KR20240100530A - Steel plate and method for manufacturing the same - Google Patents

Abstract

The present invention relates to a steel plate and a method of manufacturing the same, and more specifically, to a steel plate for an extremely thick pressure vessel with excellent cold formability and anti-lamellar tearing quality and a method of manufacturing the same.

Description

Steel plate and its manufacturing method {STEEL PLATE AND METHOD FOR MANUFACTURING THE SAME}

The present invention relates to a steel plate and a method of manufacturing the same, and more specifically, to a steel plate for an extremely thick pressure vessel with excellent cold formability and anti-lamellar tearing quality and a method of manufacturing the same.

Recently, due to the enlargement of crude oil refining and storage facilities and the demand for large-capacity storage, the demand for thicker steel materials used in these facilities has been continuously increasing. When manufacturing large structures, there is a trend to reduce the concentration of impurities such as non-metallic inclusions or segregation or to control cracks and voids on the surface and inside the material to the limit in order to improve the internal and external soundness of the steel material.

In particular, in the case of extremely thick materials exceeding 100 mm in thickness, the rolling reduction ratio is not high compared to thin materials, so the unsolidified shrinkage voids generated during continuous rolling or casting are not sufficiently compressed during the rough rolling process, resulting in residual voids in the center of the product. It remains as These residual voids act as the starting point of cracks when the structure is subjected to thickness axial stress, which can eventually cause damage to the entire facility in the form of lamellar tearing. Therefore, in the pre-rolling stage, a process is necessary to sufficiently compress the center void to ensure that no residual void exists.

Patent Document 1 related to this corresponds to a reduction technology in the heavy plate rough rolling process, and determines the limit reduction rate for each thickness at which plate bite occurs from the reduction rate for each pass set to be close to the design allowance (load and torque) of the rolling mill. Technology to distribute the reduction rate by adjusting the index of the thickness ratio for each pass to secure the target thickness of the roughing mill, and technology to modify the reduction rate to prevent plate bite based on the limit reduction rate for each thickness. It provides a manufacturing method that can apply an average reduction ratio of about 27.5% in the final three passes of rough rolling based on a thickness of 80mm. However, in the case of the above rolling method, the average reduction rate over the entire product thickness is measured, and it has the disadvantage of making it difficult to apply high strain to the center of the extremely thick material with a maximum thickness of 233 mm where residual voids exist.

One of the other methods of manufacturing extremely thick products is to use a forging machine with a higher effective deformation per pass than a rolling mill. Patent Document 2 stands the continuous slab extracted from the heating furnace vertically, gives a total width forging reduction of 400 mm or more, and performs a width forging pass with a reduction amount within 2 passes, which is a condition within the buckling limit reduction amount. By eliminating the directional edge and center pores and increasing the center strain rate, there is an advantage in that the residual voids in the center, which were a problem in Patent Document 1, can be effectively compressed, so the lamellar tearing resistance quality of the product can be improved when forging is applied.

However, there is a disadvantage that surface defects may occur due to local strain concentration during the width forging process. In particular, if surface or subsurface defects exist in the cast steel before forging, the defects may propagate during the forging process and the surface quality of the product after rolling may deteriorate.

On the other hand, in Patent Document 3, a material provided with a predetermined alloy composition is heated to 1200 to 1350 ° C., hot forging is performed with a cumulative reduction of 25% or more, heated to a temperature of not less than Ac3 and not more than 1200 ° C., and the cumulative reduction is 25% or more. Perform hot rolling to increase the content to 40% or more, reheat to a temperature of not less than Ac3 point but not more than 1050°C, rapidly cool from a temperature of not less than Ac3 point to a temperature of 350°C or less or less than Ar3 point, and then heat it to a temperature of 450°C to 700°C. It is disclosed that a thick high-strength steel plate of 100 mmt or more with a yield strength of 620 MPa or more can be manufactured through a furnace tempering process. However, in the case of the above-mentioned ultra-high strength steel sheet, not only is it vulnerable to surface cracks during casting due to its high carbon equivalent (Ceq) and hardenability index, but also it is rapidly cooled in the austenite single phase region, causing a local quenching structure of bainite or martensite in the surface layer of the plate. Because it can be formed, it has the disadvantage that it cannot be used in storage tanks and pressure vessels with excellent cold formability.

Korean Patent Publication No. 10-2012-0075246 (published on July 6, 2012) Korean Patent Publication No. 10-2012-0074039 (published on July 5, 2012) Korean Patent Publication No. 10-2017-7019203 (published on January 15, 2016)

According to one embodiment of the present invention, it is intended to provide a steel plate and a method of manufacturing the same.

According to one embodiment of the present invention, the object is to provide a steel plate for an extremely thick pressure vessel with excellent cold formability and anti-lamellar tearing quality, and a method for manufacturing the same.

The object of the present invention is not limited to the above-described contents. A person skilled in the art will have no difficulty in understanding the additional problems of the present invention from the overall content of the present specification.

According to one embodiment of the present invention, in weight percent, carbon (C): 0.10 to 0.25%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.0 to 2.0%, aluminum (Al): 0.005 to 0.1. %, phosphorus (P): 0.010% or less, sulfur (S): 0.0015% or less, niobium (Nb): 0.001 to 0.03%, vanadium (V): 0.001 to 0.03%, titanium (Ti): 0.001 to 0.03%, Chromium (Cr): 0.01~0.20%, Molybdenum (Mo): 0.01~0.15%, Copper (Cu): 0.01~0.50%, Nickel (Ni): 0.05~0.50%, Calcium (Ca): 0.0005~0.0040%, Contains the balance Fe and inevitable impurities,

The microstructure is the surface layer, which is the area from the surface to 1/8 point in the thickness direction, and contains more than 80% polygonal ferrite, and the remaining area is the area from 1/8 point to 1/2 point in the thickness direction. The phosphorus center area is % and contains more than 50% of bainite and residual ferrite,

The difference between the porosity of the area from the surface to point 3/8 in the thickness direction and the porosity of the area from point 3/8 to 5/8 in the thickness direction is 0.1 mm ³ /g or less,

A steel plate having a hardness value of 200 to 220 HB in the surface layer can be provided.

The surface layer may include one or more types of pearlite and bainite as a residual structure.

The steel sheet may contain at least 10 types of fine NbC, NbCN, VC, or CVN precipitates with a diameter of 5 to 50 nm per 1 μm ² .

The steel sheet may have a tensile strength of 510 to 690 MPa, a thickness direction elongation (ZRA) of 35% or more, and a Charpy impact absorption energy value of 50 J or more at -50°C.

When the steel plate is subjected to a 180° bending test at room temperature, the maximum depth of surface cracks may be 1 μm or less.

The steel plate may have a thickness of 133 to 233 mm.

According to one embodiment of the present invention, in weight percent, carbon (C): 0.10 to 0.25%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.0 to 2.0%, aluminum (Al): 0.005 to 0.1. %, phosphorus (P): 0.010% or less, sulfur (S): 0.0015% or less, niobium (Nb): 0.001 to 0.03%, vanadium (V): 0.001 to 0.03%, titanium (Ti): 0.001 to 0.03%, Chromium (Cr): 0.01~0.20%, Molybdenum (Mo): 0.01~0.15%, Copper (Cu): 0.01~0.50%, Nickel (Ni): 0.05~0.50%, Calcium (Ca): 0.0005~0.0040%, Primary reheating of the steel slab containing residual Fe and unavoidable impurities;

Forging the first reheated steel slab at a cumulative reduction rate of 35 to 65% and a strain rate of 1 to 4/s;

Secondary reheating of the forged steel slab;

Hot rolling the secondary reheated steel slab at a finish rolling temperature of 900 to 1100°C;

Heating the hot-rolled steel sheet to a temperature range of 820 to 900°C and maintaining it for 10 to 40 minutes, followed by primary cooling at an average cooling rate of 0.1 to 5°C/s to 700°C based on the surface temperature of the steel sheet;

Secondary cooling the primarily cooled steel sheet to room temperature based on the surface temperature of the steel sheet at an average cooling rate of 10°C/s or more; and

A tempering heat treatment step of heating the secondary cooled steel sheet to a temperature range of 550 to 700 ° C and maintaining it for 5 to 60 minutes,

It is possible to provide a method of manufacturing a steel sheet in which the cumulative reduction rate below the recrystallization temperature during forging is 20% or less.

The first reheating step is performed at a temperature range of 1100 to 1300°C,

The second reheating step can be performed at a temperature range of 1000 to 1200°C.

During the first reheating, the thickness of the steel slab is 650 to 750 mm,

After the above forging, the thickness of the steel slab is 350~450mm,

After the hot rolling, the thickness of the steel sheet may be 133 to 233 mm.

According to one embodiment of the present invention, a steel plate and a manufacturing method thereof can be provided.

According to one embodiment of the present invention, a steel plate for an extremely thick pressure vessel with excellent cold formability and anti-lamellar tearing quality and a method for manufacturing the same can be provided.

According to one embodiment of the present invention, it is possible to provide a steel plate for an extremely thick pressure vessel with excellent cold formability and anti-lamellar tearing quality that can be used in petrochemical manufacturing facilities, storage tanks, etc., and a method for manufacturing the same.

Figure 1 is a microstructure photograph of the surface layer and 1/4 point in the thickness direction of Inventive Example 1 according to an embodiment of the present invention.
Figure 2 is a microstructure photograph of the surface layer and 1/4 point in the thickness direction of Comparative Example 6, which deviates from an embodiment of the present invention.

Below, preferred embodiments of the present invention will be described. Embodiments of the present invention may be modified in various forms, and the scope of the present invention should not be construed as limited to the embodiments described below. These embodiments are provided to explain the present invention in more detail to those skilled in the art.

Hereinafter, the present invention will be described in detail.

Below, the steel composition of the present invention will be described in detail.

In the present invention, unless otherwise specified, the % indicating the content of each element is based on weight.

The steel sheet according to an embodiment of the present invention has, in weight percent, carbon (C): 0.10 to 0.25%, silicon (Si): 0.05 to 0.50%, manganese (Mn): 1.0 to 2.0%, aluminum (Al): 0.005. ~0.1%, phosphorus (P): 0.010% or less, sulfur (S): 0.0015% or less, niobium (Nb): 0.001~0.03%, vanadium (V): 0.001~0.03%, titanium (Ti): 0.001~0.03 %, chromium (Cr): 0.01~0.20%, molybdenum (Mo): 0.01~0.15%, copper (Cu): 0.01~0.50%, nickel (Ni): 0.05~0.50%, calcium (Ca): 0.0005~0.0040 %, may include balance Fe and unavoidable impurities.

Carbon (C): 0.10~0.25%

Carbon (C) is the most important element in securing basic strength, so it needs to be contained in steel within an appropriate range. To achieve this added effect, more than 0.10% of carbon (C) can be added. According to one embodiment of the present invention, 0.12% or more of carbon (C) may be added. On the other hand, if the content exceeds a certain level, during QT heat treatment, the fraction of low-temperature transformed structure may increase and the base material strength and hardness may be excessively exceeded, which may result in the formation of a low-temperature structure in the surface structure rather than stretched ferrite. Uniform elongation may decrease. Therefore, in the present invention, the upper limit of carbon (C) content can be limited to 0.25%. According to one embodiment of the present invention, the upper limit may be 0.20%.

Silicon (Si): 0.05~0.50%

Silicon (Si) is a substitutional element that improves the strength of steel through solid solution strengthening and has a strong deoxidation effect, so it is an essential element in the production of clean steel. Therefore, silicon (Si) may be added in an amount of 0.05% or more. According to one embodiment of the present invention, 0.20% or more may be added. On the other hand, when a large amount of silicon (Si) is added, it creates an MA phase and excessively increases the strength of the ferrite matrix, which can lead to deterioration in the surface quality of extremely thick products, so the upper limit of its content can be limited to 0.5%. . According to one embodiment of the present invention, the upper limit may be 0.40%.

Manganese (Mn): 1.0~2.0%

Manganese (Mn) is a useful element that improves strength through solid solution strengthening and improves hardenability to generate a low-temperature transformation phase. Therefore, in order to secure a tensile strength of 450 MPa or more, 1.0% or more can be added. According to one embodiment of the present invention, the manganese (Mn) content may be 1.1% or more. On the other hand, as the manganese (Mn) content increases, manganese (Mn) forms MnS, an elongated non-metallic inclusion with S, which reduces toughness and acts as a factor in reducing elongation when stretched in the thickness direction, thereby reducing the anti-lamellar tearing quality. This may be a factor in its rapid decline. Therefore, the manganese (Mn) content can be limited to 2.0% or less. According to one embodiment of the present invention, the content may be 1.5% or less.

Aluminum (Al): 0.005~0.1%

Aluminum (Al), along with Si, is one of the powerful deoxidizers in the steelmaking process, and can be added in amounts of 0.005% or more to achieve this effect. According to one embodiment of the present invention, the lower limit of aluminum (Al) content may be 0.01%. On the other hand, when the aluminum (Al) content is excessive, the fraction of Al ₂ O ₃ in the oxidizing inclusions generated as a result of deoxidation increases excessively, making their size coarse, and removing the inclusions during refining becomes difficult. This can be a factor that reduces anti-ramellar tearing characteristics. Therefore, the aluminum (Al) content can be limited to 0.1% or less. According to one embodiment of the present invention, the aluminum (Al) content may be 0.07% or less.

Phosphorus (P): 0.010% or less

Phosphorus (P) is an element that causes embrittlement at grain boundaries or forms coarse inclusions, so in order to improve brittle crack propagation resistance, phosphorus (P) can be limited to 0.010% or less. However, considering the level of unavoidable content during steel manufacturing, 0% is excluded.

Sulfur (S): 0.0015% or less

Sulfur (S) is an element that causes embrittlement at grain boundaries or forms coarse inclusions, so in order to improve brittle crack propagation resistance, sulfur (S) can be limited to 0.0015% or less. However, considering the level of unavoidable content during steel manufacturing, 0% is excluded.

Niobium (Nb): 0.001~0.03%

Niobium (Nb) is an element that improves the strength of the base material by precipitating in the form of NbC or NbCN. In addition, Nb dissolved in solid solution when reheated to a high temperature precipitates very finely in the form of NbC during rolling, which has the effect of suppressing recrystallization of austenite and refining the structure. Therefore, niobium (Nb) may be added in an amount of 0.001% or more. According to one embodiment of the present invention, it may be 0.005% or more. On the other hand, when niobium (Nb) is added excessively, undissolved niobium (Nb) is generated in the form of TiNb (C, N), which is a factor that inhibits the anti-lamellar tearing characteristics, so the upper limit of the content is 0.03%. It can be limited to . According to one embodiment of the present invention, the niobium (Nb) content may be 0.02% or less.

Vanadium (V): 0.001~0.03%

Since vanadium (V) is almost completely redissolved upon reheating, the strengthening effect due to precipitation or solid solution during subsequent rolling is minimal, but it has the effect of improving strength by precipitating as very fine carbonitride during subsequent heat treatment such as PWHT. To fully obtain this effect, vanadium (V) can be added in an amount of 0.001% or more. According to one embodiment of the present invention, the lower limit of vanadium (V) content may be 0.01%. On the other hand, if the content is excessive, the strength and hardness of the base material and weld zone are excessively increased, which may act as a factor such as surface cracks during pressure vessel processing, and the manufacturing cost rises rapidly, which is not commercially advantageous. Therefore, the vanadium (V) content may be 0.03% or less. According to one embodiment of the present invention, the vanadium (V) content may be 0.02% or less.

Titanium (Ti): 0.001~0.03%

Titanium (Ti) is a component that greatly improves low-temperature toughness by precipitating as TiN upon reheating and inhibiting the growth of grains in the base metal and weld heat-affected zone. Therefore, it can be added in an amount of 0.001% or more to achieve this addition effect. On the other hand, if titanium (Ti) is added excessively, low-temperature toughness may be reduced due to clogging of the playing nozzle or crystallization at the center, and it combines with N to form coarse TiN precipitates in the center of the thickness, thereby reducing the elongation of the product. , the anti-lamellar tearing properties of the final material may be reduced. Therefore, the titanium (Ti) content may be 0.03% or less. The titanium (Ti) content according to one embodiment of the present invention may be 0.025% or less, and may be 0.018% or less.

Chromium (Cr): 0.01~0.20%

Chromium (Cr) increases hardenability and increases yield and tensile strength by forming a low-temperature transformation structure, and has the effect of preventing a decrease in strength by slowing down the decomposition rate of cementite during tempering after quenching or heat treatment after welding. For this effect, more than 0.01% of chromium (Cr) can be added. On the other hand, when the chromium (Cr) content is excessive, the size and fraction of Cr-Rich coarse carbides such as M ₂₃ C ₆ increase and the impact toughness of the product decreases, and the solid solubility of Nb in the product and the fraction of fine precipitates such as NbC decrease. As this decreases, a decrease in the strength of the product may be a problem, so the upper limit of the content can be limited to 0.20%. According to one embodiment of the present invention, the upper limit of chromium (Cr) content may be 0.15%.

Molybdenum (Mo): 0.01~0.15%

Molybdenum (Mo) is an element that increases grain boundary strength and has a significant solid solution strengthening effect in ferrite, effectively contributing to increasing the strength and ductility of products. In addition, molybdenum (Mo) has the effect of preventing a decrease in toughness due to grain boundary segregation of impurities such as P. For this effect, more than 0.01% molybdenum (Mo) can be added. However, molybdenum (Mo) is an expensive element and if added excessively, the manufacturing cost can increase significantly, so the upper limit of its content can be limited to 0.15%.

Copper (Cu): 0.01~0.50%

Copper (Cu) is an advantageous element in the present invention because it can not only greatly improve the strength of the matrix phase through solid solution strengthening in ferrite, but also has the effect of suppressing corrosion in a wet hydrogen sulfide atmosphere. For this effect, it may contain more than 0.01% copper (Cu). A more desirable copper (Cu) content may be 0.03% or more. However, if the content of copper (Cu) is excessive, the possibility of causing star cracks on the surface of the steel sheet increases, and as it is an expensive element, the manufacturing cost increases significantly. Therefore, the present invention limits the upper limit of the content to 0.50%. can do. According to one embodiment of the present invention, the upper limit may be 0.30%.

Nickel (Ni): 0.05~0.50%

Nickel (Ni) is an important element in improving impact toughness by increasing stacking faults at low temperatures, facilitating cross slip of dislocations, and improving strength by improving hardenability. For this effect, more than 0.05% nickel (Ni) may be added. According to one embodiment of the present invention, the nickel (Ni) content may be 0.10% or more. On the other hand, if excessive nickel (Ni) is added, the manufacturing cost may increase due to the high cost, so the upper limit of the content can be limited to 0.50%. According to one embodiment of the present invention, the upper limit of nickel (Ni) content may be 0.30%.

Calcium (Ca): 0.0005~0.0040%

When calcium (Ca) is added after deoxidation by Al, it combines with S forming MnS inclusions to suppress the production of MnS, and at the same time forms spherical CaS, which causes cracks due to hydrogen-induced cracking. It has a suppressing effect. In order to sufficiently form S contained as an impurity into CaS, calcium (Ca) can be added in an amount of 0.0005% or more. However, if the addition amount is excessive, the calcium (Ca) remaining after forming CaS combines with oxygen to form coarse oxidative inclusions, which are elongated or broken during rolling, causing a problem of deterioration in anti-ramellar tearing properties. The upper limit of (Ca) content can be limited to 0.0040%.

The steel material of the present invention may contain remaining iron (Fe) and inevitable impurities in addition to the composition described above. Since unavoidable impurities may be unintentionally introduced during the normal manufacturing process, they cannot be excluded. Since these impurities are known to anyone skilled in the field of steel manufacturing, all of them are not specifically mentioned in this specification.

Below, the steel microstructure of the present invention will be described in detail.

In the present invention, unless specifically stated otherwise, the % indicating the fraction of microstructure is based on area.

The microstructure of the steel sheet according to an embodiment of the present invention is that the surface layer, which is the area from the surface to 1/8 point in the thickness direction, contains 80% or more polygonal ferrite in area%, and the remaining area, which is the area up to 1/8 of the point in the thickness direction, is 1/8 of the area. The central area, which is the area from point 8 to 1/2 point, may contain 50% or more of bainite and residual ferrite in terms of area percentage.

In the present invention, in order to improve cold workability through the formation of a soft microstructure in the surface layer, the surface layer portion, which is the area from the surface to 1/8 point in the thickness direction, can be limited to contain 80% or more of polygonal ferrite. In one embodiment of the present invention, polygonal ferrite may be 90% or more. In addition to polygonal ferrite, the remaining structure may include one or more of pearlite and bainite.

The remaining area, the central area from the 1/8 point to the 1/2 point, may contain more than 50% of bainite to ensure the strength proposed in the present invention, and may include ferrite as the remaining structure.

According to one embodiment of the present invention, the difference between the porosity of the area from the surface to the 3/8 point in the thickness direction and the porosity of the area from the 3/8 point to the 5/8 point in the thickness direction may be 0.1 mm ³ /g or less. there is.

That is, the value obtained by subtracting the porosity of the area from the surface to the 3/8th point in the thickness direction from the porosity of the area from 3/8 to 5/8 in the thickness direction of the steel sheet may be 0.1 mm ³ /g or less. Here, the porosity can be obtained by measuring the density (g/mm ³ ) and taking the reciprocal (mm ³ /g).

In the present invention, the difference in porosity between the surface layer and the center can be limited for anti-ramellar tearing characteristics. If the difference in porosity exceeds 0.1 mm ³ /g, there is a problem of deteriorating the thickness direction elongation (ZRA) because the residual voids act as crack initiation points that lower the elongation during the thickness direction tensile test. According to one embodiment of the present invention, the difference may be 0.05mm ³ /g or less. According to one embodiment of the present invention, the difference may be 0.03mm ³ /g or less. There is no need to specifically limit the lower limit of the difference, but in one embodiment, the lower limit may be 0 mm ³ /g.

According to one embodiment of the present invention, one or more types of fine NbC, NbCN, VC, or CVN precipitates with a diameter of 5 to 50 nm may be 10 or more per 1 μm ² .

In the present invention, the amount of precipitates can be limited to secure strength through precipitation strengthening. If the number of precipitates is less than 10 per 1 μm ² , the strengthening effect is insufficient, making it difficult to secure the appropriate strength proposed in the present invention. On the other hand, if the diameter of the precipitate is too small (less than 5 nm), it is difficult to secure strength similar to the effect of a small number of precipitates, and if it exceeds 50 nm, the pinning effect is also reduced, so the strength improvement effect is reduced. Additionally, if the precipitate size exceeds 50 nm, the coarse precipitate may act as a brittle fracture initiation point during the impact test, so it is appropriate that the diameter of the precipitate is 5 to 50 nm.

The steel plate according to an embodiment of the present invention has a thickness of 133 to 233 mm, a hardness value of the surface layer of 200 to 220 HB, a tensile strength of 510 to 690 MPa, a thickness direction elongation (ZRA) of 35% or more, and -50 The Charpy impact absorption energy value at ℃ is more than 50J, and when bending 180° at room temperature, the maximum depth of surface cracks can be less than 1μm.

Below, the steel manufacturing method of the present invention will be described in detail.

A steel plate according to an embodiment of the present invention can be manufactured by primary reheating, forging, secondary reheating, hot rolling, primary cooling, secondary cooling, and tempering of a steel slab that satisfies the alloy composition of the present invention.

1st reheat

A steel slab with a thickness of 650 to 750 mm that satisfies the alloy composition of the present invention can be first reheated at a temperature range of 1100 to 1300°C.

The complex carbonitrides of Ti or Nb or TiNb(C,N) coarse crystals formed during casting are re-dissolved, and the austenite is heated and maintained above the recrystallization temperature before forging to homogenize the structure and ensure a sufficiently high forging end temperature. Therefore, in order to minimize surface cracks that may occur during the forging process, primary reheating can be performed in a temperature range of 1100°C or higher. On the other hand, when reheating a steel slab at an excessively high temperature, problems may occur due to oxidized scale at high temperatures and manufacturing costs may increase excessively due to increased costs due to heating and maintenance. Therefore, the upper limit of the steel slab heating temperature must be set. It can be limited to 1300℃.

According to one embodiment of the present invention, during primary reheating, the thickness of the steel slab may be 650 to 750 mm. According to one embodiment of the present invention, the thickness of the steel slab may be 700 to 750 mm.

minor

The primary reheated steel slab can be forged at a cumulative reduction rate of 35 to 65% and a strain rate of 1 to 4/s.

Forging is a step in which heated steel slabs are processed into the final shape of the desired intermediate material. Since high-strain, low-speed forging is essential to sufficiently compress the voids, forging can be performed under the conditions of a cumulative reduction rate of 35 to 65% and a strain rate of 1/s to 4/s.

In the present invention, strain rate means strain rate per unit time, and the unit means %/s.

If the cumulative reduction ratio of forging is less than 35%, the remaining voids in the steel slab cannot be sufficiently compressed, resulting in residual voids, which may deteriorate the anti-lamellar tearing characteristics of the product. The cumulative reduction rate of forging according to an embodiment of the present invention may be 40% or more. If the cumulative reduction ratio exceeds 65% and the number of forging passes increases, the surface temperature continues to decrease and surface cracks may occur at low temperatures.

On the other hand, if the dislocation density is recovered or the cumulative reduction ratio below the non-recrystallization temperature that is not offset by recrystallization exceeds 20%, 20% is the uniform elongation value of the steel surface layer at about 850°C to 1150°C where forging is performed. exceeds , and therefore, the uniform elongation of the surface is extremely reduced due to work hardening of overlapping dislocations, and surface cracks may occur during the forging process. Therefore, the cumulative reduction rate below the recrystallization temperature can be limited to 20% or less. According to one embodiment of the present invention, it may be 15% or less.

During forging, if the strain rate is less than 1/s, forging productivity may decrease and the temperature of the surface layer may decrease excessively during the forging process, so surface cracks may occur during the final forging process. On the other hand, if the strain rate exceeds 4/s, surface quality may deteriorate due to excessive work hardening.

According to one embodiment of the present invention, the thickness of the steel slab after forging may be 350 to 450 mm.

Secondary reheating

The forged steel slab can be secondarily reheated to a temperature range of 1000 to 1200°C.

The complex carbonitrides of Ti or Nb or TiNb (C, N) coarse crystals formed during casting are re-dissolved, the structure is homogenized by heating and maintaining the austenite to the recrystallization temperature or higher before hot rolling, and the rolling end temperature is maintained. Reheating can be performed in a temperature range of 1000°C or higher to ensure that the temperature is high enough to minimize crushing of inclusions during the rolling process. On the other hand, if the slab is heated to an excessively high temperature, problems may occur due to oxidized scale at high temperature and manufacturing costs may increase excessively due to increased costs due to heating and maintenance. Therefore, the upper limit of the secondary heating temperature is set to 1200. It can be limited to ℃.

hot rolling

The secondary reheated steel slab can be hot rolled at a finish rolling temperature of 900 to 1100°C.

During hot rolling, if the finish rolling temperature is less than 900°C, the deformation resistance value increases excessively as the temperature drops, making it difficult to sufficiently refine the austenite grains in the center of the product's thickness direction, and as a result, the anti-ramellar tearing characteristics of the final product deteriorate. You can. On the other hand, when the finish rolling temperature exceeds 1100°C, there is a risk that the strength and impact toughness may be deteriorated because the austenite grains become excessively coarse.

According to one embodiment of the present invention, the thickness of the steel sheet after hot rolling may be 133 to 233 mm.

Primary cooling

The hot-rolled steel sheet can be heated to a temperature range of 820 to 900°C and maintained for 10 to 40 minutes, and then first cooled to 700°C based on the surface temperature of the steel sheet at an average cooling rate of 0.1 to 5°C/s.

In the present invention, the primary cooling can be performed to form the desired polygonal ferrite from the surface of the steel material to an area corresponding to 1/8 of the point in the thickness direction. During primary cooling, the temperature and cooling rate are based on the surface temperature of the steel plate.

When reheating, if the temperature is less than 820℃ or the holding time is less than 10 minutes, the re-dissolution of the carbides generated during cooling after rolling or the impurity elements segregated at the grain boundaries does not occur smoothly, and the thickness direction elongation (ZRA) and Low-temperature toughness may be greatly reduced. On the other hand, when the temperature exceeds 900°C or the holding time exceeds 40 minutes, internal lamellae occur due to austenite coarsening and coarsening of precipitated phases such as Nb(C,N) and V(C,N). Tearing quality may deteriorate.

During primary cooling, if the average cooling rate is less than 0.1°C/s, the waiting time in the air becomes too long, which may prolong manufacturing time and reduce productivity. On the other hand, if the cooling rate exceeds 5°C/s, polygonal ferrite is not formed and bainite or martensite, which is a low-temperature transformation structure, is formed, and the uniform elongation of the surface layer of the final product is lowered, so cold workability may rapidly deteriorate.

secondary cooling

The primary cooled steel sheet may be secondary cooled to room temperature based on the surface temperature of the steel sheet at an average cooling rate of 10°C/s or more.

The secondary cooling may be performed to ensure that the microstructure of areas other than the surface layer includes bainite or an abnormal structure of bainite and ferrite through strong cooling.

During the secondary cooling, the average cooling rate is based on the surface temperature of the steel plate. If the average cooling rate is less than 10°C/s, it may be difficult to obtain the above-described low-temperature transformed structure. In the present invention, there is no particular limitation on the upper limit of the average cooling rate during the secondary cooling, but it may be limited to 200°C/s or less. Meanwhile, the secondary cooling can be achieved through quenching, slowing down the sheeting speed of the steel material, and increasing the flow rate of the sprayed water.

tempering heat treatment

The secondary cooled steel sheet can be subjected to tempering heat treatment by heating it to a temperature range of 550 to 700°C and maintaining it for 5 to 60 minutes.

Through the tempering heat treatment, the dislocation density of bainite or a mixed structure of bainite and ferrite, which is a low-temperature transformation structure, can be reduced, and the strength and toughness can be improved by diffusing carbon to a short range.

If the tempering heat treatment temperature is less than 550°C, diffusion of carbon is not sufficient, so the strength may increase excessively and the toughness may decrease. On the other hand, when the temperature exceeds 700°C, fresh martensite is formed due to reverse transformation at a temperature above Ac1, and impact toughness and cold workability may be extremely deteriorated.

If the tempering heat treatment time is less than 5 minutes, there is not enough time for carbon to sufficiently diffuse during the tempering process, so the strength may be excessively exceeded, resulting in reduced toughness and cold workability, and may fall outside the appropriate strength range required by the present invention. there is. On the other hand, if the tempering heat treatment time exceeds 60 minutes, the cementite may spheroidize due to excessive heating and the strength may rapidly decrease.

Hereinafter, the present invention will be described in more detail through examples. However, it is important to note that the examples below are only for illustrating and explaining the present invention in more detail, and are not intended to limit the scope of the present invention.

(Example)

A cast plate with a thickness of 700 mm having the alloy components shown in Table 1 was produced. According to the process conditions in Table 2, forging, hot rolling, primary and secondary cooling, and then tempering heat treatment were performed, and finally, a steel plate with a thickness of 133 mm was manufactured. At this time, the first reheating temperature was 1200°C and the second reheating temperature was 1100°C, and the reheating time and tempering time for first cooling were 30 minutes.

Steel grade Alloy composition (% by weight) C Si Mn Al P S Nb V Ti Cr Mo Cu Ni Ca A 0.16 0.35 1.32 0.03 90 10 0.017 0.007 0.011 0.06 0.09 0.05 0.19 19 B 0.14 0.31 1.43 0.027 82 11 0.011 0.009 0.007 0.08 0.08 0.07 0.23 22 C 0.13 0.32 1.35 0.031 77 10 0.018 0.012 0.005 0.12 0.1 0.05 0.21 18 D 0.16 0.35 1.22 0.028 84 9 0.017 0.023 0.01 0.1 0.09 0.1 0.22 18 E 0.15 0.4 1.21 0.03 89 9 0.015 0.019 0.011 0.09 0.06 0.08 0.19 17 F 0.28 0.35 1.21 0.03 74 8 0.019 0.009 0.009 0.12 0.08 0.09 0.25 18 G 0.03 0.31 0.3 0.03 80 10 0.013 0.009 0.005 0.2 0.08 0.07 0.19 19 H 0.16 0.35 2.35 0.03 79 9 0.0005 0.014 0.004 0.25 0.09 0.06 0.18 18 I 0.17 0.3 1.21 0.029 82 35 0.012 0.15 0.005 0.23 0.1 0.1 0.19 20 J 0.06 0.25 1.32 0.31 85 10 0.015 0.016 0.012 0.15 0.07 0.06 0.23 2

* Units for P, S, Ca and O are ppm.

Psalter
number steel grade minor hot rolling Primary cooling secondary cooling tempering Cumulative reduction rate
(%) recrystallization temperature
below
Reduction rate (%) Deformation speed
(/s) finish
rolling temperature
(℃) speed
(℃/s) speed
(℃/s) temperature
(℃) One A 45 15 1.7 915 1.3 24.5 620 2 B 47 13 2.5 930 0.9 16.6 618 3 C 51 11 3.9 1050 1.6 18.4 625 4 D 61 14 2.7 987 2.4 26.5 636 5 E 39 9 2.4 968 3.6 29.3 641 6 A 84 10 1.8 915 4.1 28.0 632 7 B 12 13 1.9 921 4.3 27.7 629 8 B 48 39 2.6 925 3.8 16.6 621 9 C 47 14 8 964 1.5 19.4 618 10 C 51 16 3.3 710 1.8 26.5 599 11 D 50 13 2.7 981 83 34.4 580 12 D 59 15 3.5 990 0.7 7.4 664 13 E 51 14 3.4 980 1.5 27.3 788 14 F 49 17 3.8 912 1.5 36 636 15 G 39 12 3.9 1043 3.2 47 629 16 H 42 11 2.9 1054 2.6 40 627 17 I 46 17 1.8 983 3.0 35 640 18 J 50 14 1.6 991 2.7 33 633

The microstructure and physical properties of each specimen prepared above were measured and shown in Table 3 below. Microstructure is obtained by collecting specimens from the surface area (the area from the surface to 1/8 point in the thickness direction) and the remaining area (the area from 1/8 point to 1/2 point in the thickness direction) using an automatic image analyzer. Measured. At this time, the observed microstructure shows polygonal ferrite (P.F), martensite (M), and bainite (B), and their area percentage means 100%.

In the present invention, the difference in porosity is expressed as the value obtained by subtracting the porosity of the area from the surface to the 3/8th point in the thickness direction from the porosity of the area from 3/8 to 5/8 in the thickness direction. Here, the porosity was obtained by measuring the density (g/mm ³ ) and taking the reciprocal (mm ³ /g).

The number of precipitates measuring 5 to 50 nm observed in the cross section of the steel was measured using TEM. NbC precipitates were confirmed through diffraction patterns and EDX mapping of NbC and VC, and the number of precipitates located in 1 μm ² was counted.

In addition, the hardness value was measured at three surface areas using a Brinell hardness tester, and the average value was expressed. Tensile strength was evaluated through a room temperature tensile test, and the impact toughness of each specimen was determined by using the average of the absorbed energy values measured three times at the corresponding temperature through the Charpy V-Notch Test.

In addition, after visually observing the surface of each specimen, grinding was performed at the point where the surface crack was formed, and the grinding length until the crack disappeared was measured as the surface crack length.

city
side
th
like river
bell microstructure Properties division surface layer
(area%) center
(area%) Porosity difference
( ^mm3 /g) precipitate
(dog/
μm ² ) Hardness
(HB) Seal
robbery
(MPa) elongation
(%) Shock
tenacity
(-50℃,
J) crack
depth
(μm) B F One A P.F. 92 8 0.002 24 202 514 72 315 0 Invention Example 1 2 B P.F. 89 11 0.002 29 215 534 59 302 0 Invention Example 2 3 C P.F. 88 12 0.013 25 207 524 64 298 0 Invention Example 3 4 D P.F. 76 24 0.015 36 205 508 64 300 0 Invention Example 4 5 E P.F. 89 11 0.007 19 203 517 59 315 0 Invention Example 5 6 A P.F. 91 9 0.012 34 287 526 66 307 19 Comparative Example 1 7 B P.F. 84 16 0.36 33 207 500 12 320 0 Comparative example 2 8 B P.F. 76 24 0.017 45 276 543 59 257 16 Comparative Example 3 9 C P.F. 88 12 0.015 27 274 529 61 289 18 Comparative example 4 10 C P.F. 86 14 0.007 19 283 693 21 14 9 Comparative Example 5 11 D M 84 16 0.006 18 274 534 54 288 15 Comparative Example 6 12 D P.F. 22 78 0.007 23 206 433 57 304 0 Comparative example 7 13 E P.F. 81 19 0.015 4 182 461 64 13 0 Comparative example 8 14 F M 100 0 0.017 41 215 798 8 12 17 Comparative Example 9 15 G P.F. 9 91 0.015 35 208 434 30 289 0 Comparative Example 10 16 H P.F. 84 16 0.009 One 210 451 61 301 0 Comparative Example 11 17 I B 98 2 0.011 89 214 687 55 25 8 Comparative Example 12 18 J P.F. 30 70 0.016 38 211 431 57 286 0 Comparative Example 13

* P.F: polygonal ferrite, M: martensite, B: bainite

As shown in Table 3, in the case of the invention example that satisfies the alloy composition and manufacturing conditions of the present invention, the microstructure characteristics proposed in the present invention were satisfied, and the physical properties desired in the present invention were also secured.

Figure 1 is a microstructure photograph of the surface layer and 1/4 point in the thickness direction of Inventive Example 1 according to an embodiment of the present invention. In the case of Invention Example 1, as shown in FIG. 1, polygonal ferrite was formed in the surface layer and 1/4 of the layer was formed of bainite, which was advantageous in securing cold formability and strength.

On the other hand, Comparative Examples 1 and 2 are examples where the cumulative reduction ratio during forging is outside the scope of the present invention. In Comparative Example 1, forging cracks occurred due to excessive reduction during forging, and the surface quality was poor even during cold bending. Additionally, the hardness value of the surface layer was excessive. In Comparative Example 2, during forging, the reduction ratio was insufficient and the remaining voids in the center of the steel sheet were not sufficiently compressed, resulting in a decrease in elongation.

Comparative Example 3 is an example in which the reduction ratio exceeds the range of the present invention below the recrystallization temperature. As a result, surface defects occurred.

Comparative Example 4 is an example in which the strain rate during forging exceeds the range proposed in the present invention. As a result, the hardness value of the surface layer was excessive, and cracks also occurred.

Comparative Example 5 is an example in which the finish rolling temperature was insufficient during hot rolling. As a result, the strength was too high, the low-temperature impact toughness was deteriorated, and the cold formability was also poor.

Comparative Example 6 is an example in which the cooling rate exceeds the range proposed in the present invention during primary cooling. As a result, martensite, a hard structure, was formed in the surface layer, which reduced cold formability and resulted in surface defects in the product during the cold bending process.

Figure 2 is a microstructure photograph of the surface layer and 1/4 point in the thickness direction of Comparative Example 6, which deviates from an embodiment of the present invention. As shown in Figure 2, it can be confirmed that the surface layer is formed of marsitete tissue.

Comparative Example 7 was a case where the cooling rate was insufficient during secondary cooling, and appropriate strength was not secured in the center, so the strength desired in the present invention was not secured.

Comparative Example 8 was a case in which the tempering temperature was excessively high, and the formation of precipitates was not smooth, and as a result, the desired strength of the present invention was not secured.

Comparative Examples 9 and 10 are examples in which the carbon content exceeds the range proposed in the present invention. For this reason, in Comparative Example 9, martensite, a hard structure, was formed in the surface layer, making it difficult to secure the desired impact toughness. Comparative Example 10 is an example in which the Mn content was also low, and it was difficult to secure appropriate strength in the center, resulting in a decrease in strength.

Comparative Example 11 is an example in which the Nb content is below the content proposed in the present invention. As a result, it was not easy to form precipitates, and the desired strength could not be secured.

Comparative Example 12 is an example in which the V content exceeds the range of the present invention. As a result, a bainite structure was formed in the surface layer, and as a result, low-temperature impact toughness was inferior.

Comparative Example 13 is an example that falls short of the C content proposed in the present invention. As a result, bainite was not properly formed in the center, and strength also decreased.

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

Claims (9)

By weight percent, carbon (C): 0.10-0.25%, silicon (Si): 0.05-0.50%, manganese (Mn): 1.0-2.0%, aluminum (Al): 0.005-0.1%, phosphorus (P): 0.010. % or less, Sulfur (S): 0.0015% or less, Niobium (Nb): 0.001 to 0.03%, Vanadium (V): 0.001 to 0.03%, Titanium (Ti): 0.001 to 0.03%, Chromium (Cr): 0.01 to 0.20 %, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, including the balance Fe and inevitable impurities. ,
The microstructure is the surface layer, which is the area from the surface to 1/8 point in the thickness direction, and contains more than 80% polygonal ferrite, and the remaining area is the area from 1/8 point to 1/2 point in the thickness direction. The phosphorus center area is % and contains more than 50% of bainite and residual ferrite,
The difference between the porosity of the area from the surface to the 3/8th point in the thickness direction and the porosity of the area from the 3/8th point to the 5/8th point in the thickness direction is 0.1mm ³ /g or less,
A steel plate with a hardness value of 200 to 220 HB in the surface layer.
In claim 1,
The surface layer portion is a steel plate containing at least one type of pearlite and bainite as a residual structure.
In claim 1,
The steel sheet contains at least 10 types of fine NbC, NbCN, VC, or CVN precipitates with a diameter of 5 to 50 nm per 1 μm ² .
In claim 1,
The steel sheet has a tensile strength of 510 to 690 MPa, a thickness direction elongation (ZRA) of 35% or more, and a Charpy impact absorption energy value of 50 J or more at -50°C.
In claim 1,
The steel sheet has a maximum surface crack depth of 1 μm or less when subjected to a 180° bending test at room temperature.
In claim 1,
The steel plate is a steel plate with a thickness of 133 to 233 mm.
By weight percent, carbon (C): 0.10-0.25%, silicon (Si): 0.05-0.50%, manganese (Mn): 1.0-2.0%, aluminum (Al): 0.005-0.1%, phosphorus (P): 0.010. % or less, Sulfur (S): 0.0015% or less, Niobium (Nb): 0.001 to 0.03%, Vanadium (V): 0.001 to 0.03%, Titanium (Ti): 0.001 to 0.03%, Chromium (Cr): 0.01 to 0.20 %, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, calcium (Ca): 0.0005 to 0.0040%, including the balance Fe and inevitable impurities. primary reheating of the steel slab;
Forging the first reheated steel slab at a cumulative reduction rate of 35 to 65% and a strain rate of 1 to 4/s;
Secondary reheating of the forged steel slab;
Hot rolling the secondary reheated steel slab at a finish rolling temperature of 900 to 1100°C;
Heating the hot-rolled steel sheet to a temperature range of 820 to 900°C and maintaining it for 10 to 40 minutes, followed by primary cooling at an average cooling rate of 0.1 to 5°C/s to 700°C based on the surface temperature of the steel sheet;
Secondary cooling the primarily cooled steel sheet to room temperature based on the surface temperature of the steel sheet at an average cooling rate of 10°C/s or more; and
A tempering heat treatment step of heating the secondary cooled steel sheet to a temperature range of 550 to 700 ° C and maintaining it for 5 to 60 minutes,
A method of manufacturing a steel plate in which the cumulative reduction rate below the recrystallization temperature during forging is 20% or less.
In claim 7,
The first reheating step is performed at a temperature range of 1100 to 1300°C,
A method of manufacturing a steel plate in which the second reheating step is performed in a temperature range of 1000 to 1200°C.
In claim 7,
During the first reheating, the thickness of the steel slab is 650 to 750 mm,
After the above forging, the thickness of the steel slab is 350~450mm,
A method of manufacturing a steel sheet where, after the hot rolling, the thickness of the steel sheet is 133 to 233 mm.