JP2025500911A - Extra-thick steel material for flanges with excellent strength and low-temperature impact toughness and its manufacturing method - Google Patents

Abstract

An object of the present invention is to provide an extra-thick steel material for flanges having excellent strength and low-temperature impact toughness, and a manufacturing method thereof.
[Solution] The extra-thick steel material for flanges of the present invention, which has excellent strength and low-temperature impact toughness, contains, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0. .3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, with the remainder being Fe and other unavoidable impurities, the prior austenite grain size is 35 μm or less, and it has a fine structure containing 90 area % or more of one or more of bainite and martensite, and residual ferrite or pearlite.
[Selection diagram] None

Description

The present invention relates to steel materials that can be used in wind power towers and systems, and to a method for manufacturing the same, and more specifically to an extremely thick steel material for flanges that has excellent strength and low-temperature impact toughness, and a method for manufacturing the same.

Wind turbines are attracting attention as an environmentally friendly means of generating electricity, and include components such as tower flanges, bearings, and main shafts. Among these, tower flanges are joint parts required for connecting towers, and usually 5 to 7 flanges are used for one tower. They are installed at sea or in extreme regions, so high durability is required. In particular, in response to the demand for large-capacity energy production and high efficiency, the scale of wind towers is also increasing, and accordingly, the steel materials used are continuously required to be high-strength, high-toughness, and thick. As the thickness of the material increases, the total deformation amount decreases, and the microstructure becomes larger, and there is a tendency for the material to deteriorate due to defects in the material such as inclusions and segregation. Therefore, in order to improve the external soundness resistance of steel materials, there is a trend to reduce the concentration of impurities such as non-metallic inclusions and segregations, and to control cracks and voids on the surface and inside the material to the utmost.

In particular, in the case of extremely thick materials with a thickness of over 200 mm, the amount of deformation in the center of the material is not large, so if unsolidified shrinkage holes that occur during continuous casting or casting are not sufficiently compressed in the forging process, they will remain in the form of residual voids in the center of the flange.

When the structure is subjected to axial stress through its thickness, these residual voids can act as the initiation point for cracks, eventually causing damage to the entire equipment in the form of lamellar tearing. Therefore, a process is required to fully crimp the central void so that no residual voids remain before piercing (hole forging) and ring forging (product forming), which involve little deformation.

Related Patent Document 1 is a technology for applying strong reduction in the rough rolling process of thick plates. Specifically, it utilizes a technology for determining the limit reduction rate for each thickness at which plate bite occurs from the strong reduction rate for each pass set to be close to the design allowable value (load and torque) of the rolling mill, a technology for distributing the reduction rate by adjusting the index of the thickness ratio for each pass to ensure the target thickness of the rough rolling mill, and a technology for correcting the reduction rate based on the limit reduction rate for each thickness so that plate bite does not occur, and provides a manufacturing method that can apply an average reduction rate of about 27.5% in the final three passes of rough rolling based on 80 mmt as a reference. However, in the case of the above rolling method, the average reduction rate of the entire product thickness is measured, and in the case of an extremely thick material with a maximum thickness of 200 mmt or more, it is technically difficult to apply a high deformation to the center where residual voids exist.

Another method for manufacturing extremely thick products is to utilize a forging machine, which has a higher effective deformation per pass than a rolling machine. In Patent Document 2, a slab containing, by mass%, C: 0.08-0.20%, Si: 0.40% or less, Mn: 0.5-5.0%, P: 0.010% or less, S: 0.0050% or less, Cr: 3.0% or less, Ni: 0.1-5.0%, Al: 0.010-0.080%, N: 0.0070% or less, and O: 0.0025% or less, which satisfies the relationships in Formulas 1 and 2, with the remainder being Fe and unavoidable impurities, is hot forged with a cumulative reduction of 25% or more. It clearly states that by heating the material to between the Ac3 point and 1200°C, hot rolling with a cumulative reduction of 40% or more, quenching it at a temperature of the Ar3 point or higher to a low temperature of 350°C or lower or below the Ar3 point, and passing it through a tempering heat treatment process at a temperature of 450 to 700°C, it is possible to produce a thick, high-toughness, high-strength material with a plate thickness of 100 mmt or more, a yield strength of 620 MPa or more, and an absorbed energy of 70 J or more when evaluated for low-temperature impact toughness at -40°C.

However, in the above manufacturing method, if the cumulative reduction is too high, localized deformation concentration can cause surface defects. In particular, if surface or subsurface defects exist in the cast slab before forging, the defects can propagate during the forging process, causing further deterioration of the surface quality in the product state after rolling. In addition, if the forging reduction per pass is insufficient, it is difficult to fully close the voids remaining in the center even if the cumulative reduction is high, and the rolling process is also not suitable for controlling the voids and structure in the center of extremely thick materials, since the effective deformation in the center is smaller than the surface deformation.

Meanwhile, Patent Document 3 discloses that a thick, high-strength steel plate of 100 mmt or more and with a yield strength of 620 MPa or more can be manufactured through a process in which a material provided with a specified alloy composition is heated to 1200 to 1350°C, hot forged with a cumulative reduction of 25% or more, heated to between the Ac3 point and 1200°C, hot rolled with a cumulative reduction of 40% or more, reheated to between the Ac3 point and 1050°C, quenched at a temperature of the Ac3 point or more to a temperature of 350°C or less or Ar3 point or less, whichever is lower, and tempered at a temperature of 450°C to 700°C.

However, in the case of the above-mentioned ultra-high strength steel plate, the carbon equivalent (Ceq) and hardenability index (DI) are high, making it vulnerable to surface cracks during casting, and in the case of flange steel manufactured by normalizing heat treatment, the process conditions cannot be easily applied. In addition, when the carbon equivalent (Ceq) and hardenability index (DI) are high, cracks are likely to occur on the surface of the cast piece due to the generation of surface hard structure during the secondary cooling process of steelmaking, and the cracks may propagate during the forging process, degrading the surface quality of the final product.

Therefore, a forging method has been proposed to close the void in the center and improve the internal integrity of the final product, but no practical method has been presented to ensure both the appropriate material and excellent surface quality of the flange steel.

Korean Patent Publication No. 10-2012-0075246 Korean Patent Publication No. 10-2017-0095307 Korean Patent Publication No. 10-2017-0095307

Therefore, the object of the present invention is to provide an extra-thick steel material for flanges that has excellent strength and low-temperature impact toughness, and a manufacturing method thereof.

The object of the present invention is not limited to the above. A person skilled in the art would have no difficulty in understanding the further object of the present invention from the overall content of this specification.

The extra-thick steel material for flanges of the present invention is
The alloy contains, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, and the remainder being Fe and other unavoidable impurities,
The prior austenite grain size is 35 μm or less, and the microstructure contains at least one of bainite and martensite at 90 area % or more and residual ferrite or pearlite, and the low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more,
Having 10 or more deformation-induced NbC precipitates of 5 to 50 nm per ¹ μm2 and 5 or less coarse precipitates of 100 nm or more,
The steel material is characterized in that the porosity of the center portion of the steel material, which is the region from the surface to 3/8t to 5/8t in the thickness direction, is 0.05 mm ³ /g or less.

The above steel further contains Zr: 0.001-0.15%.

The above steel material has a thickness of 200 to 500 mm.

The steel also has a tensile strength of 590 to 820 MPa, a yield strength of 440 MPa or more, and an absorbed energy value of 50 J or more in a -50°C Charpy impact test.

The maximum surface crack depth of the above steel material is 0.1 mm or less (including 0).

In addition, the method for producing an extra-thick steel material for flanges of the present invention includes the steps of:
preparing a slab containing, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, with the remainder being Fe and other unavoidable impurities; and heating the slab to a temperature range of 1100-1300°C.
performing a first upsetting of the heated slab at a forging ratio of 1.3 to 2.4, and then bloom forging at a forging ratio of 1.5 to 2.0;
reheating the bloom forged material to a temperature range of 1100 to 1300°C;
Secondarily upsetting the reheated bloom material with a forging ratio of 1.3 to 2.3, and then round forging with a forging ratio of 1.65 to 2.25;
The round forged material is subjected to a third upsetting at a temperature equal to or lower than a recrystallization temperature defined by the following Relation 1, with a forging ratio of 2.0 to 2.8 so that the cumulative reduction is 10% or more;
After the tertiary upset material is drilled, it is reheated to a temperature range of 1100 to 1300 ° C., and then ring forged with a forging ratio of 1.0 to 1.6;
The ring-forged material is heated to a temperature range of 820 to 930°C as a temperature measurement standard at the center of the material, and held for 5 to 600 minutes, and then air-cooled to room temperature, and then heated to 550 to 700°C and held there.

[Relationship 1]
T _nr (℃)=887+464×C+890×Ti+363×Al-357×Si+(6445×Nb-644×Nb ^1/2 )+(732×V-230×V ^1/2 )
Here, C, Ti, Al, Si, Nb, and V are in weight percent.

The slab is manufactured using one of the following processes: continuous casting, semi-continuous casting, and ingot casting.

If the size of the forged surface punched out during the above primary upsetting is initially 700mm x 1800mm, it will be 1000-1200mm x 1800-2000mm.

In the case of the above bloom forging, if the size of the forged surface at the end of forging is initially 1000-1200mm x 1800-2000mm, it will be 1450-1850mm x 2100-2500mm.

When the above secondary upsetting and round forging are completed, the product size is 1450-1850Φ x 1300-1700mm.

When the above tertiary upsetting is completed, the product size is 2300-2800Φ x 400-800mm.

The maximum thickness of a flange made from the above steel material can be 200-500 mm, the inner diameter is 4000-7000 mm, and the outer diameter is 5000-8000 mm.

By optimizing the forging process, the present invention, configured as described above, can crimp the void in the center of the steel material, improving the internal integrity of the final product, and can effectively provide extra-thick steel material that can be used for flanges, with excellent strength as well as low-temperature impact toughness.

The present invention relates to an extremely thick steel material for flanges having excellent strength and low-temperature impact toughness, and a manufacturing method for the product. A preferred embodiment of the present invention will be described below. The embodiment of the present invention can be modified in various forms, and the scope of the present invention should not be construed as being limited to the embodiment described below. The embodiment is provided to allow a person having ordinary knowledge in the technical field to which the present invention pertains to understand the present invention in more detail.

The following is a more detailed explanation of the extra-thick steel material for flanges of the present invention, which has excellent strength and low-temperature impact toughness.

The extra-thick steel material for flanges of the present invention, which is excellent in strength and low-temperature impact toughness, contains, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: The steel material has a prior austenite grain size of 35 μm or less, and a microstructure containing at least one of bainite and martensite at 90 area % or more and residual ferrite or pearlite. The low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more, and has 10 or ^more deformation-induced NbC precipitates of 5 to 50 nm and 5 or less coarse precipitates of 100 nm or more per μm2. The porosity of the center of the steel material, which is a region of 3/8t to 5/8t from the surface in the thickness direction, is 0.05 ^mm3 /g or less.

The alloy composition of the present invention will be described in more detail below. Unless otherwise specified, the percentages and ppm described in the alloy composition are based on weight.

・Carbon (C): 0.05-0.20%
Carbon (C) is the most important element for ensuring basic strength, and therefore must be contained in the steel within an appropriate range. In order to obtain this additive effect, 0.05% Carbon (C) of 0.10% or more can be added. Preferably, carbon (C) of 0.10% or more can be added. On the other hand, when the carbon (C) content exceeds a certain level, the QT heat treatment Sometimes the hardening capacity increases too much and can exceed the strength and hardness of the base material, which can lead to surface cracks during forging and reduced low-temperature impact toughness in the final product. Therefore, in the present invention, the carbon (C) content can be limited to 0.20%, and more preferably, the upper limit of the carbon (C) content can be 0.18%.

Silicon (Si): 0.05 to 0.50%
Silicon (Si) is a substitutional element that improves the strength of steel by solid solution strengthening and has a strong deoxidizing effect, and is therefore an essential element for the production of clean steel. Therefore, silicon (Si) may be added in an amount of 0.05% or more, and more preferably 0.20% or more. On the other hand, when a large amount of silicon (Si) is added, it may generate a MA (Martensite-Austenite) phase, which may excessively increase the strength of the matrix and cause deterioration of the surface quality of extremely thick products, so the upper limit of the content may be limited to 0.50%. A more preferable upper limit of the silicon (Si) content may be 0.40%.

Manganese (Mn): 1.0 to 2.0%
Manganese (Mn) is a useful element that improves strength by solid solution strengthening and improves hardening ability so that a low-temperature transformation phase is generated. Therefore, in order to ensure a tensile strength of 590 MPa or more, it is preferable to add 1.0% or more of manganese (Mn). More preferably, the manganese (Mn) content is 1.1% or more. Meanwhile, manganese (Mn) forms MnS, which is a non-metallic inclusion stretched together with sulfur (S), and reduces toughness and can act as an impact initiation point, which can be a factor in rapidly reducing the low-temperature impact toughness of the product. Therefore, it is preferable to control the manganese (Mn) content to 2.0% or less, and more preferably, the manganese (Mn) content can be 1.5% or less.

Aluminum (Al): 0.005 to 0.1%
Aluminum (Al) is one of the powerful deoxidizing agents in the steelmaking process along with silicon (Si), and in order to obtain this effect, it is preferable to add 0.005% or more. A more preferable lower limit of the aluminum (Al) content may be 0.01%. On the other hand, if the aluminum (Al ₎ content is excessive, the fraction of _Al2O3 in the oxidized inclusions generated as a result of deoxidation increases excessively, and the size of the inclusions becomes coarse, which makes it difficult to remove the inclusions during refining, which may cause a decrease in low-temperature impact toughness. Therefore, it is preferable to control the aluminum (Al) content to 0.1% or less. A more preferable aluminum (Al) content may be 0.07% or less.

Phosphorus (P): 0.010% or less (including 0%), Sulfur (S): 0.0015% or less (including 0%)
Phosphorus (P) and sulfur (S) are elements that induce brittleness at grain boundaries or form coarse inclusions that induce brittleness, so in order to improve brittle crack propagation resistance, it is preferable to limit phosphorus (P) to 0.010% or less and sulfur (S) to 0.0015% or less.

Niobium (Nb): 0.005 to 0.07%
Niobium (Nb) is an element that precipitates in the form of NbC or NbCN to improve the strength of the base material. In addition, niobium (Nb) dissolved during high-temperature reheating precipitates very finely in the form of deformation induction below the recrystallization temperature during forging, suppressing the growth of austenite, and thus has the effect of refining the structure. Therefore, niobium (Nb) is preferably added at 0.005% or more, and a more preferable niobium (Nb) content may be 0.01% or more. On the other hand, when niobium (Nb) is added in excess, undissolved niobium (Nb) is generated in the form of TiNb (C, N), which is a factor that inhibits low-temperature impact toughness, so the upper limit of the niobium (Nb) content is preferably limited to 0.07%. A more preferable niobium (Nb) content may be 0.065% or less.

Vanadium (V): 0.001 to 0.3%
Vanadium (V) is almost completely dissolved again during reheating, so the strengthening effect due to precipitation or dissolution during subsequent rolling is insufficient. However, in the case of very thick forged materials, the air cooling rate is very slow, so it precipitates as very fine carbonitrides during the cooling process or tempering heat treatment process, which has the effect of improving strength. In order to fully obtain such an effect, it is necessary to add 0.001% or more of vanadium (V). A more preferable lower limit of the vanadium (V) content may be 0.01%. On the other hand, if the content is excessive, the surface hardness of the slab increases excessively due to its high hardening ability, which may act as a cause of surface cracks during flange processing, and the manufacturing cost increases rapidly, which is not commercially useful. Therefore, the vanadium (V) content may be limited to 0.3% or less. A more preferable vanadium (V) content may be 0.25% or less.

Titanium (Ti): 0.001 to 0.05%
Titanium (Ti) is a component that precipitates as TiN during reheating, suppresses the growth of prior austenite grains at high temperatures, and greatly improves low-temperature toughness. In order to obtain such an effect, it is preferable to add 0.001% or more of titanium (Ti). On the other hand, if titanium (Ti) is added excessively, there is a possibility that the low-temperature toughness will decrease due to clogging of the continuous casting nozzle or crystallization in the center. In addition, titanium (Ti) combines with nitrogen (N) to form coarse TiN precipitates in the thickness center, which reduces the elongation of the product, and therefore may reduce the uniform elongation in the forging process and cause surface cracks. Therefore, the titanium (Ti) content may be 0.05% or less. A preferred titanium (Ti) content may be 0.03% or less, and a more preferred titanium (Ti) content may be 0.018% or less.

Chromium (Cr): 0.01 to 0.30%
Chromium (Cr) is an ingredient that has the effect of preventing a decrease in strength by retarding the spheroidization rate of cementite, and improves hardenability during the cooling process. For this effect, 0.01% or more of chromium (Cr) can be added. On the other hand, if the chromium (Cr) content is excessive, the size and fraction of Cr-rich coarse carbides such as _{M23C6 etc.} increase, thereby decreasing the impact toughness of the product, and the solid solubility of niobium (Nb) in the product and the fraction of fine precipitates such as NbC decrease, which may cause a problem of a decrease in strength of the product. Therefore, in the present invention, the upper limit of the chromium (Cr) content can be limited to 0.30%. The preferred upper limit of the chromium (Cr) content can be 0.25%.

Molybdenum (Mo): 0.01 to 0.12%
Molybdenum (Mo) is an element that increases grain boundary strength, increases hardening ability, and dissolves in precipitates to improve strength, and is an element that effectively contributes to increasing the strength and ductility of products. Molybdenum (Mo) also has the effect of preventing a decrease in toughness due to grain boundary segregation of impurity elements such as phosphorus (P). For this effect, 0.01% or more of molybdenum (Mo) can be added. However, since molybdenum (Mo) is an expensive element and excessive addition may significantly increase manufacturing costs, the upper limit of the molybdenum (Mo) content can be limited to 0.12%.

・Copper (Cu): 0.01-0.60%
Copper (Cu) is an element that can greatly improve the strength of the base phase by solid solution strengthening in ferrite. For this effect, copper (Cu) can be contained in an amount of 0.01% or more. More preferably, the copper (Cu) content is 0.03% or more. However, if the copper (Cu) content is excessive, the possibility of inducing star cracks on the surface of the steel sheet increases. Since copper (Cu) is an expensive element, there is a possibility that the production cost will increase significantly. Therefore, in the present invention, the upper limit of the copper (Cu) content can be limited to 0.60%. A preferred upper limit for the copper (Cu) content may be 0.5%.

Nickel (Ni): 0.05 to 4.00%
Nickel (Ni) is an element that effectively contributes to improving strength by increasing stacking faults at low temperatures, facilitating cross slip of dislocations to improve impact toughness, improving hardenability, and improving the degree of solid solution strengthening. To achieve this effect, 0.05% or more of nickel (Ni) may be added. A preferred nickel (Ni) content may be 0.10% or more. Meanwhile, when nickel (Ni) is added excessively, the manufacturing cost may increase due to the high cost, so the upper limit of the nickel (Ni) content may be limited to 4.00%. A preferred upper limit of the nickel (Ni) content may be 3.5%.

Calcium (Ca): 0.0005 to 0.0040%
Adding calcium (Ca) after deoxidation with aluminum (Al) has the effect of suppressing the generation of MnS by combining with sulfur (S) forming MnS inclusions, and suppressing the generation of cracks due to hydrogen-induced cracking by forming spherical CaS. In order to sufficiently form sulfur (S) contained as an impurity into CaS, it is preferable to add 0.0005% or more of calcium (Ca). However, if the amount added is excessive, the calcium (Ca) remaining after forming CaS combines with oxygen (O) to generate coarse oxidized inclusions, which are stretched and broken during forging, and may be a factor in reducing the elongation and low-temperature impact toughness properties. Therefore, the upper limit of the calcium (Ca) content can be limited to 0.0040%.

Zirconium (Zr): 0.001 to 0.15%
The steel material of the present invention may optionally contain Zr in the range of 0.001 to 0.15. Zirconium (Zr) is a strong carbide forming element and may exist in the form of ZrC, and may improve the strength of the matrix phase in the form of precipitation strengthening like VC and NbC. For this effect, 0.001% or more of zirconium (Zr) may be added. However, if added in excess, it may increase the manufacturing cost due to its high cost, so the upper limit of the Zr content may be 0.15%.

The extra-thick flange steel material and products thereof, which have excellent strength and low-temperature impact toughness, of the present invention may contain the remaining Fe and other inevitable impurities in addition to the above-mentioned components. However, since unintended impurities may be unavoidably mixed in from the raw materials or the surrounding environment during normal manufacturing processes, these cannot be completely excluded. Since these impurities are known to anyone with ordinary knowledge in this technical field, this specification does not specifically mention all of their contents. Furthermore, the addition of further effective components in addition to the above-mentioned components is not completely excluded.

On the other hand, the extra-thick steel material of the present invention has a prior austenite grain size of 35 μm or less, and has a fine structure containing at least one of bainite and martensite at 90 area % or more, with the remainder being ferrite or pearlite.

In the present invention, the prior austenite grain size in the microstructure of the final product is 35 μm or less. If the prior austenite grain size exceeds 35 μm, the length of the crack path during impact fracture becomes shorter, the DBTT (Ductile Brittle Transition Temperature) increases, and low-temperature impact toughness deteriorates. Therefore, it is preferable that the prior austenite grain size is 35 μm or less.

The steel material of the present invention has a microstructure containing at least one of bainite and martensite at 90% by area or more, with the remainder being ferrite or pearlite. If the phase fraction of low-temperature transformation phases such as bainite and martensite is less than 95% by area, the strength of the base phase decreases, and the material cannot meet the tensile strength of 590 to 820 MPa and yield strength of 440 MPa or more proposed in the present invention.

In addition, in the present invention, the above-mentioned low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more. When the matrix base structure is bainite or martensite, cracks during impact testing propagate along packet boundaries based on high-angle grain boundaries, so if the packet size is large, the DBTT may increase and the impact toughness may deteriorate. Therefore, in order to ensure an absorbed energy value of 50 J or more in a -50°C Charpy impact test required in the present invention, it is appropriate that the packet size is 15 μm or less.

In addition, the steel material of the present invention can have 10 or more deformation-induced NbC precipitates of 5 to 50 nm per 1 μm ² and 5 or less coarse precipitates of 100 nm or more in its base structure. If there are less than 10 deformation-induced NbC precipitates of 5 to 50 nm, the precipitation strengthening effect is weakened, and if the number of coarse precipitates of 100 nm or more exceeds 5, the pinning effect and the precipitation strengthening effect are similarly lost, so that it is not easy to ensure the tensile strength of 590 to 820 MPa and the yield strength of 440 MPa or more required in the present invention.

Furthermore, the steel material of the present invention has a porosity of 0.05 mm ³ /g or less in the center of the steel material, which is a region from the surface to 3/8t to 5/8t in the thickness direction.

In addition, the extra-thick steel material of the present invention can have a thickness of 200 to 500 mm.

In addition, the extra-thick steel material of the present invention can have a tensile strength of 590 to 820 MPa, a yield strength of 440 MPa or more, and an absorbed energy value of 50 J or more in a Charpy impact test at -50°C.

The maximum surface crack depth of the steel material can be 0.1 mm or less (including 0).

Next, we will explain in detail another aspect of the present invention, a method for manufacturing extra-thick steel material for flanges that has excellent strength and low-temperature impact toughness.

The manufacturing method of the extra-thick steel material of the present invention includes the steps of preparing a slab having the compositional components as described above, heating it to a temperature range of 1100 to 1300°C, performing a first upset of the heated slab with a forging ratio of 1.3 to 2.4, and then performing bloom forging at a ratio of 1.5 to 2.0, reheating the bloom forged material to a temperature range of 1100 to 1300°C, performing a second upset of the reheated bloom material with a forging ratio of 1.3 to 2.3, and then performing round forging at a forging ratio of 1.65 to 2.25, and then performing round forging of the round forged material. The method includes a step of tertiary upsetting at a temperature equal to or lower than the recrystallization temperature defined by the following relational expression 1 with a forging ratio of 2.0 to 2.8 so that the cumulative reduction is 10% or more, a step of reheating the tertiary upset material to a temperature range of 1100 to 1300°C after drilling holes, and then ring forging with a forging ratio of 1.0 to 1.6, and a step of heating the ring forged material to a temperature range of 820 to 930°C as a temperature measurement standard for the center part and holding it for 5 to 600 minutes, air cooling it to room temperature, and then heating it to 550 to 700°C and holding it there.

Slab Heating First, in the present invention, a slab having the above-mentioned composition is prepared, and then heated to a temperature range of 1100 to 1300°C.

The slab needs to be heated above a certain temperature range in order to redissolve titanium (Ti) and niobium (Nb) carbonitrides or TiNb (C, N) coarse crystals formed during casting. In addition, it is preferable to heat the slab above a certain temperature range in order to homogenize the structure by heating the slab to the recrystallization temperature or higher before the primary upset forging and hold it there, ensure a sufficiently high forging end temperature, and minimize surface cracks that may occur during the forging process. Therefore, it is preferable to heat the slab in the present invention at a temperature range of 1100°C or higher.

On the other hand, if the slab heating temperature is excessively high, high-temperature oxide scale may be generated excessively, and high-temperature heating and holding may result in excessive increases in manufacturing costs. Therefore, it is preferable that the primary heating of the slab in the present invention is performed in the range of 1300°C or less.

Primary Upsetting and Bloom Forging Next, in the present invention, the heated slab is subjected to a primary upsetting with a forging ratio of 1.3 to 2.4, and then to bloom forging with a forging ratio of 1.5 to 2.0.

Upsetting is a method of performing strong plastic deformation perpendicular to the longitudinal axis, and the forging ratio during the first upsetting is preferably 1.3 to 2.4, and more preferably 1.5 to 2.0. Here, the forging ratio refers to the ratio of the cross-sectional area that changes due to forging. During such a first upsetting, if the size of the forged surface to be punched out is initially 700 mm x 1800 mm, it can be 1000 to 1200 mm x 1800 to 2000 mm.

If the forging ratio in the first upset is less than 1.3, it is difficult to sufficiently compress the porosity remaining in the center of the slab. Therefore, it is difficult to control the porosity required in the final product of the present invention to an appropriate level of 0.05 ^mm3 /g or less, and it is not easy to ensure low-temperature impact toughness in the center. On the other hand, if the forging ratio in the first upset is more than 2.4, buckling occurs during the forging process, making it difficult to perform appropriate shape control required for surface quality and flange products. Therefore, the forging ratio in the first upset is preferably 1.3 to 2.4.

In the present invention, the above-mentioned primary upset material is bloom forged with a forging ratio of 1.5 to 2.0.

Bloom forging is a method of further compressing the primary upset material into a bloom shape, and enlarging the area by processing both the top and bottom surfaces in a fixed direction in width or length. In the case of bloom forging, if the size of the forged surface at the end of forging is initially 1000-1200mm x 1800-2000mm, it can be 1450-1850mm x 2100-2500mm. In the case of bloom forging, the forging ratio is preferably 1.5 to 2.0. This is because, like upset forging, if the forging ratio is less than 1.5, it is difficult to ensure the appropriate void quality required in the present invention, and if it exceeds 2.0, surface cracks may occur.

The forging can proceed either lengthwise or widthwise. However, in the case of the lengthwise direction, the cast structure is more dense, so the elongation rate of the surface layer structure is high and the workability is excellent. Therefore, from the viewpoint of surface cracks, bloom forging in the lengthwise direction may be more appropriate than that in the widthwise direction.

Reheating In the present invention, the bloom forged material is reheated to a temperature range of 1100 to 1300°C.

When the bloom forging is completed, the bloom surface temperature is 950°C or less, and if processing is continued, there is a possibility that surface cracks or material breakage may occur. Therefore, after bloom forging, the material can be heated again to a temperature range of 1100 to 1300°C. As mentioned above, it is preferable to heat to 1100°C or more for reasons such as re-dissolving the crystallized particles, homogenizing the structure, and preventing surface cracks, but it is preferable to control the temperature to 1300°C or less due to problems such as excessive scaling and coarsening of crystal grains.

Secondary Upsetting-Round Forging Next, in the present invention, the reheated bloom material is subjected to a secondary upsetting with a forging ratio of 1.3 to 2.3, and then round forging with a forging ratio of 1.65 to 2.25.

That is, after the bloom has been heated, a secondary upset is performed with a forging ratio of 1.3 to 2.3, and then round forging is performed with a forging ratio of 1.65 to 2.25 to process the bloom into a circular flange edge. When the secondary upsetting and round forging are completed, the product size can be 1450 to 1850Φ x 1300 to 1700mm.

If the forging ratio during the secondary upset and round forging is less than the level required by the present invention, it is difficult to control the porosity of the center of the final product to 0.05 ^mm3 /g or less, and it is not easy to ensure low-temperature impact toughness at the center of the steel material. On the other hand, if the forging ratio exceeds the standard of the present invention, it is not possible to process into the desired flange product shape due to problems such as buckling, surface cracks, and shape defects.

Tertiary upset (generation of deformation-induced precipitates)
Subsequently, in the present invention, the round forged material is tertiary upset at a temperature equal to or lower than the recrystallization temperature defined by the following Relation 2 with a forging ratio of 2.0 to 2.8 so that the cumulative reduction is 10% or more.

The cylindrically processed material can be processed to the appropriate flange thickness by a tertiary upsetting before hole drilling (piercing). When the tertiary upsetting is completed, the product size can be 2300-2800Φ x 400-800mm.

The forging ratio of the tertiary upset can be 2.0 to 2.8, and if the forging ratio is insufficient or excessive, problems such as the above-mentioned inability to control residual voids and surface cracks/shape control may occur.

What is important in this tertiary upsetting process is the cumulative reduction at a temperature below the recrystallization temperature (Rst) of the steel, and forging and rolling are performed so that the cumulative reduction is 10% or more. In this case, the recrystallization temperature can be calculated using the following relational expression 2.

[Relationship 2]
T _nr (℃)=887+464×C+890×Ti+363×Al-357×Si+(6445×Nb-644×Nb ^1/2 )+(732×V-230×V ^1/2 )
Here, C, Ti, Al, Si, Nb, and V are in weight percent.

If the cumulative reduction is less than 10% at a temperature below the recrystallization temperature, it is not easy to generate deformation-induced ultrafine NbC or NbCN precipitates of 5 to 50 nm, and the number of precipitates per ¹ μm2 may be less than 10, or the number of coarse precipitates with a size of 100 nm or more may exceed 5. If the amount of fine precipitates is reduced or the size is increased, the precipitation strengthening effect is small, and the pinning effect is reduced during the subsequent reheating for quenching, making it difficult to ensure that the average prior austenite grain size in the center of the product is 35 μm or less. Therefore, it is preferable to control the cumulative reduction at a temperature below the recrystallization temperature to 10% or more, more preferably 15% or more, and most preferably 20% or more.

Hole Machining and Ring Forging Next, in the present invention, the tertiary upset material is subjected to hole machining, and then reheated to a temperature range of 1100 to 1300° C., and subsequently ring forged with a forging ratio of 1.0 to 1.6.

After the above tertiary upsetting is completed, a hole can be drilled in the center of the material using a 500 to 1000Φ punch.

The material with the holes machined is reheated again to the aforementioned temperature range of 1100 to 1300°C, and then processed into the final flange ring shape. The maximum thickness of the flange made of the steel material can be 200 to 500 mm, the inner diameter can be 4000 to 7000 mm, and the outer diameter can be 5000 to 8000 mm. Since ring forging is a process in which control of the final shape and dimensions is more important than gap pressing, severe plastic processing is not applied. Therefore, the forging ratio can be 1.0 to 1.6, and more preferably 1.2 to 1.4.

On the other hand, in the present invention, the deformation speed in all of the forging processes presented above can be 1/s to 4/s. At a deformation speed of less than 1/s, the finish forging temperature drops, and there is a possibility that surface cracks may occur. In contrast, when a high deformation speed of more than 4/s is applied in the unrecrystallized region, there is a possibility that surface cracks will be induced due to a decrease in elongation caused by excessive local work hardening.

Quenching and tempering heat treatment Finally, in the present invention, the ring-forged material is heated to a temperature range of 820 to 930°C as a temperature measurement standard at its center (t/2) and held for 5 to 600 minutes, then air-cooled to room temperature, and then heated to 550 to 700°C and held there, thereby performing quenching and tempering heat treatment.

During quenching heat treatment, if the heating temperature is less than 820°C or the holding time is less than 5 minutes, the carbides generated during cooling after forging and the impurity elements segregated at the grain boundaries may not be smoothly redissolved, and the low-temperature toughness of the steel after heat treatment may be significantly reduced. On the other hand, if the heating temperature exceeds 930°C or the holding time exceeds 600 minutes, the prior austenite grain size may exceed 35 μm required in the present invention, or the precipitate phases such as Nb(C,N) and V(C,N) may become coarse, resulting in a deterioration in strength and low-temperature impact toughness.

On the other hand, the cooling rate can be 0.5°C/s to 30°C/s based on the center (t/2) of the product. If the cooling rate is less than 0.5°C/s, it is difficult to ensure adequate strength because the fraction of bainite or martensite, which is the low-temperature transformation structure required in the present invention, cannot be ensured to be 90% or more, and if it exceeds 30°C/s, the strength increases excessively and low-temperature impact toughness may deteriorate. Therefore, it is preferable that the cooling rate during quenching is 0.5°C/s to 30°C/s.

Meanwhile, tempering can be performed at a temperature range of 550-700°C for 5-600 minutes. If the tempering temperature is below 550°C, carbon does not diffuse properly after quenching, resulting in excessive strength, which may result in deterioration of low-temperature impact toughness at -50°C. On the other hand, if the tempering temperature exceeds 700°C, fresh-martensite formed during the air-cooling process by heating in the two-phase region may also result in deterioration of low-temperature impact toughness. Therefore, it is preferable that the tempering temperature is 550-700°C.

If the tempering holding time is less than 5 minutes, the dislocation density after quenching does not decrease appropriately, and carbon does not diffuse sufficiently, resulting in excessively high strength and reduced low-temperature impact toughness, for the same reason that the temperature is low. Also, if the holding time exceeds 600 minutes, carbon becomes excessively spheroidized and coarsened, resulting in reduced impact toughness. Therefore, it is preferable that the appropriate holding time for tempering heat treatment is 5 to 600 minutes.

The present invention will be described in detail below with reference to examples. However, it should be noted that the following examples are intended to illustrate and explain the present invention in more detail, and are not intended to limit the scope of the invention.

＊表１において、成分元素の含量単位は重量％であるが、Ｐ、Ｓ及びＣａの単位はｐｐｍである。そして、残りの成分はＦｅと不可避不純物である。

*In Table 1, the content units of the component elements are in weight percent, but the units of P, S, and Ca are in ppm, and the remaining components are Fe and unavoidable impurities.

A 700 mm thick slab was manufactured with the alloy composition shown in Table 1 above. This slab was used to manufacture a final 320 mmt flange through the process conditions shown in Tables 2-3 below, including slab preparation, forging (reheating and first upset, bloom forging, reheating, second upset-round forging, third upset, reheating and ring forging), and quenching and tempering. At this time, after bloom forging was completed, the reheating temperature for the second upset was 1230°C ± 10°C, and the forging ratio for the round forging after the second upset was 2.0, which was also applied. All other process conditions that meet the scope of the present invention were applied to the processes other than those shown in Tables 2-3 below.

The physical properties of each test piece produced above were then measured and are shown in Table 4 below.

Here, the prior austenite grain size and the fraction of low-temperature transformation phases (bainite and/or martensite) were measured using an automatic image analyzer by taking a test specimen from the central structure test specimen after QT heat treatment, and the packet size of bainite was automatically analyzed by setting the boundary condition to 15° through EBSD (Electron Back Scattered Diffraction). Meanwhile, in this embodiment, the residual structure excluding the low-temperature transformation phase in both the invention example and the comparative example is ferrite and/or pearlite.

The yield/tensile strength was evaluated by a room temperature tensile test according to ASTM A370, and a 0.2% offset was applied to the yield strength. The impact toughness of each test piece was calculated by using the average absorbed energy value measured three times at the corresponding temperature through the Charpy V-Notch Test.

The number of deformation-induced NbC precipitates of 5 to 50 nm observed in the cross section of the steel material was measured using a TEM. The NbC precipitates were confirmed by the NbC diffraction pattern and EDX mapping, and the number of NbC precipitates located within 1 μm ² was counted.

The porosity of the center of the product was measured by measuring the density (g/mm ³ ) and taking the reciprocal (mm ³ /g).

After visually observing the surface of each test piece, the test piece was ground at the point where a surface crack had formed, and the grinding length until the crack disappeared was measured as the length of the surface crack.

＊表３において、ａ＊は鋼材の断面で観察される５～５０ｎｍの変形誘起ＮｂＣ析出物の１μｍ^２当たりの個数であり、ｂ＊は１００ｎｍ以上の粗大析出物の個数を示す。

*In Table 3, a * indicates the number of deformation-induced NbC precipitates of 5 to 50 nm observed per 1 ^μm2 in the cross section of the steel material, and b * indicates the number of coarse precipitates of 100 nm or more.

As can be seen from Tables 1 to 3 above, in the case of Examples 1 to 5, which satisfy the alloy composition and manufacturing conditions proposed by the present invention, not only do they have excellent strength and excellent low-temperature impact toughness at -50°C, but they also ensure good surface quality in the flange product state.

In contrast, Comparative Examples 1 to 12 are cases in which the alloy composition proposed by the present invention is satisfied but the manufacturing conditions are not satisfied, and since the properties proposed by the present invention in the flange product state, such as the prior austenite grain size, the fraction and packet size of the low-temperature transformation phase, and the porosity, are not satisfied, it can be seen that the strength and low-temperature impact toughness values are at a low level. Furthermore, even if the material quality is good, if the forging ratio conditions are not satisfied at each stage of forging, poor surface quality properties in the product state can be confirmed by the occurrence of surface cracks or through cracks.

On the other hand, Comparative Examples 13 to 16 meet the manufacturing conditions proposed by the present invention, but do not meet the alloy composition, and are of low quality, with excessive strength (insufficient impact toughness) or insufficient strength.

As described above, the detailed description of the present invention has been given of the preferred embodiment of the present invention, but it goes without saying that a person having ordinary skill in the art to which the present invention pertains can make various modifications without departing from the scope of the present invention. Therefore, the scope of the present invention should not be limited to the described embodiment, but should be determined by the claims below as well as equivalents thereto.

Claims (12)

The alloy contains, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, and the remainder being Fe and other unavoidable impurities,
The prior austenite grain size is 35 μm or less, and the microstructure contains at least one of bainite and martensite at 90 area % or more, and residual ferrite or pearlite. The low-temperature transformation phase has a packet size of 15 μm or less based on a high-angle grain boundary of 15° or more.
Having 10 or more deformation-induced NbC precipitates of 5 to 50 nm per ¹ μm2 and 5 or less coarse precipitates of 100 nm or more,
The extra-thick steel material for flanges is characterized in that the porosity of the center of the steel material, which is a region from the surface to 3/8t to 5/8t in the thickness direction, is 0.05 mm ³ /g or less. The extra-thick steel material for flanges according to claim 1, characterized in that the steel material further contains, by weight percent, Zr: 0.001 to 0.15%. The extra-thick steel material for flanges according to claim 1, characterized in that the steel material has a thickness of 200 to 500 mm. The extra-thick steel material for flanges according to claim 1, characterized in that the steel material has a tensile strength of 590 to 820 MPa, a yield strength of 440 MPa or more, and an absorbed energy value of 50 J or more in a -50°C Charpy impact test. The extra-thick steel material for flanges according to claim 1, characterized in that the maximum surface crack depth of the steel material is 0.1 mm or less (including 0). preparing a slab containing, by weight, C: 0.05-0.2%, Si: 0.05-0.5%, Mn: 1.0-2.0%, Al: 0.005-0.1%, P: 0.01% or less, S: 0.015% or less, Nb: 0.005-0.07%, V: 0.001-0.3%, Ti: 0.001-0.05%, Cr: 0.01-0.3%, Mo: 0.01-0.12%, Cu: 0.01-0.6%, Ni: 0.05-4.0%, Ca: 0.0005-0.004%, with the remainder being Fe and other unavoidable impurities; and heating the slab to a temperature range of 1100-1300°C.
performing a first upsetting of the heated slab at a forging ratio of 1.3 to 2.4, and then bloom forging at a forging ratio of 1.5 to 2.0;
reheating the bloom forged material to a temperature range of 1100 to 1300°C;
Secondarily upsetting the reheated bloom material at a forging ratio of 1.3 to 2.3, and then round forging at a forging ratio of 1.65 to 2.25;
The round forged material is subjected to a third upsetting at a temperature equal to or lower than the recrystallization temperature defined by the following Relation 1, with a forging ratio of 2.0 to 2.8 so that the cumulative reduction is 10% or more;
After the tertiary upset material is drilled, it is reheated to a temperature range of 1100 to 1300 ° C., and then ring forged with a forging ratio of 1.0 to 1.6;
The method for manufacturing an extra-thick steel material for flanges includes a step of heating the ring-forged material to a temperature range of 820 to 930°C as a temperature measurement standard at the center of the material, holding the temperature range for 5 to 600 minutes, air-cooling the material to room temperature, and then heating the material to 550 to 700°C and holding the temperature range for 5 to 600 minutes.
[Relationship 1]
T _nr (℃)=887+464×C+890×Ti+363×Al-357×Si+(6445×Nb-644×Nb ^1/2 )+(732×V-230×V ^1/2 )
(Here, C, Ti, Al, Si, Nb, and V are in weight percent.) The method for manufacturing the extra-thick steel material for flanges according to claim 6, characterized in that the slab is manufactured using one of a continuous casting process, a semi-continuous casting process, and an ingot casting process. The method for manufacturing the extra-thick steel material for flanges described in claim 6 is characterized in that the size of the forged surface punched out during the first upsetting is 1000-1200mm x 1800-2000mm when the initial size is 700mm x 1800mm. The method for manufacturing the extra-thick steel material for flanges described in claim 6 is characterized in that, in the case of bloom forging, the size of the forged surface at the completion of forging is 1450-1850mm x 2100-2500mm when it is initially 1000-1200mm x 1800-2000mm. The manufacturing method for extra-thick steel material for flanges described in claim 6, characterized in that when the secondary upsetting and round forging are completed, the product size is 1450-1850Φ x 1300-1700mm. The manufacturing method for extra-thick steel material for flanges described in claim 6, characterized in that when the tertiary upsetting is completed, the product size is 2300-2800Φ x 400-800mm. The manufacturing method of the extra-thick steel material for flanges described in claim 6, characterized in that the maximum thickness of the flange made of the steel material can be 200 to 500 mm, the inner diameter is 4000 to 7000 mm, and the outer diameter is 5000 to 8000 mm.