expression 2, Mn, Mo, Cr, and B are the weight % of the corresponding alloying elements, and are entered as 0 when not added.

[0043] (III) Beneficial effects

[0044] The present invention having the above-mentioned structure can effectively provide a high-strength steel sheet having a tensile strength of 590 MPa or more by controlling the fine structure and average dislocation density of the steel. 0 The HER value is excellent and is 100% or more, so that the product of tensile strength×HER is 45000 MPa% or more in a wide blanking gap range of 5-20% as the actual hole expansion and flange forming property. BRIEF DESCRIPTION OF THE DRAWINGS

[0045] Figure 1 It is a graph showing the relationship of HER according to the punching clearance of Inventive Examples 1 to 10 and Comparative Examples 5 to 14 in the examples of the present invention.

[0046] Best Mode for Carrying Out the Invention

[0047] Hereinafter, the present invention will be described.

[0048] In order to solve the problems of the prior art described above, the inventors of the present invention studied the changes in local formability and hole expansion and flanging properties (hole expansion ratio (HER)) of various steels with different compositions and microstructures according to the composition, manufacturing process and microstructure characteristics. As a result, it was confirmed that the detailed characteristics of each component phase are correlated with the local formability and hole expansion and flanging properties, thereby obtaining the following Relationships 1 and 2. In addition, based on the obtained relationship, it was confirmed that the main phase of the steel microstructure is a ferrite low-temperature phase transformation phase, and contains a bainite phase and a polygonal ferrite phase with a total of less than 40% and a residual pearlite, martensite and MA phase with a total of less than 3%, the polygonal ferrite is more than 10%, and the average dislocation density (geometric required dislocation) value satisfies 1.0×10 14 -2.5×10 14 m -2 When the tensile strength is above 590MPa, HER 0 The present invention is proposed because the HER value as the actual hole expansion and flanging formability is excellent and is 100% or more, so that the product of tensile strength×HER is 45000 MPa% or more in a wide blanking gap range of 5-20%.

[0049] Hereinafter, the components of the hot-rolled steel sheet of the present invention and the reasons for limiting their contents will be described, wherein "%" means weight % unless otherwise specified.

[0050] C: 0.03-0.08%

[0051] The C is the most economical and effective element for strengthening steel, and has a great influence on the hardness value of each component phase. When the addition amount of the C increases, the hardenability increases, and the phase fraction of hard phases such as bainite phase and martensite phase in the fine structure increases, thereby increasing the tensile strength. In addition, C forms fine precipitates together with Ti and Nb with high affinity, and the yield strength and tensile strength are increased through precipitation strengthening. However, when the C content is less than 0.03%, it is difficult to obtain a sufficient strengthening effect. When the C content exceeds 0.08%, the hardness of each phase including the bainite phase and the martensite phase increases, and there are problems of excessive strength increase and reduced formability, and weldability will also deteriorate. Therefore, the C content is preferably in the range of 0.03-0.08%. In the present invention, in order to stably ensure the target level of strength and formability, the C content is more preferably limited to 0.04-0.07%.

[0052] Si: 0.01-1.0%,

[0053] The Si has the effect of deoxidizing the molten steel and strengthening the solid solution, and is beneficial to improving the formability by delaying the formation of coarse carbides. In the present invention, in order to obtain such an effect, it is preferred to contain 0.01% or more of Si. However, when the content of the Si exceeds 1.0%, red oxide scale caused by Si is formed on the surface of the steel sheet during hot rolling, which not only significantly deteriorates the surface quality of the steel sheet, but also reduces the ductility and weldability. Therefore, it is preferred to limit the content of Si to less than 1.0%.

[0054] Mn: 1.0-2.0%

[0055] Like Si, Mn is an effective element for solid solution strengthening steel. By increasing the hardenability of steel, bainite and martensite phases as hard phases are easily formed during cooling after hot rolling. However, when the content of Mn is less than 1.0%, the above-mentioned effect of adding Mn cannot be obtained. When the content of Mn exceeds 2.0%, the hardenability is greatly increased, and the hardness of each phase including the bainite phase and the martensite phase increases, and there are problems of excessive strength increase and reduced formability. In addition, when the slab is cast in the continuous casting process, the segregation part in the center of the thickness is greatly developed, and an uneven fine structure is formed in the thickness direction during cooling after hot rolling, resulting in poor stretch flangeability. In particular, it is difficult to uniformly produce a fine structure over the entire length and width of the hot-rolled plate during cooling. More preferably, the upper limit of the Mn content is limited to 1.8%.

[0056] Acid soluble aluminum: 0.01-0.1%,

[0057] The acid-soluble aluminum is mainly added for deoxidation. When the content of the acid-soluble aluminum is less than 0.01%, the effect of adding the acid-soluble aluminum is insufficient. When the content of the acid-soluble aluminum exceeds 0.1%, the acid-soluble aluminum combines with nitrogen to form AlN, which is easy to cause corner cracks in the slab during continuous casting, and defects caused by the formation of inclusions are easy to occur. Therefore, it is preferred to limit the content of the acid-soluble aluminum to 0.01-0.1%, and the upper limit of the amount of the acid-soluble aluminum added is more preferably limited to 0.06%.

[0058] Cr: 0.005-0.5%,

[0059] The Cr plays a role in solid solution strengthening of steel and delays ferrite transformation during cooling to help form bainite. However, when the amount of Cr added is less than 0.005%, the above-mentioned effect of adding Cr cannot be obtained, and when the amount of Cr added exceeds 0.5%, the ferrite transformation is excessively delayed to form a martensite phase, so the elongation deteriorates. In addition, similar to Mn, the segregation part in the center of the thickness is significantly developed, resulting in uneven microstructure in the thickness direction, so the stretch flangeability deteriorates. More preferably, the upper limit of the Cr content is limited to 0.3%.

[0060] Mo: 0.005-0.3%

[0061] Mo increases the hardenability of steel and easily forms a bainite structure. However, when the amount of Mo added is less than 0.005%, the above-mentioned effect of adding Mo cannot be obtained. When the amount of Mo added exceeds 0.3%, the martensite phase is formed due to the excessive increase in hardenability, resulting in a sharp deterioration in formability. In addition, it is economically disadvantageous and also harmful to weldability. Therefore, the content of Mo is preferably limited to 0.005-0.3%. More preferably, the content of Mo is limited to 0.05-0.2%.

[0062] P: 0.001-0.05%

[0063] Like Si, P has both solid solution strengthening and ferrite transformation promoting effects. However, in order to control the P content to less than 0.001%, more manufacturing costs are required, which is economically disadvantageous and insufficient to obtain strength. On the other hand, when the P content exceeds 0.05%, brittleness caused by grain boundary segregation occurs, and fine cracks are easily generated during molding, which significantly deteriorates ductility, stretch flangeability and impact resistance. Therefore, the P content is preferably limited to 0.001-0.05%.

[0064] S: 0.001-0.01%

[0065] The S is an impurity present in steel. When the S content exceeds 0.01%, it combines with Mn and the like to form non-metallic inclusions, which easily causes fine cracks during steel cutting, and there is a problem of greatly reducing stretch flangeability and impact resistance. In the present invention, the lower limit of the S content is not particularly limited, but in order to manufacture with an S content of less than 0.001%, a lot of time is required during steelmaking operations, thereby reducing productivity. Therefore, considering this situation, the lower limit of the S content can be limited to 0.001%. Therefore, in the present invention, the S content is preferably limited to 0.001-0.01%, and the upper limit of the S content is more preferably limited to 0.005%.

[0066] N: 0.001-0.01%

[0067] Together with C, N is a representative solid solution strengthening element, and forms coarse precipitates together with Ti, Al, etc. Generally, the solid solution strengthening effect of N is better than that of carbon, but as the amount of N in steel increases, there is a problem of significantly reduced toughness. In addition, in order to control the content of N to less than 0.001%, a lot of time is required during steelmaking operations, thereby reducing productivity. Therefore, in the present invention, it is preferred to limit the content of N to 0.001-0.01%.

[0068] Ti: 0.005-0.12%

[0069] Together with Nb and V, Ti is a representative precipitation strengthening element, and due to its strong affinity with N, coarse TiN is formed in the steel. TiN has the effect of inhibiting grain growth during the heating process for hot rolling. In addition, the remaining Ti after reacting with nitrogen is dissolved in the steel and combined with carbon to form TiC precipitates, so Ti is a useful component for improving the strength of the steel. However, when the Ti content is less than 0.005%, the above-mentioned effect cannot be obtained. When the Ti content exceeds 0.12%, coarse TiN and coarsening of the precipitates are generated, and there is a problem of deterioration of the stretch flangeability during forming. Therefore, in the present invention, the Ti content is preferably limited to 0.005-0.12%. More preferably, the Ti content is limited to 0.045-0.11%.

[0070] Nb: 0.005-0.06%

[0071] Together with Ti and V, Nb is a representative precipitation strengthening element, which precipitates during hot rolling and effectively improves the strength and impact toughness of steel by delaying recrystallization to bring about grain refinement. However, when the Nb content is less than 0.005%, the above effect cannot be obtained, and when the Nb content exceeds 0.06%, during hot rolling, due to excessive delay of recrystallization, extended grains and coarse composite precipitates are formed, so there is a problem of deterioration of stretch flangeability. Therefore, in the present invention, the Nb content is preferably limited to 0.005-0.06%.

[0072] V: 0.005-0.2%

[0073] Together with Nb and Ti, V is a representative precipitation strengthening element. It hardly precipitates during hot rolling, but forms precipitates after coiling, thereby playing a role in improving the strength of steel. Therefore, the strength is effectively further improved without the increase in deformation resistance and rolling load caused by delayed recrystallization during hot rolling. In the present invention, in order to obtain this effect, the V content is preferably 0.005% or more. However, when the V content is too much, there is a problem of deterioration of stretch flangeability due to the formation of coarse precipitates, and it is also economically disadvantageous. Therefore, in the present invention, the upper limit of V is preferably limited to 0.2%, and more preferably limited to 0.15%.

[0074] B: 0.0003-0.003%

[0075] When the B exists in the steel in a solid solution state, it is mainly segregated at the grain boundaries, which has the effect of stabilizing the grain boundaries and improving the brittleness of the steel, and stabilizes the solid solution N to inhibit the formation of coarse AlN nitrides. In addition, by delaying the ferrite phase transformation, the bainite phase and the martensite phase as hard phases are effectively formed. In the present invention, the above-mentioned effect is only achieved when the B content is 0.0003% or more. When the B content exceeds 0.003%, the effect no longer increases, and the ductility is reduced, and there is a problem of poor formability. Therefore, in the present invention, the upper limit of the B content is preferably limited to 0.003%, and more preferably limited to 0.002%.

[0076] In addition to the above composition, the rest is Fe. However, in the usual manufacturing process, unnecessary impurities may inevitably be mixed from the raw materials or the surrounding environment, so these impurities cannot be excluded. These impurities are well known to those skilled in the art, so all of them are not particularly mentioned in this specification. In addition, the addition of effective components other than the above composition is not excluded.

[0077] Relation 1

[0078] In addition, in the present invention, in the steel material having the above-mentioned composition range, the value of the following relational expression 1 composed of the contents of C, N, S, Nb, Ti and V is preferably controlled to be greater than 0.2 and less than 0.6, and more preferably controlled to be greater than 0.3 and less than 0.5.

[0079] When the value of X in the following relational expression 1 exceeds 0.6, the formation of precipitates increases and the strength increases slightly, but during the hot rolling process, due to the delay of recrystallization, a fine structure extending in the rolling direction is easily formed, so the elongation in the vertical direction of rolling decreases. In addition, when the hot-rolled steel plate is cooled, the solid solution C and solid solution N atoms in the untransformed phase become insufficient, it is difficult to stably form a hard phase, and the grain boundaries become fragile, resulting in deterioration of the shear surface quality. On the other hand, when the value of X is less than 0.2, grain growth is easy during reheating, and recrystallization becomes uneven during hot rolling, coarse grains are formed locally, and the solid solution C and solid solution N become unnecessarily excessive. The hardness value of the hard phase shows a tendency to increase, and the stretched flangeability may eventually deteriorate.

[0080] [Equation 1]

[0081] 0.2≤X≤0.6

[0082] X=(Nb / 93+Ti* / 48+V / 51) / (C / 12+N / 14)

[0083] Ti*=Ti-3.42N-1.5S

[0084] In the above-mentioned relational formula 1, Nb, Ti, C, N and S are the weight % of the corresponding alloying elements, and are entered as 0 when not added.

[0085] Relation 2

[0086] Furthermore, in the present invention, in the steel material having the composition range as described above, the value of the following relational expression 2 composed of Mn, Mo, Cr and B is preferably controlled to be greater than 1.5 and less than 3.5. The following relational expression 2 is a factorization of the combination of alloy elements that can maintain the formation of bainite, martensite and MA phases as hard phases at an appropriate level in the fine structure of the steel of the present invention. As the "T" value of the following relational expression 2 increases, the formation of bainite, martensite and MA phases as hard phases increases, and the hardness value of each hard phase also increases. Therefore, the larger the value, the more conducive it is to ensure strength, but if the T value is too large, the ductility of the steel decreases, the hardness difference between the soft phase and the hard phase increases, resulting in poor stretch flangeability, and there is a problem of increased material deviation over the entire length and width of the hot-rolled steel plate.

[0087] [Equation 2]

[0088] 1.5≤T≤3.5

[0089] T=[Mn]+2.8[Mo]+1.5[Cr]+500[B]

[0090] In the above relational expression 2, Mn, Mo, Cr, and B are the weight % of the corresponding alloying elements, and are entered as 0 when not added.

[0091] Hereinafter, the microstructure of the high-strength hot-rolled steel sheet of the present invention will be described.

[0092] The main phase of the microstructure of the steel proposed in the present invention is a ferrite system low temperature phase transformation phase, and contains a total of less than 40% of bainite phase and polygonal ferrite phase and less than 3% of residual pearlite, martensite and MA phase, and the average dislocation density (geometric required dislocation) value of the microstructure satisfies 1.0×10 14 -2.5×10 14 m -2 range.

[0093] When the steel plate microstructure is controlled within the above range proposed in the present invention, the desired tensile strength of the steel of 590 MPa or more can be ensured, and the sample hole expansion ratio (HER) of the processed hole can be 0 ) is excellent and is more than 100%, so that the HER value as the actual hole expansion and flanging formability can meet the product of tensile strength × HER of more than 45000 MPa% in a wide blanking gap range of 5-20%.

[0094] However, in the microstructure composed of the polygonal ferrite phase, the bainite phase and the ferrite system low temperature phase transformation phase, when the phase fraction of the bainite with high dislocation density increases and the sum of the bainite and polygonal ferrite phases exceeds 40%, the elongation decreases significantly, even if HER 0 If the ratio is 100% or more, the dependence of the shear surface quality on the change of the blanking gap becomes large, so it is difficult for the actual tensile strength × HER product to meet 45000 MPa% or more. In addition, when the phase fraction of polygonal ferrite in the microstructure increases excessively and the sum of polygonal ferrite and bainite phases exceeds 40%, the strength decreases and it is difficult to meet the desired strength. When the blanking gap is large, the generation of burrs becomes serious and the HER deteriorates, so it is difficult for the tensile strength × HER product to meet 45000 MPa% or more. More preferably, the sum of the fractions of the bainite and polygonal ferrite phases is limited to 10-40%.

[0095] In addition, in the present invention, the ferrite system low temperature phase transformation phase refers to the phase and organization based on ferrite generated by low temperature phase transformation such as needle-shaped ferrite, granular bainitic ferrite or bainitic ferrite. These ferrite system low temperature phase transformation phases are distinguished from polygonal ferrite in that the shape of the grains is not equiaxed, the grain boundaries are irregular, and the dislocation density is high, and the ferrite system low temperature phase transformation phase is distinguished from bainite in that iron carbides are not precipitated inside or on the grain boundaries. In addition, the above-mentioned low temperature phase transformation phase can also include two or more phases and organizations.

[0096] The present invention does not particularly limit the area ratio of each of the ferrite phase, bainite phase and ferrite low temperature transformation phase in the steel microstructure. However, polygonal ferrite contributes to the ductility of the steel and the formation of fine precipitates, so it is preferred to limit the area ratio of polygonal ferrite to 10% or more.

[0097] Furthermore, since bainite is a harder phase than the ferrite low temperature transformation phase, the increase of the bainite phase leads to a decrease in ductility. Therefore, in consideration of this, the area ratio of the bainite phase is preferably limited to 20% or less.

[0098] In addition to the above structures, the steel of the present invention may contain less than 3 area % of pearlite, martensite, retained austenite and martensite and austenite (MA) as residual structures. When the content of the above phases alone or in sum is 3 area % or more, the stretch flangeability of the steel is significantly deteriorated, and thus it may be difficult to satisfy the requirement of the present invention that the product of tensile strength × HER is 45000 MPa % or more.

[0099] In addition, in the present invention, the distinction and area fraction of the polygonal ferrite phase, bainite phase, ferrite low temperature phase transformation phase, martensite phase and retained austenite phase formed in the steel can be measured using Electron Back Scattered Diffraction (EBSD, (JEOL JSM-7001F)), and the area fraction can be measured based on the results of analysis at a magnification of 1000 to 3000.

[0100] In addition, the above-mentioned ferrite low temperature phase transformation phase shows a higher hardness value than polygonal ferrite. In addition, the bainite phase is a phase containing fine iron carbides, and is a hard phase having a higher hardness value than the ferrite system low temperature phase transformation phase. When the above-mentioned phase is mixed and formed with the soft polygonal ferrite phase, the steel has the characteristics of showing high strength and stable HER value. However, the physical characteristics including the hardness value of each constituent phase vary greatly depending on the components constituting the steel. Therefore, it is difficult to clearly distinguish how excellent the original HER value of the steel is, and whether it shows a stable HER value when the blanking gap changes, only by the proportion of the constituent phases.

[0101] Therefore, in the present invention, it is confirmed that the average dislocation density (geometric required dislocation) of the microstructure of the steel material is an important factor affecting the balance between the strength and HER value of the steel material among the microstructure characteristics of the steel material.

[0102] That is, in the present invention, a microstructure having an average dislocation density (geometric required dislocation) satisfying 1.0×10 14 -2.5×10 14 m -2 High-strength hot-rolled steel sheet. Maintaining the ferrite system low-temperature transformation phase matrix structure and bainite phase with a certain dislocation density at a specified level is beneficial to maintaining high strength, high ductility and high HER.

[0103] The average dislocation density (geometric required dislocation) can be calculated as follows using kernel average misorientation (KAM) data measured by EBSD.

[0104]

[0105] where θ is the average misorientation (KAM values),

[0106] u is the unit length (step size in the EBSD measurement),

[0107] b is the Burgers vector.

[0108] For convenience, the calculations described above can be performed using OIM analysis software, which is software for analyzing the EBSD measurement results. TM (EDAX), etc. EBSD measurement was evaluated based on a cross section parallel to the rolling direction at a position of 1 / 4 of the thickness of the hot-rolled steel sheet.

[0109] Next, the method for producing a high-strength hot-rolled steel sheet of the present invention will be described in detail.

[0110] The method for manufacturing a hot-rolled steel plate of the present invention comprises the following steps: reheating a steel billet satisfying the above composition and the above relational expressions 1 and 2 to 1150-1350° C.; hot rolling the reheated steel billet at a temperature in the range of 850-1150° C.; cooling the hot-rolled steel plate once at an average cooling rate of 50-100° C. / second to a temperature in the range of 400-500° C.; and reheating the steel plate to a temperature in the range of 450-550° C. by holding the once-cooled steel plate for 0.5-3 seconds. ℃; cooling the reheated steel plate to a temperature in the range of 400-500 ℃ at an average cooling rate of 1-30 ℃ / second, and then coiling; and cooling the coiled plate to a temperature in the range of room temperature to 200 ℃ at an average cooling rate of 0.1-25 ℃ / hour, wherein in one or more of the primary cooling process and the secondary cooling process, by making the temperature (T E ) becomes 450-550℃ and the temperature (T C ) becomes 400-500°C, so that the average temperature (T A ) can be maintained in the range of 400-500°C.

[0111] First, in the present invention, the steel slab satisfying the above composition and the above relational expressions 1 and 2 is reheated to 1150-1350°C. At this time, when the reheating temperature is lower than 1150°C, the precipitates are not fully re-solidified, so that the formation of precipitates in the process after hot rolling is reduced and coarse TiN remains, and the heat storage of the slab is insufficient, so it is difficult to constantly control the temperature of the steel plate during hot rolling. In addition, when the reheating temperature exceeds 1350°C, the strength decreases due to abnormal grain growth of austenite grains, so the reheating temperature is preferably limited to 1150-1350°C.

[0112] Next, in the present invention, the reheated steel slab is hot rolled at a temperature in the range of 850-1150°C.

[0113] The hot rolling is performed at a temperature in the range of 850-1150°C. When the hot rolling is started at a temperature higher than 1150°C, the temperature of the hot rolled steel sheet increases, the grain size becomes coarse, and the surface quality of the hot rolled steel sheet deteriorates. On the other hand, when the hot rolling is terminated at a temperature lower than 850°C, due to excessive delay of recrystallization, elongated grains develop, anisotropy becomes severe, and formability also deteriorates.

[0114] Furthermore, in the present invention, the hot-rolled steel sheet is primarily cooled at an average cooling rate of 50-100°C / sec to a temperature in the range of 400-500°C.

[0115] As described below, in the present invention, coiling is performed after a short cooling rest period and secondary cooling after the termination of the primary cooling. Therefore, when the cooling termination temperature is too low during the primary cooling, unnecessary hard phases may be formed in the subsequent process. In particular, when cooled to below 400°C, a polygonal ferrite phase is not formed in the fine structure, or the area fraction of the formed polygonal ferrite phase is less than 10%, so the elongation of the steel is significantly insufficient. On the other hand, when cooled to more than 500°C, too many fine precipitates are formed, the yield strength increases, and the solid solution C and solid solution N required for the formation of the hard phase are reduced, so that the area fraction of the bainite phase is reduced, and it is difficult to ensure the target strength. Therefore, it is preferred to cool to a temperature of 400-500°C for the first time.

[0116] In addition, the average cooling rate during the primary cooling is preferably set to 50-100°C / sec. When the cooling rate is less than 50°C / sec, the fraction of the polygonal ferrite phase formed may be too large, which is not conducive to ensuring strength. When the cooling rate exceeds 100°C / sec, the fraction of the polygonal ferrite phase is greatly reduced in the region where the primary cooling termination temperature is low, so the elongation may be insufficient.

[0117] Next, in the present invention, the steel sheet is reheated to a temperature in the range of 450-550° C. by holding the primary cooled steel sheet for 0.5-3 seconds.

[0118] That is, the steel plate that has been cooled once is stopped for 0.5-3 seconds, and the temperature of the steel plate is adjusted to a temperature in the range of 450-550° C. by internal latent heat and phase change heat. At this time, when the cooling stop time is less than 0.5 seconds, there is no reheating effect, and when the cooling stop time exceeds 3 seconds, the polygonal ferrite phase fraction in the fine structure is greatly increased, and the bainite and bainitic ferrite phases as hard phases are reduced.

[0119] In addition, pearlite and coarse carbides may form in the thickness center of the steel plate at a high temperature region near 600°C, thereby deteriorating stretch flangeability. On the other hand, excessive cooling occurs in the region near 400°C, where the steel plate temperature is low, resulting in insufficient elongation.

[0120] Next, in the present invention, the reheated steel sheet is secondarily cooled at an average cooling rate of 1-30° C. / sec to a temperature in the range of 400-500° C., and then coiled.

[0121] The termination temperature during secondary cooling is preferably in the range of 400-500° C., more preferably 430-480° C. When the secondary cooling termination temperature is too high, bainite is not sufficiently formed, making it difficult to ensure strength, while when the secondary cooling termination temperature is too low, a large amount of bainite phase, martensite phase, and MA phase are unnecessarily formed, so that the ductility and stretch flangeability of the steel are deteriorated.

[0122] In addition, the average cooling rate during secondary cooling is preferably 1-30°C / sec. When the cooling rate is too high, the MA phase is easily formed and too much bainite phase is formed, so the elongation decreases. The lower limit of the cooling rate is not particularly limited, but in order to control the cooling rate to be slow and less than 1°C / sec, separate cooling and heat preservation equipment is required, which may be economically disadvantageous. Therefore, considering this point, the lower limit of the cooling rate can be limited to 1°C / sec.

[0123] In the present invention, in one or more of the primary cooling process and the secondary cooling process, the temperature (T E ) becomes 450-550℃ and the temperature (T C ) is 400-500℃ to terminate the cooling, and the average temperature of the coil after coiling (T A ) is controlled to be 400-500° C. The edge portions on both sides (W×60%) are the sum of the widths W×30% of the edge portions on one side and W×30% of the edge portions on the other side relative to the entire width (W) of the steel plate in the width direction of the steel plate, and the width center portion (W×40%) is the width dimension of the center portion of the steel plate excluding the width dimensions of the two edge portions.

[0124] At this time, (T E ) and (T C ) refers to the temperature of each position of the steel plate after the primary cooling and / or secondary cooling is completed. When T is used as the edge temperature EIf the temperature is too low, MA phase and martensite phase are easily formed in the edge of the coiled sheet where the cooling rate is fast after coiling, so the elongation and stretch flangeability are deteriorated. C If it is too high, pearlite and coarse carbides will form in the center of the coil where the cooling rate is very slow after coiling, so the shear quality and stretch flangeability will deteriorate. E The upper limit and T C The lower limit of T is not particularly limited, but when T E Too high or T C When it is too low, the above secondary cooling termination temperature standard cannot be met, so T E and T C The upper and lower limits are defined as 550°C and 400°C, respectively.

[0125] In addition, in the present invention, the steel plate after the secondary cooling is terminated is coiled. At this time, the average temperature (T A ) can be controlled to 400-500℃.

[0126] Thereafter, in the present invention, the coiled sheet is cooled to a temperature ranging from room temperature to 200° C. at an average cooling rate of 0.1 to 25° C. / hour.

[0127] At this time, when the cooling rate exceeds 25°C / hour, the MA phase is easily formed in the steel, and the stretch flangeability of the steel deteriorates. In addition, in order to control the cooling rate to less than 0.1°C / hour, a separate heating device is required, which is economically disadvantageous. It is preferably cooled at 1-10°C / hour.

[0128] Alternatively, the present invention may further include the steps of pickling and oiling the steel plate that has been coiled after the secondary cooling.

[0129] In addition, the method may further include heating the pickled and oiled steel sheet to a temperature range of 450-740° C. and then hot-dip galvanizing.

[0130] In the present invention, the hot-dip galvanizing may use a plating bath containing 0.01-30 wt % of magnesium (Mg), 0.01-50 wt % of aluminum (Al), and the balance of Zn and inevitable impurities. DETAILED DESCRIPTION

[0131] Hereinafter, the present invention will be described in more detail by way of examples. However, the description of these examples is only for illustrating the practice of the present invention, and the present invention is not limited to the description of these examples.

[0132] (Example)

[0133] Steel slabs having the composition shown in Table 1 below were prepared. Then, the prepared steel slabs were made into hot-rolled steel sheets using the manufacturing conditions shown in Table 2 below. In Table 2 below, FDT represents the hot rolling temperature, the first CR (CR1st) represents the first cooling rate, the hold time (Hold Time) represents the cooling interruption time after the first cooling, MT represents the temperature of the steel sheet before the start of the second cooling, and the second CR (CR2nd) represents the second cooling rate. After the first cooling as described above is terminated, the average temperature of the two edge portions of the steel sheet is represented as T E The average temperature of the width center after cooling is expressed as T C , the average temperature of the coil after the secondary cooling is terminated is expressed as T A In addition, the coiled sheet is cooled to a temperature ranging from room temperature to 200° C. at an average cooling rate of 0.1 to 25° C. / hour.

[0134] [Table 1]

[0135]

[0136] * The unit of the content of each component in Table 1 is weight %, and the remaining components are Fe and inevitable impurities. In Table 1, X represents the X value defined by Relational Formula 1, and T represents the T value defined by Relational Formula 2.

[0137] [Table 2]

[0138] category Steel Type FDT One CR Keep Time MT Secondary CR TE Tc TA Comparative Example 1 Comparison Steel 1 885 68 0.7 465 15 490 448 470 Comparative Example 2 Comparison Steel 2 862 65 0.7 470 18 485 432 450 Comparative Example 3 Comparison Steel 3 865 70 0.7 443 20 465 430 442 Comparative Example 4 Comparison Steel 4 860 69 0.7 477 22 472 443 463 Comparative Example 5 Inventing Steel 1 873 58 0.7 560 15 485 430 460 Comparative Example 6 Inventing Steel 1 866 75 0.7 350 14 493 425 465 Comparative Example 7 Inventing Steel 1 866 115 0.7 435 18 477 435 451 Comparative Example 8 Inventing Steel 1 872 20 0.7 460 18 480 452 462 Comparative Example 9 Inventing Steel 1 880 78 4.2 472 15 491 438 460 Comparative Example 10 Inventing Steel 1 876 86 0.7 471 22 550 530 540 Comparative Example 11 Inventing Steel 1 855 58 0.7 483 21 420 350 380 Comparative Example 12 Inventing Steel 1 884 60 0.7 468 55 485 430 445 Comparative Example 13 Inventing Steel 1 890 62 0.6 475 25 390 465 420 Comparative Example 14 Inventing Steel 1 878 60 0.7 454 22 455 540 485 Invention Example 1 Invention of Steel 2 873 72 0.7 452 16 460 455 460 Invention Example 2 Invention Steel 3 887 75 0.6 482 22 500 460 485 Invention Example 3 Invention of Steel 4 892 65 0.7 469 25 470 425 454 Invention Example 4 Invention Steel 5 874 62 0.7 462 25 500 458 482 Invention Example 5 Invention Steel 6 891 64 0.7 450 18 475 435 456 Invention Example 6 Invention Steel 7 888 60 0.7 463 19 450 432 445 Invention Example 7 Invention Steel 8 885 69 0.6 450 16 466 420 448 Invention Example 8 Inventing Steel 9 892 82 0.7 473 19 495 453 472 Invention Example 9 Invent Steel 10 892 82 0.7 454 20 505 451 477 Invention Example 10 Inventing Steel 11 885 72 0.7 465 22 482 426 440

[0139] The microstructure of each hot-rolled steel plate manufactured as described above was analyzed in detail and is shown in Table 3 below. At this time, the microstructure was analyzed at the 1 / 4 to 1 / 2 position of the thickness of the cross section in the rolling direction of the steel plate. In addition, the distinction and area fraction measurement of the polygonal ferrite phase (PF), bainite phase (B), ferrite low temperature phase transformation phase (BF), martensite phase (M) and MA phase formed in the steel plate were analyzed using electron backscatter diffraction (EBSD, (JEOL JSM-7001F)) and at a magnification of 3000 to 5000. In addition, the average dislocation density (geometric required dislocation (GND)) was analyzed using OIM after EBSD measurement at the 1 / 4 position of the thickness of the hot-rolled steel plate based on the cross section parallel to the rolling direction. TM (EDAX) was calculated. In addition, as shown in Table 3, the microstructure of the hot-rolled steel sheet of this Example was observed, and no retained austenite was observed.

[0140] [Table 3]

[0141] category Steel Type PF BF B P M MA <![CDATA[GND(×10 14 )]]> Comparative Example 1 Comparison Steel 1 39 54 7 0 0 0 0.8 Comparative Example 2 Comparison Steel 2 16 69 8 0 5 2 2.6 Comparative Example 3 Comparison Steel 3 11 69 7 0 7 6 2.6 Comparative Example 4 Comparison Steel 4 33 67 0 0 0 0 0.9 Comparative Example 5 Inventing Steel 1 15 78 3 4 0 0 0.95 Comparative Example 6 Inventing Steel 1 6 71 18 0 3 2 2.8 Comparative Example 7 Inventing Steel 1 5 85 10 0 0 0 2.6 Comparative Example 8 Inventing Steel 1 34 58 8 0 0 0 0.9 Comparative Example 9 Inventing Steel 1 32 60 4 5 0 0 0.9 Comparative Example 10 Inventing Steel 1 29 64 3 4 0 0 0.95 Comparative Example 11 Inventing Steel 1 11 63 15 0 8 3 3.3 Comparative Example 12 Inventing Steel 1 14 68 14 0 0 4 2.8 Comparative Example 13 Inventing Steel 1 12 62 22 0 2 2 3.1 Comparative Example 14 Inventing Steel 1 31 64 0 8 0 0 0.95 Invention Example 1 Invention of Steel 2 25 68 5 0 0 2 1.5 Invention Example 2 Invention Steel 3 21 75 4 0 0 1 1.2 Invention Example 3 Invention of Steel 4 25 72 3 0 0 0 1.3 Invention Example 4 Invention Steel 5 20 67 13 0 0 0 2.3 Invention Example 5 Invention Steel 6 25 72 3 0 0 0 1.6 Invention Example 6 Invention Steel 7 23 65 12 0 0 0 1.9 Invention Example 7 Invent Steel 8 21 68 11 0 0 0 2.2 Invention Example 8 Inventing Steel 9 29 62 9 0 0 0 1.7 Invention Example 9 Invent Steel 10 20 73 7 0 0 0 1.8 Invention Example 10 Inventing Steel 11 17 71 12 0 0 0 2.4

[0142] * In Table 3, PF refers to polygonal ferrite phase, B refers to bainite phase, BF refers to ferrite low temperature transformation phase, M refers to martensite phase, and MA phase refers to martensite and austenite phase. In addition, GND represents the average dislocation density (geometric required dislocation).

[0143] Table 4 below shows the results of evaluating the mechanical properties of each of the hot-rolled steel sheets obtained.

[0144] Mechanical properties are evaluated by tensile test and HER results by changing the blanking gap to 5%, 10%, and 20%, respectively, for each steel plate. YS, TS, and T-El represent 0.2% offset yield strength, tensile strength, and elongation at break, and are the results of testing JIS No. 5 standard test pieces collected in the direction perpendicular to the rolling direction. In addition, HER evaluation is based on JFST 1001-1996 standard. The results of the HER test in Table 4 below are the average values ​​after 3 times.

[0145] [Table 4]

[0146]

[0147] As shown in Tables 1 to 4, it can be confirmed that Invention Examples 1 to 10, in which the steel compositions and the manufacturing process conditions satisfy the scope of the present invention, all have excellent mechanical properties.

[0148] On the other hand, Comparative Examples 1 and 2 do not satisfy Relational Formula 1. The composition of the components of Comparative Example 1 exceeds the range shown in Relational Formula 1, and exhibits high yield strength and high tensile strength due to the precipitation strengthening effect. In addition, the average dislocation density in the microstructure of the steel is 0.9×10 14 m -2 When the thickness is smaller than the range proposed in the present invention, an excessive amount of polygonal ferrite phase is formed, and thus a worse HER value is shown when the blanking gap is large.

[0149] Comparative Example 2 is a case where the steel component composition does not fall within the range shown in Relational Formula 1, and excessive martensite phase and MA phase are formed, and the average dislocation density is as high as 2.6×10 14 m -2 , and the HER value is poor compared to the material strength.

[0150] Comparative Examples 3 and 4 are cases where the composition of the steel components does not satisfy Relational Formula 2, and Comparative Example 3 is a case where the composition of the steel exceeds the range shown in Relational Formula 1. In this case, too much alloy component with high hardenability is contained, thereby ensuring stable strength, but the elongation is insufficient and the HER value is also poor. In addition, Comparative Example 4 is a case where the composition of the steel does not reach the range shown in Relational Formula 1, the alloy component with high hardenability is insufficient, the bainite phase is not formed in the fine structure, and the target strength cannot be ensured.

[0151] In addition, in Comparative Examples 5 to 14, the steel compositions are within the range of the present invention, but the manufacturing process conditions are not within the range of the present invention.

[0152] Specifically, Comparative Examples 5 and 6 are cases where the termination temperature during the first cooling after hot rolling is not within the range proposed by the present invention. In Comparative Example 5, the termination temperature during the first cooling exceeds the range of the present invention, thereby forming an unnecessary pearlite structure, and the HER value is poor. In addition, in Comparative Example 6, the termination temperature during the first cooling does not reach the range of the present invention, the polygonal ferrite phase is insufficient, the ferrite system low temperature phase transformation phase forms an unnecessary MA phase in the grains, and a martensite structure is also formed in a part of the grain boundaries, and the HER value is poor.

[0153] Comparative Examples 7 and 8 are cases where the cooling rate during the primary cooling is not within the range proposed by the present invention. Specifically, in Comparative Example 7, the cooling rate exceeds the range proposed by the present invention, the polygonal ferrite phase is insufficient, and most of the fine structure is composed of a ferrite system low temperature phase transformation phase, so the elongation is very poor. In addition, in Comparative Steel 8, the cooling rate does not reach the proposed range, the polygonal ferrite phase fraction increases significantly, and the elongation is good, but when the blanking gap is large, the HER is poor.

[0154] Comparative Example 9 is a case where the cooling stop time after the primary cooling is unnecessarily long. The polygonal ferrite phase fraction in the fine structure of the steel sheet is greatly increased, and an unnecessary pearlite structure is formed, so the HER value is poor when the blanking gap is small or large.

[0155] Comparative Examples 10 and 11 are cases where the cooling end temperature during secondary cooling is not within the range proposed by the present invention. Comparative Example 10 is a case where the secondary cooling end temperature exceeds the temperature range of the present invention, the polygonal ferrite phase fraction increases significantly, and an unnecessary pearlite structure is formed, so the HER value is poor. In addition, Comparative Example 11 is a case where the temperature of the steel plate after secondary cooling does not reach the range of the present invention, and martensite and MA phase are unnecessarily formed in the fine structure, so the HER value is poor.

[0156] In Comparative Example 12, the cooling rate during the secondary cooling exceeded the range proposed by the present invention, too much MA phase was formed in the fine structure, and the HER value was poor.

[0157] Comparative Examples 13 and 14 are cases where the temperature of the steel plate in the width direction before coiling does not meet the standard of the present invention. Specifically, Comparative Example 13 is a case where the temperature of the edge of the steel plate does not meet the standard of the present invention, and too many low-temperature phase transformation phases including martensite phases are formed at the edge, so the strength is high, but the elongation is insufficient, and the HER is also poor. In addition, Comparative Example 14 is a case where the temperature of the center of the steel plate exceeds the standard of the present invention, and a pearlite structure is formed in the center of the steel plate, and the bainite structure is also degraded, making it difficult to distinguish from the pearlite structure, and the HER characteristics are poor.

[0158] in addition, Figure 1 : is a graph showing the HER relationship according to the punching clearance of Inventive Examples 1 to Inventive Examples 10 and Comparative Examples 5 to Comparative Examples 14 in the embodiment of the present invention. Figure 1 As shown, it can be confirmed that in the case of Inventive Examples 1 to 10, the tensile strength is excellent and is 590 MPa or more compared with Comparative Examples 5 to 14, and the HER 0 The HER value of the actual hole expansion and flanging formability is excellent and is 100% or more, and the product of tensile strength × HER is 45000 MPa% or more in a wide blanking gap range of 5-20%.

[0159] As mentioned above, the preferred embodiments of the present invention are described in the detailed description of the present invention, but various modifications can be made by those skilled in the art without departing from the scope of the present invention. Therefore, the scope of the rights of the present invention should not be limited to the described embodiments, but should be determined by the claims and their equivalents.

Claims

1. A hot-rolled steel sheet, comprising, by weight, C: 0.03-0.08%, Si: 0.01-1.0%, Mn: 1.0-2.0%, acid-soluble aluminum (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%, and the balance of Fe and unavoidable impurities, and the hot-rolled steel sheet satisfies the following equations 1 and 2, The main phase of the microstructure is a ferrite system low temperature phase transformation phase composed of one or more of acicular ferrite, granular bainitic ferrite and bainitic ferrite, and the microstructure contains less than 40% of bainite phase and polygonal ferrite phase and less than 3% of retained pearlite, martensite, retained austenite and MA phase in total, and the average dislocation density of the microstructure, i.e., the geometric required dislocation, satisfies 1.0×10 14 -2.5×10 14 m -2 , [Equation 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 formula 1, Nb, Ti, C, N, and S are the weight % of the corresponding alloying elements, and are entered as 0 when not added. [Equation 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 % of the corresponding alloying elements, and are entered as 0 when not added.

2. The hot rolled steel sheet according to claim 1, wherein: The sum of the area fractions of the bainite phase and the polygonal ferrite phase is 10-40%.

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, wherein: The hot-rolled steel sheet has a tensile strength of 590 MPa or more, and a HER value that satisfies a product of tensile strength×HER of 45000 MPa% or more within a range of a blanking gap of 5-20%.

6. A method for manufacturing a hot-rolled steel plate, comprising the following steps: The steel billet is reheated to 1150-1350° C., and the steel billet comprises, in terms of weight%, C: 0.03-0.08%, Si: 0.01-1.0%, Mn: 1.0-2.0%, acid-soluble aluminum: 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%, and the balance is Fe and unavoidable impurities, and the steel billet satisfies the following Relationships 1 and 2; hot rolling the reheated steel billet at a temperature in the range of 850-1150° C.; Cooling the hot-rolled steel plate at an average cooling rate of 50-100° C. / s to a temperature in the range of 400-500° C.; Reheating the steel plate to a temperature in the range of 450-550° C. by holding the steel plate after the primary cooling for 0.5-3 seconds; The reheated steel plate is subjected to secondary cooling at an average cooling rate of 1-30°C / second to a temperature in the range of 400-500°C, and then coiled; as well as The coiled sheet is cooled to a temperature ranging from room temperature to 200° C. at an average cooling rate of 0.1-25° C. / hour, In one or more of the primary cooling process and the secondary cooling process, the temperature T of the two edge portions (W×60%) relative to the entire width W of the steel plate is set to E The temperature T of the center of the width (W×40%) is 450-550°C. C The cooling is terminated in such a way that the temperature reaches 400-500°C, so that the average temperature T of the coiled sheet after coiling in the coiling step is A Can be maintained in the range of 400-500℃, [Equation 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 formula 1, Nb, Ti, C, N, and S are the weight % of the corresponding alloying elements, and are entered as 0 when not added. [Equation 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 % of the corresponding alloying elements, and are entered as 0 when not added.

7. The method for producing a hot-rolled steel sheet according to claim 6, wherein: The manufacturing method further comprises the steps of pickling and oiling the steel plate that is coiled after secondary cooling.

8. The method for producing a hot-rolled steel sheet according to claim 6, wherein: The manufacturing method further comprises the step of heating the pickled and oiled steel sheet to a temperature range of 450-740° C. and then hot-dip galvanizing.

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