High-strength thick steel plate with excellent formability and method for manufacturing the same

Abstract

The present invention relates to a high-strength thick steel plate having excellent formability and a method for manufacturing the same. More specifically, an object of the present invention is to provide a high-strength thick steel plate and a method for manufacturing the same, wherein the high-strength thick steel plate ensures a uniform microstructure, thereby having excellent yield strength and elongation, and thus does not generate cracks during forming.

Description

Technical Field

[0001] The present invention relates to a thick steel plate and a method for manufacturing the same, and relates to a thick steel plate having high strength characteristics and excellent formability and a method for manufacturing the same. Background Art

[0002] Conventional commercial vehicles and heavy equipment structural components primarily use thick plate materials with a thickness of 12-14 mm and a tensile strength of 440 MPa or higher, manufactured using the thick plate process. However, in recent years, developments are underway to utilize high-strength steel materials with a tensile strength of 550 MPa or higher, in pursuit of both lightweighting and increased strength. In particular, the ultra-thick 15-25 mm thick steel used in large commercial vehicles, special-purpose vehicles, and heavy equipment components is manufactured using the thick plate process, but to ensure price competitiveness, a method utilizing hot rolling is required.

[0003] However, when manufacturing high-strength ultra-thick steel through the hot rolling process, it is difficult to achieve large reduction during rolling and form a uniform fine structure, so it is difficult to ensure a stable yield strength. In addition, cracks are easily generated when manufacturing parts, and local stress concentration occurs during use, resulting in a problem of poor durability.

[0004] In response to this, conventional steel materials, such as those described in Patent Document 1, undergo conventional hot rolling in the austenite region and then are coiled at high temperatures, maintaining strength and ductility by forming fine precipitates with a ferrite phase as the matrix. Alternatively, as described in Patent Document 2, the coiling temperature is cooled to a temperature where a bainite phase forms the matrix, thereby preventing the formation of coarse pearlite. Furthermore, as described in Patent Document 3, a technique is proposed for refining austenite grains by applying a reduction of 20-40% or more in the non-recrystallized region during hot rolling using Ti and Nb, among other factors.

[0005] However, while alloying elements such as Si, Mn, Al, Mo, and Cr, primarily used in the aforementioned technologies for producing thick, high-strength steel, effectively improve strength, excessive addition of these elements can lead to segregation and microstructural inhomogeneity, resulting in poor formability. Microcracks generated on shear surfaces are more likely to propagate in fatigue environments, leading to component failure. In particular, as thickness increases, microstructural inhomogeneity increases between the surface and deeper layers, increasing local stress concentration and the crack propagation rate in fatigue environments, leading to poor durability.

[0006] Furthermore, in order to refine the grains of thick materials and obtain precipitation strengthening effects, it is effective to use precipitate-forming elements such as Ti, Nb, and V. However, when coiling is performed at high temperatures of 500-700°C, where precipitates are easily formed, or when the cooling rate is not controlled during cooling after hot rolling, coarse carbides are formed in the center of the thickness of the thick material, thereby deteriorating the quality of the shear surface.

[0007] Furthermore, during hot rolling, a 20-40% reduction can be easily applied to thin products more than twice in the non-recrystallized zone, but this is difficult to apply when manufacturing thick products with a smaller total reduction than thin products.

[0008] [Prior art literature]

[0009] (Patent Document 1) Japanese Patent Publication No. 2002-322541

[0010] (Patent Document 2) Korean Patent Gazette No. 10-1528084

[0011] (Patent Document 3) Japanese Patent Publication No. 1997-143570 Summary of the Invention

[0012] Technical problems to be solved

[0013] According to one aspect of the present invention, an object is to provide a high-strength thick steel plate and a method for manufacturing the same, wherein the high-strength thick steel plate has excellent yield strength and elongation by ensuring a uniform fine structure during a hot rolling process of a thick material, thereby preventing cracks from being generated during forming.

[0014] The technical problems of the present invention are not limited to the above contents. Anyone skilled in the art can understand the additional technical problems of the present invention from the entire contents of this specification.

[0015] Technical Solution

[0016] One aspect of the present invention may provide a thick steel plate comprising, in weight percent, C: 0.05-0.15%, Si: 0.01-1.0%, Mn: 1.0-2.0%, Cr: 0.005-1.0%, Al: 0.01-0.1%, P: 0.001-0.02%, S: 0.001-0.01%, N: 0.001-0.01%, Ti: 0.005-0.11%, Nb: 0.005-0.07%, and the balance being Fe and unavoidable impurities, wherein the R value of the thick steel plate, as defined in the following equation 1, satisfies 0.3 to 1. 0, based on the cross section, in terms of area %, a surface portion in a range of 0 to t / 4 (wherein t represents the thickness of the steel plate) and a deep portion in a range of t / 4 to t / 2 (excluding t / 4) each contain, as a microstructure, a total of 90% or more of ferrite and bainite, less than 5% of pearlite and carbides having a diameter of 0.5 μm or more, and less than 5% of MA phase (martensite and austenite), the product of the yield strength and elongation (YS×T-El) of the thick steel plate is 16000 MPa·% or more, and the thickness of the thick steel plate is 10 mm or more.

[0017] [Equation 1]

[0018] R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]-0.9×[Si]-1.5×[Nb]

[0019] [C]*=[C]-[C]×Q

[0020] Q=([Nb] / 93+[Ti] / 48) / ([C] / 12)

[0021] ([C], [Mn], [P], [S], [Si], [Nb], and [Ti] in the above-mentioned relational formula 1 are the weight % of the corresponding alloying elements.)

[0022] The thickness of the steel plate may be greater than 15 mm.

[0023] In the deep portion of the steel plate, the pearlite and carbides having a diameter of 0.5 μm or more may account for 3% or less, and the MA phase may account for 3% or less, in terms of area%.

[0024] In the surface layer portion of the steel plate, the bainite may be 20% or less, the pearlite and carbides having a diameter of 0.5 μm or more may be less than 2%, and the MA phase may be 3% or less, in terms of area%.

[0025] Taking any line perpendicular to the thickness section of the steel plate as a reference, the difference between the average hardness value and the maximum hardness value of the hardness values ​​measured at intervals of 0.5 mm from a position 0.5 mm directly below the surface of the test piece to a position 0.5 mm directly below the surface of the other side can be less than 20 Hv.

[0026] Another aspect of the present invention may provide a method for manufacturing a thick steel plate, comprising the steps of: reheating a steel slab, wherein the steel slab comprises, in terms of weight %, C: 0.05-0.15%, Si: 0.01-1.0%, Mn: 1.0-2.0%, Cr: 0.005-1.0%, Al: 0.01-0.1%, P: 0.001-0.02%, S: 0.001-0.01%, N: 0.001-0.01%, and Ti: 0.005-0. 11%, Nb: 0.005-0.07%, and the balance Fe and inevitable impurities, and the steel slab satisfies the R value defined in the following relational expression 1 of 0.3 to 1.0; a hot rolling step, hot rolling the reheated steel slab at a temperature range of 800-1150° C. at a reduction rate of 20-50% to a thickness of 10 mm or more, and terminating the rolling at a temperature range of Tn-50 to Tn defined in the following relational expression 2; and subjecting the hot-rolled steel plate to a CR defined in the following relational expression 3. 最小 The steel plate is cooled at the above cooling rate to a temperature range of 450-550° C., and then coiled; and the coiled steel plate is cooled for a second time.

[0027] [Equation 1]

[0028] R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]-0.9×[Si]-1.5×[Nb]

[0029] [C]*=[C]-[C]×Q

[0030] Q=([Nb] / 93+[Ti] / 48) / ([C] / 12)

[0031] ([C], [Mn], [P], [S], [Si], [Nb], and [Ti] in the above-mentioned relational formula 1 are the weight % of the corresponding alloying elements.)

[0032] [Equation 2]

[0033] Tn=730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]-80×[Si]-1.4×(t-8)

[0034] (The unit of Tn in the above relational expression 2 is °C, and [C], [Mn], [Cr], [Nb], [Ti], and [Si] are the weight % of the corresponding alloying elements.)

[0035] (In the above equation 2, t is the thickness (mm) of the final rolled plate.)

[0036] [Equation 3]

[0037] CR 最小 =76.6-157×[C]-25.2×[Si]-14.1×[Mn]-27.3×[Cr]+61×[Ti]+448×[Nb]

[0038] (CR of the relational expression 3 最小 The unit is °C / s, and [C], [Si], [Mn], [Cr], [Ti], and [Nb] are the weight % of the corresponding alloying elements.)

[0039] The reheating may be performed at a temperature in the range of 1200-1350°C.

[0040] During the primary cooling, the cooling rate may be 80° C. / second or less.

[0041] During the secondary cooling, air cooling or water cooling may be performed to a temperature range from room temperature to 200°C.

[0042] Beneficial effects

[0043] According to one aspect of the present invention, a thick steel plate having excellent tensile strength, yield strength, and elongation, and thus high strength characteristics and excellent formability, and a method for manufacturing the same can be provided.

[0044] According to another aspect of the present invention, a high-strength thick steel plate for structural components such as rims, wheel discs, components, and frames of large commercial vehicles and a method for manufacturing the same are provided. BRIEF DESCRIPTION OF THE DRAWINGS

[0045] Figure 1 It is a graph showing the product of yield strength and elongation (YS×T-E1) and the difference (ΔH) between the average hardness value and the maximum hardness value of the thickness cross section of the invention steel and the comparative steel.

[0046] Figure 2 and Figure 3 Graphs showing the distribution of hardness values ​​at thickness cross sections of the inventive steel and the comparative steel, respectively.

[0047] Best Practice

[0048] Below, preferred specific embodiments of the present invention are described. The specific embodiments of the present invention can be modified into various forms, and the scope of the present invention should not be construed as being limited to the specific embodiments described below. This specific embodiment is provided to explain the present invention in more detail to those skilled in the art.

[0049] To solve the above-mentioned problems, the inventors of the present invention studied the distribution of microstructures in each thickness direction and detailed changes in material properties depending on the composition, hot rolling, and cooling conditions of ultra-thick rolled steel materials having various compositions.

[0050] As a result, a method for imparting excellent yield strength and ductility to thick hot-rolled steel plates was identified. In particular, uniformity was ensured in the microstructure of thick steel plates having a thickness greater than a certain value, and it was confirmed that the hardness distribution in the thickness direction could be constant, thereby completing the present invention.

[0051] Hereinafter, the present invention will be described in detail.

[0052] Hereinafter, the steel composition of the present invention will be described in detail.

[0053] Unless otherwise specified in the present invention, % indicating the content of each element is based on weight.

[0054] The steel plate according to one aspect of the present invention may contain, in terms of weight %, C: 0.05-0.15%, Si: 0.01-1.0%, Mn: 1.0-2.0%, Cr: 0.005-1.0%, Al: 0.01-0.1%, P: 0.001-0.02%, S: 0.001-0.01%, N: 0.001-0.01%, Ti: 0.005-0.11%, Nb: 0.005-0.07%, and the balance Fe and inevitable impurities.

[0055] Carbon (C): 0.05-0.15%

[0056] Carbon (C) is the most economical and effective element for strengthening steel. As the amount added increases, the precipitation strengthening effect or the fraction of the bainite phase increases, thereby increasing tensile strength. As the thickness of the hot-rolled steel sheet increases, the cooling rate in the center of the thickness slows down during cooling after hot rolling. High carbon (C) content increases the likelihood of coarse carbides or pearlite forming. A sufficient strengthening effect is difficult to achieve with a carbon (C) content of less than 0.05%. However, a carbon (C) content exceeding 0.15% leads to poor formability and reduced durability due to the formation of coarse carbides or pearlite phases and banded structures in the center of the thickness. Furthermore, weldability also deteriorates.

[0057] Therefore, the carbon (C) content may be 0.05-0.15%, more preferably 0.06% or more, and 0.12% or less.

[0058] Silicon (Si): 0.01-1.0%

[0059] Silicon (Si) deoxidizes molten steel, has a solid solution strengthening effect, and delays the formation of coarse carbides, making it an element that helps improve formability. When the silicon (Si) content is less than 0.01%, the solid solution strengthening effect is insufficient and the effect of delaying carbide formation is not significant, making it difficult to improve formability. When the silicon (Si) content exceeds 1.0%, due to the increase in phase transformation temperature, when hot rolling extremely thick materials in the low temperature zone, it is easy to roll through the local ferrite zone to form coarse grains in the surface part. Due to silicon (Si), red oxide scale forms on the surface of the steel plate, not only the surface quality of the steel plate becomes very poor, but also there is a problem of reduced ductility and weldability.

[0060] Therefore, the content of silicon (Si) may be 0.01-1.0%, more preferably 0.1% or more and 0.9% or less.

[0061] Manganese (Mn): 1.0-2.0%

[0062] Manganese (Mn), like Si, is an element that is effective in solid solution strengthening of steel and increases the hardenability of steel, making it easier to form a bainite phase during cooling after hot rolling. When the manganese (Mn) content is less than 1.0%, the above-mentioned effect of addition cannot be achieved. When the manganese (Mn) content exceeds 2.0%, due to the significant increase in hardenability, martensitic transformation is likely to occur, and pearlite formation is promoted during high-temperature coiling. In addition, when casting slabs in the continuous casting process, a segregation area in the center of the thickness is very developed, and a fine structure is unevenly formed in the thickness direction during cooling after hot rolling, resulting in poor formability and durability.

[0063] Therefore, the content of manganese (Mn) may be 1.0-2.0%, and more preferably may be 1.1% or more.

[0064] Chromium (Cr): 0.005-1.0%

[0065] Chromium (Cr) is a solid solution strengthening element for steel. It delays the ferrite transformation during cooling, thus contributing to the formation of bainite. When the chromium (Cr) content is less than 0.005%, the aforementioned effects of addition are not achieved. When the chromium (Cr) content exceeds 1.0%, the ferrite transformation is excessively delayed, resulting in the formation of a martensite phase, which deteriorates elongation. Furthermore, similar to Mn, segregation develops significantly in the center of the thickness, making the microstructure uneven across the thickness, thereby deteriorating formability and durability.

[0066] Therefore, the content of chromium (Cr) may be 0.005-1.0%, more preferably 0.1% or more and 0.9% or less.

[0067] Aluminum (Al): 0.01-0.1%

[0068] Aluminum (Al) is an element added primarily for deoxidation. When the Al content is less than 0.01%, the added effect is insufficient. When the Al content exceeds 0.1%, it combines with nitrogen to form AlN, which can easily cause corner cracks in the slab during continuous casting and defects caused by the formation of inclusions.

[0069] Therefore, the content of aluminum (Al) may be 0.01-0.1%.

[0070] Phosphorus (P): 0.001-0.02%

[0071] Phosphorus (P), like Si, is an element that simultaneously exhibits solid solution strengthening effects and promotes ferrite transformation. When the phosphorus (P) content exceeds 0.02%, brittleness results from grain boundary segregation, microcracks are easily generated during molding, and formability and durability are significantly degraded. Furthermore, reducing the phosphorus (P) content to less than 0.001% requires significant manufacturing costs, making it uneconomical and insufficient to achieve sufficient strength.

[0072] Therefore, the content of phosphorus (P) may be 0.001-0.02%.

[0073] Sulfur (S): 0.001-0.01%

[0074] Sulfur (S) is an impurity present in steel. When the sulfur (S) content exceeds 0.01%, it combines with Mn and other materials to form non-metallic inclusions. This can easily cause fine cracks during steel cutting, leading to reduced formability and durability. Furthermore, reducing the sulfur (S) content to less than 0.001% requires significant time during steelmaking operations, reducing productivity.

[0075] Therefore, the content of sulfur (S) may be 0.001-0.01%.

[0076] Nitrogen (N): 0.001-0.01%

[0077] Nitrogen (N), along with carbon, is a representative solid solution strengthening element, forming coarse precipitates along with titanium, aluminum, and other elements. While nitrogen (N) generally exhibits a superior solid solution strengthening effect than carbon, increasing its content significantly reduces toughness, leading to a limit of 0.01% for nitrogen (N). Furthermore, reducing the nitrogen (N) content to less than 0.001% requires significant steelmaking time, reducing productivity.

[0078] Therefore, the content of nitrogen (N) may be 0.001-0.01%.

[0079] Titanium (Ti): 0.005-0.11%

[0080] Titanium (Ti) is a representative precipitation strengthening element. Due to its strong affinity with N, coarse TiN is formed in steel. TiN has the effect of inhibiting grain growth during the heating process for hot rolling. In addition, since the remaining titanium (Ti) after reacting with N is dissolved in the steel and combines with C to form TiC precipitates, it is useful for improving the strength of the steel. When the titanium (Ti) content is less than 0.005%, the above effect cannot be achieved. When the titanium (Ti) content exceeds 0.11%, coarse TiN is generated. Due to the coarsening of the precipitates, local stress concentration occurs during forming, which causes the problem of easy cracking.

[0081] Therefore, the content of titanium (Ti) may be 0.005-0.11%, more preferably 0.01% or more and 0.1% or less.

[0082] Niobium (Nb): 0.005-0.07%

[0083] Niobium (Nb), along with Ti, is a representative precipitation strengthening element. It precipitates during hot rolling and effectively improves the strength and impact toughness of steel by retarding recrystallization and resulting in grain refinement. However, Nb contents below 0.005% do not achieve these benefits. However, Nb contents exceeding 0.07% lead to excessive recrystallization retardation during hot rolling, resulting in the formation of elongated grains and coarse composite precipitates, which can impair formability and durability.

[0084] Therefore, the niobium (Nb) content may be 0.005-0.07%, and more preferably may be 0.01% or more.

[0085] In addition to the above composition, the steel material of the present invention may contain a balance of iron (Fe) and unavoidable impurities. Unavoidable impurities may be undesirably incorporated during conventional manufacturing processes and cannot be eliminated. Such impurities are well known to those skilled in the art of conventional iron and steel manufacturing, and therefore, not all of them are specifically described in this specification.

[0086] The steel of the present invention may have an R value defined in the following Relationship 1 of 0.3 to 1.0.

[0087] By controlling R in relational equation 1, the segregation of C, Mn, P, S, etc. and the formation of MnS that occur during the solidification of steel and the cooling of the slab in the continuous casting process can be minimized, thereby improving the uniformity of the microstructure. It is well known that segregation of alloying elements such as C and Mn occurs in the cast structure formed during solidification, and P has the characteristic of mainly segregating at the grain boundaries when the steel plate is kept at a high temperature after hot rolling and cooling, and thus becomes the cause of grain boundary embrittlement. The segregation as described above is highly dependent on the content of the alloying elements. In particular, C and Mn form coarse carbides and pearlite structures during the cooling process after hot rolling, and this becomes the cause of the deterioration of the shear surface quality. In addition, Mn forms MnS as a non-metallic inclusion together with S, and the MnS elongates during the rolling process, so there is a problem of significantly deteriorating the formability of the final product. Furthermore, Si suppresses the formation of coarse carbides and has a high solid solution strengthening effect even with a small alloying amount, while Nb and Ti form fine precipitates and effectively reduce the grain size, thereby effectively improving the aforementioned segregation and grain boundary embrittlement problems.

[0088] In the present invention, steels having various alloy compositions were hot rolled to produce thick materials, and the hardness of the cross section was measured. It was confirmed that there was a correlation between the uniformity of the microstructure, the alloy composition, and its content, thereby obtaining Relationship 1.

[0089] When the R value defined in the following equation 1 is less than 0.3, it is difficult to ensure the desired physical properties of the present invention. However, when the R value defined in the following equation 1 exceeds 1.0, the heterogeneity of the microstructure increases, resulting in a significant change in the hardness value at the cross section. More preferably, the lower limit of the R value may be 0.5, and the upper limit may be 0.8.

[0090] [Equation 1]

[0091] R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]-0.9×[Si]-1.5×[Nb]

[0092] [C]*=[C]-[C]×Q

[0093] Q=([Nb] / 93+[Ti] / 48) / ([C] / 12)

[0094] ([C], [Mn], [P], [S], [Si], [Nb], and [Ti] in the above-mentioned relational formula 1 are the weight % of the corresponding alloying elements.)

[0095] Hereinafter, the microstructure of the steel of the present invention will be described in detail.

[0096] Unless otherwise specified in the present invention, % indicating the fraction of fine structures is based on area.

[0097] Based on the cross section, in terms of area %, the surface layer portion in the range of 0 to t / 4 (where t represents the thickness of the steel plate) and the deep layer portion in the range of t / 4 to t / 2 (excluding t / 4) of the steel meeting the alloy composition according to one aspect of the present invention can each contain, as a fine structure, 90% or more of ferrite and bainite in total, less than 5% of pearlite and carbides with a diameter of 0.5 μm or more, and less than 5% of MA phase (martensite and austenite).

[0098] The fine structure of thick high-strength steel crystallizes during cooling, and coarse carbides and pearlite are more likely to form in the deep layer with a slow cooling rate, compared to the surface layer with a fast cooling rate, where bainite and MA phase (martensite and austenite) are more likely to form.

[0099] Typically, the MA phase formed in the surface layer is a hard phase that exhibits higher hardness than the surrounding microstructure, resulting in an uneven hardness distribution. This difference in hardness between the MA phase and the matrix can also cause microcracks during molding. Furthermore, the coarse carbides and pearlite formed in the deeper layers exhibit higher hardness than the surrounding microstructure but are also brittle, causing microcracks during shear molding.

[0100] Therefore, in the present invention, in order to simultaneously solve the problems in the surface and deep layers, the content of pearlite and carbides with a diameter of 0.5 μm or greater is limited to less than 5%, and the content of MA phase is limited to less than 5%. In this case, the content of pearlite, carbides with a diameter of 0.5 μm or greater, and MA phase can be applied equally to both the surface and deep layers.

[0101] The purpose of including 90% or more of ferrite and bainite in the present invention is to suppress the formation of unnecessary coarse carbides and pearlite, achieve a uniform hardness distribution across the thickness, and ensure excellent yield strength and elongation. If the content of ferrite and bainite is less than 90%, it is difficult to achieve the desired yield strength and elongation product (YS × T - E1) value in the present invention. Therefore, the total content of ferrite and bainite in the present invention can be 90% or more.

[0102] In terms of ensuring the desired physical properties of the present invention, it is more preferred that in the deep portion, the pearlite and carbides with a diameter of 0.5 μm or more may be 3% or less, and the MA phase may be 3% or less, and in the surface portion, the bainite may be 20% or less, the pearlite and carbides with a diameter of 0.5 μm or more may be less than 2%, and the MA phase may be 3% or less.

[0103] In the present invention, the microstructure has the same characteristics in both the surface and deep layers of the steel, and the microstructure proposed in the present invention applies uniformly to the entire steel. Furthermore, in the present invention, the surface portion refers to the region from 0 to t / 4 (t is the thickness of the steel plate), and the deep portion refers to the region from t / 4 to t / 2 (excluding t / 4), based on the cross section.

[0104] Hereinafter, the method for producing steel of the present invention will be described in detail.

[0105] The steel according to one aspect of the present invention can be produced by reheating, hot rolling, primary cooling, coiling, and secondary cooling a steel billet satisfying the above-mentioned alloy composition.

[0106] Reheating of slabs

[0107] The steel billet meeting the above alloy composition can be reheated in the temperature range of 1200-1350°C.

[0108] When the reheating temperature is lower than 1200°C, the precipitates cannot be fully redissolved, the formation of precipitates in the process after hot rolling is reduced, and coarse TiN remains. On the other hand, when the reheating temperature exceeds 1350°C, the strength decreases due to abnormal grain growth of austenite grains.

[0109] Hot Rolling

[0110] The reheated slab may be hot rolled at a reduction ratio of 20-50% within a temperature range of 800-1150° C., and the rolling may be terminated at a temperature range of Tn-50 to Tn defined in the following relation 2.

[0111] When the hot rolling temperature exceeds 1150°C, the steel sheet temperature becomes too high, the grain size becomes coarse, and the surface quality of the hot-rolled steel sheet deteriorates. On the other hand, when the hot rolling temperature is below 800°C, due to excessive recrystallization delay, elongated grains develop, anisotropy becomes severe, and formability deteriorates. When rolling is performed below the austenitic temperature range, non-uniform microstructures develop more severely. As a result, fine cracks are easily generated in non-uniform areas during forming, and ductility is also reduced.

[0112] When the rolling end temperature exceeds Tn, the fine structure of the steel becomes coarse and uneven. When the rolling end temperature is lower than Tn-50, in ultra-thick high-strength steel with a steel plate thickness corresponding to 15-25 mm, ferrite phase transformation is promoted in the surface layer where the temperature is relatively low, increasing the fraction of fine ferrite phase, but with an elongated grain shape, which becomes the cause of rapid crack propagation, and uneven fine structure may remain in the center, which is not conducive to durability.

[0113] The rolling end temperature determined by Relational Expression 2 of the present invention refers to the temperature of the hot-rolled steel sheet at the end of hot rolling.

[0114] [Equation 2]

[0115] Tn=730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]-80×[Si]-1.4×(t-8)

[0116] (The unit of Tn in the above relational expression 2 is °C, and [C], [Mn], [Cr], [Nb], [Ti], and [Si] are the weight % of the corresponding alloying elements.)

[0117] (In the above equation 2, t is the thickness (mm) of the final rolled plate.)

[0118] The reduction in the hot rolling temperature range can be 20-50%.

[0119] When the reduction is less than 20%, it is difficult to achieve the recrystallization delay effect, and uneven coarse grains are likely to form. When the reduction exceeds 50%, due to the formation of excessively elongated microstructures, carbides are formed along the grain boundaries, and cracks are likely to occur along the grain boundaries during molding. In addition, due to the reduction of fine precipitates, the precipitation strengthening effect is also reduced.

[0120] One-step cooling and winding

[0121] The hot rolled steel sheet can be subjected to the CR defined in the following equation 3: 最小 The above cooling speed is used for primary cooling, cooling to a temperature range of 450-550°C, and then winding.

[0122] The temperature range from immediately after hot rolling to the cooling end temperature corresponds to the temperature range where ferrite transformation occurs during the cooling process. The cooling rate of the thickness center portion is slower than that of the thickness surface portion of the rolled plate. Therefore, coarse ferrite phase and coarse carbide are formed in the thickness center portion, resulting in an uneven microstructure. Therefore, in order to suppress the above-mentioned problem, in the present invention, it is necessary to cool the steel sheet at a temperature higher than the specific cooling rate (CR 最小 However, when the average cooling rate in the temperature range exceeds 80°C / second (sec), the difference in cooling rate between the surface layer and the deep layer becomes too large, and the difference in hardness between the surface layer and the deep layer increases significantly, thereby deteriorating the formability and durability.

[0123] The cooling rate determined by Relational Expression 3 in the present invention represents the cooling rate of the hot-rolled steel sheet after hot rolling.

[0124] [Equation 3]

[0125] CR 最小=76.6-157×[C]-25.2×[Si]-14.1×[Mn]-27.3×[Cr]+61×[Ti]+448×[Nb]

[0126] (CR of the relational expression 3 最小 The unit is °C / s, and [C], [Si], [Mn], [Cr], [Ti], and [Nb] are the weight % of the corresponding alloying elements.)

[0127] When the cooling end temperature and the coiling temperature exceed 550°C, the pearlite phase forms a banded structure or a large amount of coarse carbides are formed, resulting in poor formability and durability of the steel. When the cooling end temperature and the coiling temperature are lower than 450°C, excessive martensite phase and MA phase are formed, thereby poor formability and durability.

[0128] Secondary cooling

[0129] The coiled steel plate may be subjected to secondary cooling to a temperature range from room temperature to 200° C., and the secondary cooling may be air cooling or water cooling.

[0130] In the present invention, air cooling refers to cooling in the atmosphere at room temperature at a cooling rate of 0.001-10°C / hour. Even if the cooling rate exceeds 10°C / hour, if the above-mentioned coiling temperature and primary cooling conditions are adhered to, a portion of the untransformed phase in the steel can be suppressed from transforming into the MA phase, so water cooling is not a problem. In the present invention, water cooling refers to loading the coil into a water tank at room temperature and cooling it. However, in order to control the cooling rate to less than 0.001°C / hour, separate heating and heat preservation devices are required, which is not conducive to economic efficiency. Therefore, the lower limit of the cooling rate can be 0.001°C / hour.

[0131] The steel plate of the present invention manufactured as described above is a thick steel plate having a thickness of 10 mm or more, more preferably 15 mm or more, and the upper limit of the thickness is 25 mm. The steel plate of the present invention can have a difference between the average hardness value and the maximum hardness value measured at intervals of 0.5 mm from a position 0.5 mm directly below the surface of the test piece to a position 0.5 mm directly below the surface of the other side, based on an arbitrary line perpendicular to the thickness section, and can be 20 Hv or less. More specifically, the average hardness value can be 160-300 Hv. In addition, the product of the yield strength and the elongation (YS×T-El) is 16000 MPa·% or more, thereby achieving high strength and excellent formability.

[0132] The present invention will be described in more detail below through examples. However, it should be noted that the following examples are only used to illustrate the present invention in more detail and are not intended to limit the scope of the present invention. DETAILED DESCRIPTION

[0133] Table 1 below shows the steel composition and final hot-rolled sheet thickness of each steel type. Table 2 below shows the rolling end temperature (FDT), the total of the reduction (%), the coiling temperature (CT), the cooling rate (CR*) from hot rolling to the coiling temperature as the cooling end temperature, Tn and Tn-50 defined in equation 2, and the minimum cooling rate (CR*) defined in equation 3 for the steel types shown in Table 1. 最小 The reheating temperature not listed in Table 2 was the same as 1250°C, the hot rolling temperature was the same as 800-1150°C, and the cooling rate of the steel plate after coiling was the same as 1°C / hour.

[0134] [Table 1]

[0135]

[0136] [Equation 1]

[0137] R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]-0.9×[Si]-1.5×[Nb]

[0138] [C]*=[C]-[C]×Q

[0139] Q=([Nb] / 93+[Ti] / 48) / ([C] / 12)

[0140] ([C], [Mn], [P], [S], [Si], [Nb], and [Ti] in the above-mentioned relational formula 1 are the weight % of the corresponding alloying elements.)

[0141] [Table 2]

[0142]

[0143] [Equation 2]

[0144] Tn=730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]-80×[Si]-1.4×(t-8)

[0145] (The unit of Tn in the above relational expression 2 is °C, and [C], [Mn], [Cr], [Nb], [Ti], and [Si] are the weight % of the corresponding alloying elements.)

[0146] (In the above equation 2, t is the thickness (mm) of the final rolled plate.)

[0147] [Equation 3]

[0148] CR 最小=76.6-157×[C]-25.2×[Si]-14.1×[Mn]-27.3×[Cr]+61×[Ti]+448×[Nb]

[0149] (CR of the relational expression 3 最小 The unit is °C / s, and [C], [Si], [Mn], [Cr], [Ti], and [Nb] are the weight % of the corresponding alloying elements.)

[0150] Tables 3 and 4 below show the microstructural characteristics and mechanical properties of the steel grades.

[0151] The microstructure of Table 3 is the result of analysis at a position of 0.5 mm just below the surface of the hot-rolled plate and in the deep part. In the present invention, the surface part refers to the range of 0 to t / 4 based on the thickness (t), and the deep part refers to the range of t / 4 to t / 2 (excluding t / 4). The microstructure of the surface part of Table 3 is the result of analysis at 0.5 mm just below the surface, and the microstructure of the deep part is the result of analysis at t / 2, which is the center of the thickness. The area fraction of the MA phase is measured by etching with the Lepera etching method, and then analyzed using an optical microscope and an image analyzer at 1000 magnifications. The area fractions of martensite and austenite phases (MA), ferrite phase (F), bainite phase (B) and pearlite phase (P) are the results of analysis using a scanning electron microscope (SEM) at 3000 magnifications and 5000 magnifications. Ferrite (F) refers to polygonal ferrite with an equiaxed crystal structure, while bainite (B) refers to ferrite phases observed in low-temperature regions, including bainite, acicular ferrite, and bainitic ferrite. Furthermore, the area fraction of pearlite (P) refers to the sum of the area fractions of pearlite and carbides larger than 0.5 μm.

[0152] YS, TS and T-El in Table 4 refer to 0.2% offset yield strength, tensile strength and elongation at break, and are the results of taking JIS No. 5 standard test pieces in a direction parallel to the rolling direction and conducting tests. In addition, the hardness at the cross section of the test piece is measured and shown together. For the measurement of hardness, an arbitrary line perpendicular to the thickness section of the test piece is used as a reference, and measurements are performed with a Micro-Vickers testing machine at intervals of 0.5 mm from a position 0.5 mm directly below the surface of the test piece to a position 0.5 mm directly below the surface of the other side, and a load of 500 g is applied. Table 4 shows the maximum hardness value and the average hardness value at the thickness section of the measured hardness values, and shows the difference between the two hardness values. Peak (Pieces) refers to the number of parts where the difference between the hardness value at the thickness position and the average hardness value exceeds 20 Hv.

[0153] [Table 3]

[0154]

[0155] [Table 4]

[0156]

[0157] As shown in Table 4, Invention Steels 1 to 7 satisfying the alloy composition, manufacturing method, and Relational Formulas 1 to 3 proposed in the present invention ensure all the mechanical properties expected by the present invention.

[0158] Figure 1 This is a graph showing the product of the yield strength and elongation of the invention steel and the comparative steel, as well as the difference between the average hardness value and the maximum hardness value of the thickness section. It can be confirmed that the difference in hardness value of the invention steel is less than 20 Hv, and the value of YS×T-El is greater than 16000 MPa·%.

[0159] Figure 2 and Figure 3 The graphs show the hardness distribution of the inventive steel and the comparative steel at the thickness cross section. In the case of the comparative steel, it can be confirmed that the hardness value is relatively low in the center of the thickness compared to the surface layer, and the hardness value varies greatly depending on the thickness position.

[0160] Comparative Steels 1 to 4 do not satisfy Relationship 1 proposed in the present invention. Comparative Steel 1 is a case where the C content satisfies the range of the present invention but exceeds the range of Relationship 1 proposed in the present invention in consideration of segregation. Therefore, in the microstructure, excessive pearlite is formed in the deep part and the surface part, and when the hardness in the thickness direction is measured, a high hardness difference is locally shown. The ductility is also insufficient, and thus results outside the range proposed in the present invention are shown. Comparative Steels 2 and 3 are cases where the Mn composition range proposed in the present invention is not exceeded and Relationship 1 is not satisfied. In Comparative Steel 2, due to the low Mn content, segregation in the thickness direction of the rolled plate does not occur, and coarse carbides and uneven pearlite are not formed, but the yield strength and tensile strength are insufficient, so the desired properties of the present invention cannot be achieved. In Comparative Steel 3, due to the excessive Mn content, bainite is formed in the surface part due to high hardenability, but excessive pearlite is formed in the deep part, and elongated MnS inclusions are also observed. In particular, when measuring hardness across the thickness, localized high hardness differences were observed, and ductility was also insufficient. Comparative Steel 4 exhibits a phosphorus content outside the range specified in the present invention and does not satisfy Relational Equation 1. While the microstructure of Comparative Steel 4 satisfies the range specified in the present invention and exhibits good strength and elongation, hardness measurements reveal localized high hardness differences, posing a high risk of brittleness during use after component manufacture.

[0161] Comparative Steel 5 satisfies Relationship 1 but does not meet the Si content range of the present invention. In the microstructure, coarse ferrite is confirmed to form in the surface portion, and MA phase is also formed in the surface and deep portions. Furthermore, the surface portion exhibits slightly low hardness, and the product of yield strength and elongation is outside the range expected by the present invention. This is because excessive Si addition increases the transformation temperature, causing ferrite to form in the surface portion during hot rolling. Consequently, two-phase rolling occurs, and some of the untransformed ferrite forms the MA phase.

[0162] Comparative Steel 6 meets the alloy composition range of the present invention but does not satisfy Relational Equation 1. In this case, no component segregation is observed, and the microstructure contains almost no MA phase or coarse carbides, with only fine pearlite observed around grain boundaries. Consequently, the hardness distribution through the thickness is relatively uniform. However, the strength values ​​expected by the present invention cannot be achieved.

[0163] Comparative Steel 7 and Comparative Steel 8 are cases where the relationship 2 and the reduction ratio are not satisfied. Comparative Steel 7 is a case where rolling is terminated within the temperature range that satisfies the relationship 2, but due to insufficient reduction ratio, an uneven microstructure is formed during cooling. Therefore, the fraction of the composition of the microstructure satisfies the present invention, but coarse grains are mixed in the ferrite matrix structure, thereby showing a low yield strength. Steel having a microstructure as described above may deteriorate in durability during the use of the component. Comparative Steel 8 is a case where the relationship 2 and the reduction ratio conditions are not satisfied. It can be confirmed that due to the large reduction amount, recrystallization is delayed during the rolling process, so an over-elongated microstructure is formed in the surface part, but equiaxed ferrite and pearlite are mainly formed in the deep part. Therefore, it is composed of an uneven microstructure according to the thickness position, which becomes a cause of the deterioration of the durability of the component, and the elongation is also poor.

[0164] Comparative Steels 9 and 10 do not meet the coiling temperature requirements. In Comparative Steel 9, the cooling end temperature and coiling temperature are above the temperature ranges specified in the present invention, resulting in the local formation of pearlite, with pearlite bands being observed particularly in the deep layers. Consequently, through-thickness hardness measurements show locally high hardness differences. Comparative Steel 10, the cooling end temperature and coiling temperature are below the ranges specified in the present invention. In Comparative Steel 10, excessive bainite forms in the surface microstructure, resulting in insufficient elongation.

[0165] Comparative Steel 11 is a case where the cooling rate does not satisfy the cooling rate condition of Relationship 3, and the cooling rate during cooling after hot rolling is lower than the range of the present invention. Pearlite and coarse carbides are formed in the deep thickness part, so a high hardness difference is locally shown.

[0166] Comparative Steel 12 and Comparative Steel 13 are cases where the conditions for reduction and cooling end temperature are not met. In Comparative Steel 12, the reduction in the temperature region where recrystallization is delayed during hot rolling is insufficient, and the coiling temperature is also low, so the grain size of the ferrite is uneven, and excessive bainite is formed in the microstructure of the surface layer. In addition, local pearlite is observed in the deep thickness portion, resulting in low elongation. Comparative Steel 13 is a case where the reduction in the temperature region where recrystallization is delayed is insufficient, the coiling temperature is high, and the cooling rate does not meet Relationship 3. Therefore, it was confirmed that the microstructure is uneven, the pearlite is formed into a banded structure, and the yield strength is low.

[0167] The present invention has been described in detail above through the embodiments, but other different embodiments are also feasible. Therefore, the technical concept and scope of the claims are not limited to the embodiments.

Claims

1. A thick steel plate comprising, in terms of weight %, C: 0.05-0.15%, Si: 0.01-1.0%, Mn: 1.0-2.0%, Cr: 0.005-1.0%, Al: 0.01-0.1%, P: 0.001-0.02%, S: 0.001-0.01%, N: 0.001-0.01%, Ti: 0.005-0.11%, Nb: 0.005-0.07%, and the balance being Fe and unavoidable impurities. The thick steel plate has an R value defined in the following relational formula 1 that satisfies 0.3 to 1.0, Based on the cross section, in terms of area %, the surface layer portion in the range of 0 to t / 4 and the deep layer portion in the range of t / 4 to t / 2 respectively contain 90% or more of ferrite and bainite in total, less than 5% of pearlite and carbides with a diameter of 0.5 μm or more, and less than 5% of MA phase, i.e., martensite and austenite, as the microstructure. t represents the thickness of the steel plate, and the deep layer does not include t / 4. The product of the yield strength and elongation of the thick steel plate, YS×T-El, is 16000 MPa·% or more. The thickness of the thick steel plate is more than 10 mm, [Equation 1] R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]-0.9×[Si]-1.5×[Nb] [C]*=[C]-[C]×Q Q=([Nb] / 93+[Ti] / 48) / ([C] / 12) [C], [Mn], [P], [S], [Si], [Nb], and [Ti] in Relationship 1 are weight % of the corresponding alloying elements.

2. The thick steel plate according to claim 1, wherein The thickness of the steel plate is greater than 15 mm.

3. The thick steel plate according to claim 1, wherein: In the deep layer portion of the steel plate, the pearlite and carbides having a diameter of 0.5 μm or more account for 3% or less, and the MA phase accounts for 3% or less, in terms of area%.

4. The thick steel plate according to claim 1, wherein In the surface layer portion of the steel plate, the bainite accounts for 20% or less, the pearlite and carbides having a diameter of 0.5 μm or more account for less than 2% by area, and the MA phase accounts for 3% or less.

5. The thick steel plate according to claim 1, wherein The difference between the average hardness value and the maximum hardness value of the hardness values ​​measured at intervals of 0.5 mm from a position 0.5 mm directly below the surface of the test piece to a position 0.5 mm directly below the surface of the other side, taking an arbitrary line perpendicular to the thickness section of the steel plate as a reference, is less than 20 Hv.

6. A method for manufacturing a thick steel plate, comprising the following steps: reheating a steel slab, wherein the steel slab comprises, in terms of weight%, C: 0.05-0.15%, Si: 0.01-1.0%, Mn: 1.0-2.0%, Cr: 0.005-1.0%, Al: 0.01-0.1%, P: 0.001-0.02%, S: 0.001-0.01%, N: 0.001-0.01%, Ti: 0.005-0.11%, Nb: 0.005-0.07%, and the balance is Fe and unavoidable impurities, and the steel slab has an R value defined in the following relational formula 1 of 0.3 to 1.0; a hot rolling step of hot rolling the reheated steel slab at a temperature range of 800-1150° C. at a reduction rate of 20-50% to a thickness of 10 mm or more, and terminating the rolling at a temperature range of Tn-50 to Tn defined in the following equation 2; The hot-rolled steel sheet is subjected to the CR defined in the following equation 3. 最小 The above cooling rate is used for primary cooling, cooling to a temperature range of 450-550°C, and then winding; and The coiled steel plate is subjected to secondary cooling, [Equation 1] R=[C]*+0.7×[Mn]+8.5×[P]+7.5×[S]-0.9×[Si]-1.5×[Nb] [C]*=[C]-[C]×Q Q=([Nb] / 93+[Ti] / 48) / ([C] / 12) [C], [Mn], [P], [S], [Si], [Nb] and [Ti] in the above-mentioned relational formula 1 are the weight percentages of the corresponding alloying elements. [Equation 2] Tn=730+92×[C]+70×[Mn]+45×[Cr]+650×[Nb]+410×[Ti]-80×[Si]-1.4×(t-8) The unit of Tn in the above relational expression 2 is °C, and [C], [Mn], [Cr], [Nb], [Ti] and [Si] are the weight % of the corresponding alloying elements. In the above equation 2, t is the thickness of the final rolled plate, where the unit of thickness is mm. [Equation 3] CR 最小 =76.6-157×[C]-25.2×[Si]-14.1×[Mn]-27.3×[Cr]+61×[Ti]+448×[N b] The CR of the relation 3 最小 The unit is °C / s, and [C], [Si], [Mn], [Cr], [Ti], and [Nb] are the weight % of the corresponding alloying elements.

7. The method for manufacturing a thick steel plate according to claim 6, wherein: The reheating is performed in a temperature range of 1200-1350°C. 8 . The method for manufacturing a thick steel plate according to claim 6 , wherein a cooling rate during the primary cooling is 80° C. / second or less.

9. The method for manufacturing a thick steel plate according to claim 6, wherein: During the secondary cooling, air cooling or water cooling is performed to a temperature range from room temperature to 200°C.

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