KR102722678B1 - Extra heavy gauged steel plate having excellent surface quality and lamellar tearing resistance and manufacturing method for the same - Google Patents

Abstract

According to one aspect of the present invention, a steel for an ultra-thick pressure vessel contains, in wt%, 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 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%, the remainder includes iron and unavoidable impurities, the elongation in the thickness direction (ZRA) is 35% or more, the tensile strength is 450 to 620 MPa, and the maximum surface crack depth may be 2 μm or less.

Description

{EXTRA HEAVY GAUGED STEEL PLATE HAVING EXCELLENT SURFACE QUALITY AND LAMELLAR TEARING RESISTANCE AND MANUFACTURING METHOD FOR THE SAME}

The present invention relates to steel that can be used in petrochemical manufacturing facilities, storage tanks, etc., and more specifically, to steel for pressure vessels having excellent surface quality and resistance to lamellar tearing and a method for 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 thick steel used in these facilities has been continuously increasing. In manufacturing large structures, the trend is to reduce the concentration of impurities such as non-metallic inclusions or segregation, and to control cracks and pores on the surface and inside the material to the extreme in order to improve the internal and external soundness of the steel.

In particular, in the case of extremely thick materials exceeding 100 mmt in thickness, the rolling reduction ratio is not high compared to thin materials, so the unsolidified shrinkage voids that occur during casting or rolling are not sufficiently compressed during the rough rolling process and remain in the form of residual voids in the center of the product.

These residual voids act as crack initiation points when the structure is subjected to axial stress along the thickness direction, and can eventually cause damage to the entire equipment in the form of lamellar tearing. Therefore, a process is necessary to sufficiently compress the center voids to ensure that no residual voids exist prior to rolling.

Patent Document 1 related to this corresponds to a technology for applying pressure in a thick plate roughing process, and it uses a technology for determining a thickness-specific limit reduction ratio at which plate biting occurs from a pass-specific reduction ratio set to be close to the design allowable value (load and torque) of the rolling mill, a technology for distributing the reduction ratio by adjusting the index of the thickness ratio of each pass to secure the target thickness of the roughing mill, and a technology for modifying the reduction ratio so that plate biting does not occur based on the thickness-specific limit reduction ratio, thereby providing a manufacturing method that can apply an average reduction ratio of approximately 27.5% in the final three passes of roughing with 80 mmt as a standard. However, in the case of the above rolling method, since it measures the average reduction ratio of the entire product thickness, there is a disadvantage in that it is difficult to apply high strain to the center of an extremely thick material with a maximum thickness of 233 mmt where residual voids exist.

Another method for manufacturing extremely thick materials is to utilize a forging machine having a higher effective strain per pass than a rolling mill. Patent Document 2 discloses a method in which a continuous casting slab extracted from a heating furnace is vertically erected to provide a total width forging reduction amount of 400 mm or more and performs width forging passes with a reduction amount of within two passes, which is within the buckling limit reduction amount, to eliminate pores in the width direction edge portion and center portion and increase center strain. Since the center residual pores, which were a problem in Patent Document 1, can be effectively compressed, the quality of the product's lamella tearing resistance can be improved when forging is applied.

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

Meanwhile, Patent Document 3 discloses that a thick-walled high-strength steel plate having a yield strength of 620 MPa or more and a thickness of 100 mmt or more can be manufactured through a process of heating a material provided with a predetermined alloy composition to 1200-1350°C, performing hot forging with a cumulative reduction amount of 25% or more, heating the material to a temperature of Ac3 point or higher and 1200°C or lower, performing hot rolling with a cumulative reduction amount of 40% or more, reheating the material to a temperature of Ac3 point or higher and 1050°C or lower, rapidly cooling the material from the temperature of Ac3 point or higher to a temperature of 350°C or lower or Ar3 point or lower, whichever is lower, and tempering the material at a temperature of 450°C to 700°C.

However, in the case of the ultra-high strength steel plate described above, the carbon equivalent (Ceq) and hardenability index are high, so not only is it vulnerable to surface cracks during casting, but it also has the disadvantage of not being able to be used in storage tanks and pressure vessels manufactured by hot working at temperatures of 800 to 950°C.

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 aspect of the present invention, a steel material for an ultra-thick pressure vessel having excellent surface quality and lamella tearing quality and a method for manufacturing the same can be provided.

The subject matter of the present invention is not limited to the above-described content. Those skilled in the art will have no difficulty in understanding additional subjects of the present invention from the overall content of this specification.

According to one aspect of the present invention, a steel for an ultra-thick pressure vessel contains, in wt%, 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 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%, the remainder includes iron and unavoidable impurities, the elongation in the thickness direction (ZRA) is 35% or more, the tensile strength is 450 to 620 MPa, and the maximum surface crack depth may be 2 μm or less.

The thickness of the above steel can be 133 to 233 mm.

The above steel may contain ferrite having an average grain size of 50㎛ or less.

According to another aspect of the present invention, a method for manufacturing a steel for an extremely thick pressure vessel comprises the following components in weight %: 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 to 0.20%, molybdenum (Mo): 0.01 to 0.15%, copper (Cu): 0.01 to 0.50%, nickel (Ni): 0.05 to 0.50%, A first heating step of heating a 700 mm thick slab containing calcium (Ca): 0.0005 to 0.0040%, the remainder iron and unavoidable impurities, to a temperature range of 1100 to 1300°C; a step of performing a first forging process on the first heated slab with a cumulative reduction amount of 3 to 15% and a strain rate of 1/s to 4/s to provide a first intermediate material; a second heating step of heating the first intermediate material to a temperature range of 1000 to 1200°C; a step of performing a second forging process on the second heated first intermediate material with a cumulative reduction amount of 3 to 30% and a strain rate of 1/s to 4/s to provide a second intermediate material; a third heating step of heating the second intermediate material to a temperature range of 1000 to 1200°C; The step of hot-rolling the secondary intermediate material that has been heated three times above in a temperature range of 900 to 1100°C to provide a hot-rolled material; and the step of heating the hot-rolled material in a temperature range of 820 to 900°C, maintaining it for 10 to 40 minutes, and then cooling it to room temperature to perform a normalizing heat treatment may be included.

The thickness of the above first intermediate material may be 450 to 550 mm, the thickness of the above second intermediate material may be 300 to 450 mm, and the thickness of the above hot-rolled material may be 133 to 233 mm.

The central porosity of the above secondary intermediate material may be 0.1 mm ³ /g or less.

The maximum surface crack depth of the above secondary intermediate material may be 5 ㎛ or less, and the maximum surface crack depth of the above hot-rolled material may be 2 ㎛ or less.

According to a preferred aspect of the present invention, it is possible to provide a steel material for an extremely thick pressure vessel having excellent surface quality and lamella tearing resistance and a method for manufacturing the same.

The present invention relates to a steel for an extremely thick pressure vessel having excellent surface quality and lamella tearing resistance and a method for manufacturing the same. Hereinafter, preferred embodiments of the present invention will be described. The embodiments of the present invention may be modified in various forms, and the scope of the present invention should not be construed as being limited to the embodiments described below. The embodiments are provided to those skilled in the art to which the present invention pertains in order to further detail the present invention.

Hereinafter, a steel material for an extremely thick pressure vessel having excellent surface quality and lamella tearing resistance according to one aspect of the present invention will be described in more detail.

According to one aspect of the present invention, a steel for an ultra-thick pressure vessel contains, in wt%, 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 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%, the remainder includes iron and unavoidable impurities, the elongation in the thickness direction (ZRA) is 35% or more, the tensile strength is 450 to 620 MPa, and the maximum surface crack depth may be 2 μm or less.

Hereinafter, the alloy composition of the present invention will be described in more detail. Hereinafter, unless specifically stated otherwise, % and ppm related to the contents of the alloy composition are based on weight.

According to one aspect of the present invention, a steel for an ultra-thick pressure vessel contains, in wt%, 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 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): 0.0005~0.0040%, the remainder may contain iron and unavoidable impurities.

Carbon (C): 0.10~0.25%

Since C is the most important element for securing basic strength, it needs to be contained in steel within an appropriate range, and in order to obtain this addition effect, 0.10% or more of C can be added. Preferably, 0.12% or more of C can be added. On the other hand, if the C content exceeds a certain level, the fraction of pearlite increases during normalizing heat treatment, which may excessively exceed the strength and hardness of the base material, causing surface cracks to occur during forging processing, and deteriorating the anti-lamellar tearing characteristics in the final product. Therefore, the upper limit of the C content may be 0.25%, and a more preferable upper limit of the C content may be 0.20%.

Silicon (Si): 0.05~0.50%

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 for manufacturing clean steel. Therefore, Si can be added at 0.05% or more, and more preferably at 0.20% or more. On the other hand, when Si is added in a large amount, it can generate an MA phase and excessively increase the strength of the ferrite matrix, which can cause deterioration in the surface quality of ultra-thick products. Therefore, it is preferable that the upper limit of the content is limited to 0.5%. The upper limit of the preferable Si content can be 0.40%.

Manganese (Mn): 1.0~2.0%

Mn is a useful element that improves strength by solid solution strengthening and enhances hardenability so that a low-temperature transformation phase is generated. Therefore, in order to secure a tensile strength of 450 MPa or more, it is preferable that Mn be added at 1.0% or more. A more preferable Mn content may be 1.1% or more. On the other hand, as the Mn content increases, Mn forms MnS, a non-metallic inclusion that is elongated together with S, which lowers the toughness, and acts as a factor that lowers the elongation during thickness-wise tension, which may be a factor that rapidly lowers the internal lamellar tearing quality. Therefore, it is preferable to manage the Mn content to 2.0% or less, and a more preferable Mn content may be 1.5% or less.

Aluminum (Al): 0.005~0.1%

Al, along with Si, is one of the powerful deoxidizers in the steelmaking process, and to obtain this effect, it is preferable to add 0.005% or more. The lower limit of the more preferable Al content may be 0.01%. On the other hand, if the Al content is excessive, the fraction of Al2O3 among the oxidizing inclusions generated as a result of deoxidation increases excessively, and the size thereof becomes coarse, and it becomes difficult to remove the inclusions during refining, which may be a factor that lowers the anti-lamellar tearing characteristics. Therefore, it is preferable to manage the Al content to 0.1% or less, and the more preferable Al content may be 0.07% or less.

Phosphorus (P): 0.010% or less, Sulfur (S): 0.0015% or less

Since P and S are elements that induce brittleness at grain boundaries or form coarse inclusions, it is desirable to limit P to 0.010% or less and S to 0.0015% or less to improve brittle crack propagation resistance.

Niobium (Nb): 0.001~0.03%

Nb is an element that improves the strength of the base material by precipitating in the form of NbC or NbCN. In addition, Nb, which is dissolved during reheating at a high temperature, is very finely precipitated in the form of NbC during rolling, which suppresses recrystallization of austenite and has the effect of refining the structure. Therefore, it is preferable that Nb is added in an amount of 0.001% or more, and a more preferable Nb content may be 0.005% or more. On the other hand, if Nb is added excessively, undissolved Nb is generated in the form of TiNb(C,N), which becomes a factor that inhibits the lamellar tearing characteristics, and therefore it is preferable that the upper limit of the Nb content is limited to 0.03%. A more preferable Nb content may be 0.02% or less.

Vanadium (V): 0.001~0.03%

Since V is almost entirely reused during reheating, the strengthening effect by precipitation or solid solution during subsequent rolling is minimal, but it has the effect of improving the strength by precipitating very fine carbonitrides during subsequent heat treatment processes such as PWHT. In order to sufficiently obtain this effect, it is necessary to add V at 0.001% or more. The lower limit of the more preferable V content may be 0.01%. On the other hand, if the content is excessive, it may excessively increase the strength and hardness of the base material and the weld, which may act as a factor of surface cracks during pressure vessel processing, and the manufacturing cost may increase rapidly, which is not commercially advantageous. Therefore, the V content may be 0.03% or less, and the more preferable V content may be 0.02% or less.

Titanium (Ti): 0.001~0.03%

Since Ti is a component that significantly improves low-temperature toughness by precipitating as TiN during reheating and suppressing the growth of grains in the parent material and the weld heat-affected zone, it is preferable to add 0.001% or more to obtain this addition effect. On the other hand, if Ti is added excessively, the low-temperature toughness may decrease due to clogging of the casting nozzle or central crystallization, and since it combines with N to form coarse TiN precipitates in the center of the thickness, the elongation of the product may decrease, so the lamellar tearing resistance of the final product may deteriorate. Therefore, the Ti content may be 0.03% or less. The preferable Ti content may be 0.025% or less, and the more preferable Ti content may be 0.018% or less.

Chromium (Cr): 0.01~0.20%

Cr increases the hardenability and forms a low-temperature transformation structure, thereby increasing the yield and tensile strength, and prevents the decrease in strength by slowing down the decomposition rate of cementite during tempering after rapid cooling or heat treatment after welding. For this effect, 0.01% or more of Cr can be added. On the other hand, when the Cr content is excessive, the size and fraction of Cr-rich coarse carbides such as M ₂₃ C ₆ increase, thereby lowering the impact toughness of the product, while the solubility of Nb in the product and the fraction of fine precipitates such as NbC decrease, so that the decrease in the strength of the product may become a problem. Therefore, the upper limit of the Cr content can be limited to 0.20%. The desirable upper limit of the Cr content can be 0.15%.

Molybdenum (Mo): 0.01~0.15%

Mo is an element that increases grain boundary strength and has a large effect of solid solution strengthening in ferrite, and thus effectively contributes to increasing the strength and ductility of the product. In addition, Mo has the effect of preventing the deterioration of toughness due to grain boundary segregation of impurities such as P. For this effect, it is desirable to add 0.10% or more of Mo. However, since Mo is an expensive element, excessive addition can significantly increase the manufacturing cost, and therefore the upper limit of the Mo content can be limited to 0.15%.

Copper (Cu): 0.01~0.50%

Cu can greatly improve the strength of the matrix by solid solution strengthening in ferrite, and also has the effect of suppressing corrosion in a wet hydrogen sulfide atmosphere, and is therefore an advantageous element in the present invention. For this effect, 0.01% or more of Cu may be included. A more preferable Cu content may be 0.03% or more. However, when the content of Cu is excessive, there is a problem that star cracks are likely to be induced on the surface of the steel sheet, and the manufacturing cost increases significantly as it is an expensive element, and therefore the upper limit of the Cu content in the present invention may be limited to 0.50%. A preferable upper limit of the Cu content may be 0.30%.

Nickel (Ni): 0.05~0.50%

Ni is an important element for improving impact toughness and hardenability by increasing stacking faults at low temperatures and facilitating cross-slip of dislocations, thereby improving strength. For this effect, 0.05% or more of Ni may be added. The preferred Ni content may be 0.10% or more. On the other hand, if Ni is added excessively, the manufacturing cost may also increase due to the high cost, so the upper limit of the Ni content may be limited to 0.50%. The preferred upper limit of the Ni content may be 0.30%.

Calcium (Ca): 0.0005~0.0040%,

When Ca is added after deoxidation by Al, it combines with S that forms MnS inclusions to suppress the formation of MnS, and at the same time, it forms spherical CaS to suppress the occurrence of cracks due to hydrogen-induced cracking. In order to sufficiently form S contained as an impurity into CaS, it is preferable to add Ca in an amount of 0.0005% or more. However, if the amount added is excessive, the Ca remaining after forming CaS combines with O to generate coarse oxidizing inclusions, which are elongated and destroyed during rolling, resulting in a problem of deterioration of the lamellar tearing resistance. Therefore, the upper limit of the Ca content can be limited to 0.0040%.

The steel for an ultra-thick pressure vessel according to one aspect of the present invention may contain residual Fe and other unavoidable impurities in addition to the above-mentioned components. However, since unintended impurities may inevitably be mixed in from raw materials or the surrounding environment during a normal manufacturing process, they cannot be completely excluded. Since these impurities are known to anyone with ordinary knowledge in the art, not all of the contents are specifically mentioned in this specification. In addition, addition of effective components other than the above-mentioned composition is not excluded.

According to one aspect of the present invention, a steel for an extremely thick pressure vessel may have a steel thickness of 133 to 233 mm, an elongation in the thickness direction (ZRA) of 35% or more, and a tensile strength of 450 to 620 MPa.

According to one aspect of the present invention, a steel for an ultra-thick pressure vessel has a microstructure of ferrite average and pearlite composite structures, and the average grain size of the ferrite may be 50 ㎛ or less.

According to one aspect of the present invention, a steel material for a pressure vessel having a central void ratio of 0.1 mm ³ /g or less and a maximum surface crack depth of 2 ㎛ or less may be used.

Hereinafter, a method for manufacturing a steel for an ultra-thick pressure vessel according to one aspect of the present invention will be described in more detail.

According to another aspect of the present invention, a method for manufacturing a steel for an extremely thick pressure vessel comprises: a first heating step of heating a slab having a thickness of 700 mm and having a predetermined alloy composition to a temperature range of 1100 to 1300°C; a step of performing a first forging process on the first heated slab at a cumulative reduction amount of 3 to 15% and a strain rate of 1/s to 4/s to provide a first intermediate material; a second heating step of heating the first intermediate material to a temperature range of 1000 to 1200°C; a step of performing a second forging process on the second heated first intermediate material to a cumulative reduction amount of 3 to 30% and a strain rate of 1/s to 4/s to provide a second intermediate material; and a third heating step of heating the second intermediate material to a temperature range of 1000 to 1200°C. The step of hot-rolling the secondary intermediate material that has been heated three times above in a temperature range of 900 to 1100°C to provide a hot-rolled material; and the step of heating the hot-rolled material in a temperature range of 820 to 900°C, maintaining it for 10 to 40 minutes, and then cooling it to room temperature to perform a normalizing heat treatment may be included.

Slavic primary heating

The slab having a predetermined alloy composition can be heated in a temperature range of 1100 to 1300°C. Since the alloy composition of the slab corresponds to the alloy composition of the aforementioned extremely thick steel, the description of the alloy composition of the slab is replaced with the description of the alloy composition of the aforementioned extremely thick steel. The thickness of the slab can be 650 to 750 mm, and preferably 700 mm.

In order to re-dissolve the composite carbonitrides of Ti or Nb or the coarse-grained TiNb(C,N) particles formed during casting, to homogenize the structure by heating and maintaining the austenite above the recrystallization temperature before the first forging, and to ensure a sufficiently high forging completion temperature to minimize surface cracks that may occur during the forging process, the first heating of the slab can be performed at a temperature range of 1100℃ or higher.

On the other hand, if the slab is heated at an excessively high temperature, problems may occur due to oxidation scale at high temperatures, and the manufacturing cost may increase excessively due to increased costs due to heating and maintenance. Therefore, the upper limit of the slab heating temperature may be limited to 1300℃.

1st forging

The first heated slab can be subjected to first forging at a cumulative reduction of 3 to 15% and a strain rate of 1/s to 4/s to provide a first intermediate product.

The first forging is a step in which a slab heated in a temperature range of 1100 to 1300℃ is forged to a thickness of 450 to 550 mm, and processed into the final desired width of the secondary intermediate material. Since high-strain, low-speed forging is essential to sufficiently compress the void, the first forging can be performed under the conditions of a cumulative reduction of 3 to 15% and a forging speed of 1/s to 4/s.

When the accumulated reduction amount of the first forging is less than 3%, the remaining voids in the slab are not sufficiently compressed, resulting in residual voids, which may deteriorate the lamella tearing resistance of the product. The preferred accumulated reduction amount of the first forging may be 5% or more, and the more preferred accumulated reduction amount of the first forging may be 7% or more. However, when the accumulated reduction amount below the non-recrystallization temperature at which the dislocation density is recovered or is not offset by recrystallization exceeds 15%, the uniform elongation of the surface is extremely reduced due to the work hardening of the superimposed dislocations, and surface cracks may occur during the forging process. The preferred accumulated reduction amount of the first forging may be 13% or less, and the more preferred accumulated reduction amount of the first forging may be 11% or less.

Secondary heating and secondary forging

The primary intermediate material is re-heated at a temperature range of 1000 to 1200℃, and the second-heated primary intermediate material is re-forged at a cumulative reduction amount of 3 to 30% and a strain rate of 1/s to 4/s to provide a secondary intermediate material with a thickness of 300 to 450 mm. The maximum surface crack depth of the secondary intermediate material can be 5㎛ or less.

This is a step of processing the primary intermediate material into the desired thickness and length of the secondary intermediate material by heating and forging it in a temperature range of 1000 to 1200℃. As in the primary forging, high-strain, low-speed forging is essential in the secondary forging in order to secure the central porosity of the secondary intermediate material to 0.1 mm ³ /g or less. Therefore, the secondary forging can be performed by applying a cumulative reduction amount of 3 to 30% and a strain rate of 1/s to 4/s.

If the accumulated reduction amount of the second forging is less than 3%, the micropores remaining after the first forging cannot be completely compressed, and when deformation is applied to the end points of the oval-shaped compressed pores, the physical properties may be worse than when the pores are in a circular shape due to the notch effect. Therefore, it is necessary to sufficiently squeeze the pores with a deformation of 3% or more during the second forging. However, if the accumulated reduction amount exceeds 30%, surface cracks may occur due to surface work hardening.

The strain rate of the second forging can be 1/s to 4/s, similar to the first forging. At a rate less than 1/s, surface cracks may occur due to the temperature drop in the final forging, and a high strain rate exceeding 4/s in the non-recrystallized region can also cause a decrease in elongation and surface cracks.

3rd heating and hot rolling

The secondary intermediate material after forging can be reheated to a temperature range of 1000 to 1200℃.

The composite carbonitrides of Ti or Nb or coarse TiNb(C,N) crystallites formed during casting are reused, and the austenite is heated to a temperature higher than the recrystallization temperature before hot rolling and maintained to homogenize the structure. In order to ensure a sufficiently high rolling end temperature and minimize crushing of inclusions during the rolling process, a third heating can be performed in a temperature range of 1000℃ or higher.

On the other hand, if the secondary intermediate material is heated at an excessively high temperature, problems may occur due to oxidation scale at high temperatures, and the manufacturing cost may increase excessively due to the increased cost due to heating and maintenance. Therefore, the upper limit of the 3rd heating temperature may be limited to 1200℃.

The third heated secondary intermediate material can be hot rolled at a temperature range of 900 to 1100℃ to provide a hot-rolled material having a thickness of 133 to 233 mm. The maximum surface crack depth of the hot-rolled material can be 2㎛ or less.

When the final hot rolling temperature is less than 900℃, the deformation resistance value increases excessively as the temperature decreases, so it is difficult to sufficiently refine the austenite grains in the center of the thickness direction of the product, and accordingly, the lamellar tearing resistance of the final product may be deteriorated. On the other hand, when the hot rolling temperature exceeds 1100℃, the austenite grains become excessively coarse, so there is a concern that the strength and impact toughness may be deteriorated. Therefore, the hot rolling temperature is preferably 900 to 1100℃.

Normalizing heat treatment

The hot rolled steel that has been hot rolled can be normalized by heating it to a temperature range of 820 to 900℃, maintaining it for 10 to 40 minutes, and then cooling it to room temperature.

During normalizing heat treatment, if the temperature is lower than 820℃ or the holding time is shorter than 10 minutes, carbides generated during cooling after rolling or impurity elements segregated at grain boundaries do not re-dissolve smoothly, which may significantly deteriorate the elongation in the thickness direction (ZRA) and low-temperature toughness of the steel after heat treatment. On the other hand, during normalizing heat treatment, if the temperature exceeds 900℃ or the holding time exceeds 40 minutes, the quality of anti-lamellar tearing may deteriorate due to the coarsening of austenite and the coarsening of precipitated phases such as Nb(C,N) and V(C,N).

Hereinafter, the present invention will be described more specifically through examples. However, it should be noted that the examples described below are only intended to illustrate and further specify the present invention, and are not intended to limit the scope of the rights of the present invention.

(Example)

A 700 mm thick cast steel having the alloy composition of Table 1 was manufactured. First forging, second forging, hot rolling, and normalizing heat treatment were performed according to the process conditions of Table 2. At this time, a first heating temperature of 1200℃, a second heating temperature of 1100℃, and a third heating temperature of 1050℃ were commonly applied, and a normalizing time of 30 minutes was commonly applied. A condition of 550 mm was applied for the thickness of the first intermediate material, and a condition of 400 mm was applied for the thickness of the second intermediate material.

division Alloy composition (wt%)
※ P, S, Ca and O are in weight ppm units River species C Si Mn Al P S Nb V Ti Cr Mo Cu Ni Ca 1 0.15 0.35 1.32 0.03 91 10 0.017 0.007 0.012 0.05 0.09 0.05 0.18 19 2 0.16 0.3 1.44 0.025 80 12 0.011 0.008 0.007 0.07 0.08 0.07 0.22 20 3 0.17 0.33 1.37 0.03 77 11 0.018 0.01 0.005 0.1 0.1 0.05 0.23 15 4 0.16 0.36 1.23 0.027 81 9 0.017 0.021 0.01 0.08 0.08 0.1 0.24 17 5 0.16 0.4 1.19 0.03 87 8 0.015 0.015 0.011 0.09 0.06 0.08 0.19 16 6 0.18 0.35 1.22 0.03 79 10 0.019 0.009 0.009 0.12 0.08 0.09 0.25 18 7 0.03 0.31 1.21 0.03 78 10 0.013 0.009 0.005 0.2 0.08 0.07 0.19 19 8 0.16 0.35 2.35 0.03 78 9 0.016 0.014 0.004 0.25 0.09 0.06 0.18 18 9 0.17 0.3 1.21 0.029 82 35 0.012 0.015 0.005 0.23 0.1 0.1 0.19 20 10 0.16 0.25 1.32 0.31 86 11 0.015 0.016 0.012 0.15 0.07 0.06 0.23 2

Psalter
No. River species 1st forging Second forging Rolling
end
temperature
(℃) Normalizing
temperature
(℃) Pressure
(%) Transformation rate
(/s) Pressure
(%) Transformation rate
(/s) A 1 8.5 1.5 8.4 2.3 983 856 B 2 10.3 1.7 11.3 3.5 1005 881 C 3 12.5 3.1 10.9 2.1 982 880 D 4 10.9 1.3 11.3 2.3 991 891 E 5 11.3 1.5 12.4 3.1 909 850 F 6 12.4 1.2 13.5 2 915 837 G 7 13.5 1.4 11.2 1.8 957 840 H 8 12.1 2.9 10.5 1.2 962 851 I 9 13.5 3.4 11.7 1.9 941 871 J 10 9.8 2.3 12.4 2.5 918 883 K 1 1.3 2.8 10.9 3.1 924 859 L 2 10.7 0.1 8.4 2.7 929 855 M 3 13.2 3.5 54 1.9 935 843 N 4 11.4 2 10.5 5.2 954 841 O 5 10.9 1.9 11.4 1.9 793 835 P 6 10.3 1.2 10.5 1.4 955 1009

For each specimen, the void ratio, microstructure fraction, average ferrite grain size, ZRA, tensile strength, and step-by-step surface crack depth were quantitatively measured after forging, and the results are shown in Table 3.

Psalter
No River species After forging
Porosity
(mm ³ /g) Pera
Eat
(area%) Perla
Eat
(area%) ferrite
Average particle size
(um) Surface crack depth
(mm) ZRA
(%) seal
robbery
(MPa) minor
after Rolling
after A 1 0.002 83 17 37 0 0 65 512 B 2 0.001 87 13 35 0 0 60 513 C 3 0.015 85 15 40 0 0 49 535 D 4 0.007 84 16 36 0 0 53 507 E 5 0.013 87 13 38 0 0 62 499 F 6 0.012 86 14 37 0 0 63 513 G 7 0.023 85 15 35 0 0 64 401 H 8 0.015 86 14 33 8.4 4.1 12 673 I 9 0.003 83 17 38 0 0 14 522 J 10 0.004 82 18 41 0 0 8 532 K 1 0.19 84 16 38 0 0 7 515 L 2 0.015 83 17 35 7.4 3 60 509 M 3 0 85 15 33 14.5 7.2 62 513 N 4 0.013 83 17 37 8.3 3.9 64 489 O 5 0.011 84 16 38 6.2 3 66 541 P 6 0.013 84 16 79 0 0 27 472

It can be confirmed that specimens that satisfy both the alloy composition and the process conditions limited by the present invention satisfy the properties targeted by the present invention, whereas specimens that do not satisfy any one of the alloy composition and the process conditions limited by the present invention fail to secure the properties at the level targeted by the present invention.

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

Claims (7)

In weight %, 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 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): 0.0005~0.0040%, including the remainder of Fe and inevitable impurities,
The elongation in the thickness direction (ZRA) is 35% or more,
The tensile strength is 450~620MPa,
Steel for extremely thick pressure vessels with a maximum surface crack depth of 2㎛ or less.
In the first paragraph,
The above steel is for use in extremely thick pressure vessels, with a thickness of 133 to 233 mm.
In the first paragraph,
The above steel is a steel for extremely thick pressure vessels, containing ferrite having an average grain size of 50㎛ or less.
In weight %, 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 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): The first heating step is to heat a 700 mm thick slab containing 0.0005 to 0.0040% iron and unavoidable impurities at a temperature range of 1100 to 1300°C;
A step of providing a first intermediate material by first forging the first heated slab at a cumulative reduction amount of 3 to 15% and a strain rate of 1/s to 4/s;
A second heating step of heating the above primary intermediate material to a temperature range of 1000 to 1200°C;
A step of providing a second intermediate material by performing a second forging process on the second heated first intermediate material at an accumulated pressure of 3 to 30% and a strain rate of 1/s to 4/s;
A third heating step of heating the above secondary intermediate material to a temperature range of 1000 to 1200°C;
A step of hot-rolling the above-mentioned thirdly heated secondary intermediate material in a temperature range of 900 to 1100℃ to provide a hot-rolled material; and
A method for manufacturing steel for pressure vessels of extremely thick masses, comprising the step of heating the hot-rolled material to a temperature range of 820 to 900°C, maintaining it for 10 to 40 minutes, and then cooling it to room temperature to perform normalizing heat treatment.
In paragraph 4,
The thickness of the above primary intermediate material is 450 to 550 mm,
The thickness of the above secondary intermediate material is 300 to 450 mm,
A method for manufacturing a steel for an extremely thick pressure vessel, wherein the thickness of the hot-rolled steel is 133 to 233 mm.
In paragraph 4,
A method for manufacturing a steel material for an extremely thick pressure vessel, wherein the porosity of the central portion of the secondary intermediate material is 0.1 mm ³ /g or less.
In paragraph 4,
The maximum surface crack depth of the above secondary intermediate material is 5㎛ or less,
A method for manufacturing a steel for an extremely thick pressure vessel, wherein the maximum surface crack depth of the hot-rolled material is 2㎛ or less.