Austenitic steel material and manufacturing method therefor

Abstract

The present invention relates to an austenitic steel material applicable to various fields such as the energy industry and shipbuilding, and a manufacturing method therefor and, more specifically, to an austenitic steel material and a manufacturing method therefor wherein the austenitic steel material can be used in a structure requiring high strength and excellent hydrogen embrittlement resistance such as a hull requiring high strength and excellent non-magnetic properties, a steel pipe or facilities transporting crude oil having high hydrogen sulfide content, and the like.

Description

Austenitic steel and its manufacturing method

[0001] The present invention relates to austenitic steels applicable to various fields, including the energy industry, shipbuilding, and marine structures, and a method for manufacturing the same. More specifically, the present invention relates to austenitic steels applicable to ship hulls requiring excellent non-magnetic properties, as well as steel pipes and equipment for transporting crude oil with a high hydrogen sulfide content, and a method for manufacturing the same.

[0002] As crude oil extraction environments become increasingly harsh, containing large amounts of hydrogen sulfide and other gases, the steel pipes and structures used to transport crude oil require materials with superior hydrogen embrittlement resistance. Therefore, transporting crude oil requires materials with superior hydrogen embrittlement resistance and sufficient toughness and strength to withstand the external pressures encountered during transportation.

[0003] Conventional high-strength carbon steels have been limited by the presence of internal structures vulnerable to hydrogen embrittlement, known as "hard spots," making them difficult to produce and limiting their use. To overcome these limitations, various manufacturing methods and methods for detecting these hard spots have been proposed. However, the inherent instability of hard spots for increased strength has made it difficult to produce steels that combine superior hydrogen embrittlement resistance with high strength.

[0004] Furthermore, in the case of structures that use a large amount of steel materials, such as military vessels such as submarines, the location of the opponent can be tracked by detecting changes in the magnetic field generated by the interaction between the hull and the Earth's magnetic field. If such structures are manufactured with austenitic steel that can maintain non-magnetism even after deformation, the probability of detection can be greatly reduced. In the case of conventional carbon steel, it goes through a process called demagnetization to periodically remove magnetism, but this time reduces the operating time of the vessel and also causes economic loss. Therefore, the demand for steel with excellent non-magnetic properties is increasing.

[0005] One aspect of the present invention is to provide an austenitic steel having high strength characteristics, excellent impact toughness and hydrogen embrittlement resistance, and excellent investment rate even after deformation, and a method for manufacturing the same.

[0006] The objectives of the present invention are not limited to the above-described scope. Those skilled in the art will have no difficulty understanding the additional objectives of the present invention from the overall content of this specification.

[0007] According to one aspect of the present invention, a steel material comprises, in wt%, C: 0.050% to 1.70%, Mn: 15.0 to 40.0%, Cr: 3.00% or less, V: 1.00 to 3.00%, N: 1.000% or less (excluding 0%), Mo: 3.50% or less, and Nb: 1.00% or less, with the remainder being iron (Fe) and inevitable impurities, and a microstructure in which a main structure is austenite, an area fraction of grain boundary carbides formed at grain boundaries of the austenite is 5.0 area% or less, and at least one or more kinds of fine precipitates of VC and VCN having a diameter of 50 nm or less are 100 / mm 2 The number of units per unit area may be included.

[0008] The above-described steel may additionally contain at least one of Ti: 1.00% or less, Al: 5.00% or less, and Si: 5.00% or less.

[0009] The above-described steel can satisfy the following equation 1.

[0010] [Equation 1] 23.6[C]+[Mn] ≥ 28.000, 33.5[C]-[Mn] ≤ 23.00

[0011] (In the above formula 1, [C] and [Mn] represent the weight percentages of C and Mn contained in the steel, respectively.)

[0012] The above-described grain boundary carbide may include at least one of Cr, Mo and Nb carbides.

[0013] The area fraction of the austenite described above may be 95 area% or more.

[0014] The room temperature yield strength of the above-mentioned steel may be 550 MPa or more, the Charpy impact energy value at -84°C may be 27 J or more, and the permeability after 20% cold plastic deformation at room temperature may be 1.200 or less.

[0015] In addition, the steel material described above may have an investment ratio of 1.100 or less when the strain is at least 2% during cold plastic deformation at room temperature, and a crack length ratio (CLR) in a hydrogen-induced cracking (HIC) test defined by the following equation 2 may be 10% or less.

[0016] [Formula 2]

[0017] CLR(Crack Length Ratio, %) = ∑(a / W) * 100

[0018] (In the above equation 2, a represents the length of a single crack (㎛), and w represents the width of the specimen (㎛).)

[0019] In addition, the steel material described above may have an area fraction of grain boundary carbides in the weld heat affected zone of 5.0 area% or less after submerged welding with a heat input of 3.0 kJ / mm.

[0020] A method for manufacturing a steel according to another aspect of the present invention may include the steps of: heating a slab comprising, in wt%, C: 0.050% to 1.70%, Mn: 15.0% to 40.0%, Cr: 3.00% or less, V: 1.00% to 3.00%, N: 1.000% or less (excluding 0%), Mo: 3.50% or less, and Nb: 1.00% or less, with the remainder being iron (Fe) and unavoidable impurities; finishing hot-rolling the slab to obtain a hot-rolled steel sheet; cooling the hot-rolled steel sheet to room temperature after solution treatment; and aging the hot-rolled steel sheet.

[0021] The above-described slab may additionally include at least one of Ti: 1.00% or less, Al: 5.00% or less, and Si: 5.00% or less, and may satisfy the following equation 1.

[0022] [Equation 1] 23.6[C]+[Mn] ≥ 28.000, 33.5[C]-[Mn] ≤ 23.00

[0023] (In the above formula 1, [C] and [Mn] represent the weight percentages of C and Mn contained in the steel, respectively.)

[0024] Finally, the above-described heating can be performed at 1000°C or more and 1300°C or less, the above-described finishing hot rolling can be performed at 700°C or more and 1050°C or less, the above-described solution treatment can be performed at 900°C or more and 1200°C or less for 30 minutes or more and 2 hours or less, and the above-described aging treatment can be performed at 500°C or more and 850°C or less for 30 minutes or more and 5 hours or less.

[0025] The present invention can provide a steel material that can be used in a structure requiring high strength characteristics, excellent impact toughness, hydrogen embrittlement resistance, and non-magnetic characteristics.

[0026] The various advantageous and beneficial advantages and effects of the present invention are not limited to the above-described contents, and will be more easily understood in the course of explaining specific embodiments of the present invention.

[0027] Figure 1 is a graph showing the range of carbon and manganese according to one aspect of the present invention.

[0028] Figure 2 is a transmission electron microscope photograph of Example 1, a steel material according to one aspect of the present invention.

[0029] Hereinafter, preferred embodiments of the present invention will be described. However, the embodiments of the present invention may be modified into various other forms, and the scope of the present invention is not limited to the embodiments described below.

[0030] In this specification, the term "including" is used to indicate that other components may be included rather than excluding other components unless specifically stated to the contrary.

[0031] Additionally, unless otherwise specifically provided in the specification of the present invention, the % unit means weight%.

[0032] Traditionally, austenitic steels with high manganese content have had high chromium content added to improve corrosion resistance and strength, and molybdenum and niobium, which play a similar role to the aforementioned chromium, have also tended to be added to improve the strength of the steel.

[0033] However, in this case, there was a problem in that it was difficult to secure excellent impact toughness of the steel because chromium, molybdenum, or niobium combined with carbon to form carbides at the austenite grain boundaries.

[0034] Accordingly, the inventors of the present invention found that when the content of chromium, molybdenum or niobium added to the steel is reduced and vanadium is added at a high concentration, the above-mentioned problem can be solved, and at the same time, the strength of the steel can be improved by the vanadium-derived fine precipitate.

[0035] From this point of view, a steel according to one embodiment of the present invention may include, in wt%, C: 0.050% to 1.70%, Mn: 15.0% to 40.0%, Cr: 3.00% or less, V: 1.00% to 3.00%, N: 1.000% or less (excluding 0%), Mo: 3.50% or less, and Nb: 1.00% or less, with the remainder being iron (Fe) and unavoidable impurities.

[0036] Below, each component is described in detail.

[0037] C: 0.050%~1.70%

[0038] C is an element that stabilizes austenite and increases strength, and in particular, it plays a role in lowering Ms and Md, which are transformation points from austenite to epsilon or alpha martensite during the cooling process or working. Therefore, if C is added insufficiently, the stability of austenite is insufficient, making it impossible to obtain stable austenite at extremely low temperatures. In addition, external stress can easily cause working-induced transformation into epsilon or alpha martensite, thereby reducing toughness and also reducing the strength of the steel. Therefore, the lower limit of C in the present invention may be 0.050%. In another embodiment, the present invention may include C in an amount of 0.070% or more, and in another embodiment, it may include C in an amount of 0.100% or more. On the other hand, if the content of C is excessive, the toughness may deteriorate rapidly due to carbide precipitation, and the workability may deteriorate due to an excessive increase in strength. Therefore, the present invention may include C in an amount of 1.70% or less. In another embodiment, the present invention may set the upper limit of C to 1.50% or less, and in another embodiment, to 1.30%.

[0039] Mn: 15.0~40.0%

[0040] Manganese (Mn) is an element that plays an important role in stabilizing austenite. In one embodiment of the present steel, it is preferable that 15.0% or more of Mn be included to stabilize austenite. If the Mn content is lower than this, epsilon martensite, a metastable phase, is formed, which easily transforms into alpha martensite through working-induced transformation at ultra-low temperatures, so that high toughness may not be secured. There is a method of increasing the C content to stabilize austenite in order to suppress the formation of epsilon martensite, but in this case, a large amount of carbides may be precipitated, which may rapidly deteriorate the physical properties, more specifically, the toughness. Therefore, the Mn content is preferably 15.0% or more. In another embodiment, the Mn content may be 18.0% or more, and in yet another embodiment, the Mn content may be 20.0% or more.

[0041] Excessive manganese content can not only slow down the corrosion rate of steel but is also undesirable from an economic perspective. Therefore, the manganese content included in the steel according to one aspect of the present invention may be 40.0% or less. In another embodiment, the manganese content may be 35.0% or less, and in another example, 30.0% or less.

[0042] Cr: 3.00% or less

[0043] Cr is an austenite stabilizing element, and can increase the strength of steel or contribute to improving corrosion resistance within the range of an appropriate addition amount. However, as described above, Cr is a carbide forming element, and if added excessively to steel, it can form carbides at austenite grain boundaries, thereby reducing the low-temperature impact toughness of the steel. In addition, if the amount of Cr added exceeds a certain level, excessive carbides may be precipitated in the heat-affected zone (HAZ) of the weld, which may result in poor ultra-low-temperature toughness. Therefore, in the present invention, the upper limit of Cr may be 3.00%. In another embodiment, the upper limit of the Cr content may be 2.80%, and in yet another embodiment, the upper limit of the Cr content may be 2.50%.

[0044] V: 1.00~3.00%

[0045] V is an element that forms at least one of VC precipitates and VCN precipitates by combining with C, and is an element that suppresses grain growth of austenite structure and delays recrystallization, thereby contributing to strength improvement. In particular, when at least one of fine VC precipitates and VCN precipitates exceeding a certain fraction is formed, it can more effectively contribute to strength improvement of steel. In the present invention, for the purpose of achieving the above-described effect, the lower limit of the V content may be set to 1.00%. In another embodiment, the present invention may include 1.10% or more of V, and in another embodiment, it may include 1.20% or more.

[0046] On the other hand, if the content of V is excessively added, it is practically impossible to completely dissolve the coarse carbide formed during the steelmaking process during the reheating process, and the coarse carbide remains in the subsequent process, which may deteriorate the physical properties. In addition, since V is an expensive element, it may not be economically desirable when added in excess. Therefore, the upper limit of the V content in the present invention may be 3.00%, and in another embodiment, the upper limit of the V content may be 2.80%.

[0047] N: 1.000% or less (excluding 0%)

[0048] Nitrogen, along with carbon, stabilizes austenite, enhancing toughness. It is particularly advantageous for enhancing strength through solid solution strengthening, similar to carbon. It is also well known as an element that effectively increases stacking fault energy and promotes slip.

[0049] However, since there is a problem that coarse nitrides are formed when added in excess of 1.000%, which deteriorates the surface quality and physical properties of the steel, it is preferable that the upper limit be limited to 1.000%. The upper limit of the preferable nitrogen (N) content may be 0.500%, and the upper limit of the more preferable nitrogen (N) content may be 0.200%. The present invention does not separately specify the lower limit of the N content, but in consideration of the case where N is inevitably added, the lower limit of the nitrogen (N) content may be 0.005% or 0.007%.

[0050] Mo: 3.50% or less

[0051] Mo is an element that is incorporated into the matrix to increase strength. However, similar to Cr, if Mo is added in excess to steel, it may form carbides at austenite grain boundaries, thereby reducing the low-temperature impact toughness of the steel. In addition, if the amount of Mo added exceeds a certain level, excessive grain boundary carbides may precipitate in the heat-affected zone (HAZ) of the weld, resulting in poor ultra-low-temperature toughness. Therefore, the present invention may include Mo at 3.50% or less. In another embodiment, the upper limit of Mo may be 3.40%, and in another embodiment, it may be 3.20%. Since the purpose of the present invention can be achieved even when Mo is not added at all, the present invention does not specifically limit the lower limit of Mo. However, as an example, the lower limit of Mo may be 0.01%.

[0052] Nb: 1.00% or less

[0053] Nb is an element that forms NbC precipitates by combining with C, and is an element that suppresses grain growth of austenite structure and increases the recrystallization temperature, thereby increasing the amount of non-recrystallized rolling in the steel manufacturing process, thereby contributing to the improvement of strength. In addition, even if Nb is not added at all, there is no problem in achieving the purpose of the present invention, but when a certain fraction or more of fine NbC precipitates are formed, it can contribute more effectively to the improvement of the strength of the steel. Therefore, the present invention may limit the preferable lower limit of the Nb content to 0.01% in consideration of the precipitation strengthening phenomenon caused by the formation of precipitates.

[0054] If the content of Nb is excessive, coarse carbides formed during the steelmaking process may be vulnerable to external forces during the continuous casting process, causing cracks, which may deteriorate the quality of the cast steel. In addition, if a large amount of Nb is added, a large amount of coarse carbides may be precipitated in the weld heat-affected zone, which may reduce the impact toughness, which is not desirable. Therefore, the upper limit of the Nb content in the present invention may be limited to 1.00%, and the preferred upper limit of the Nb content may be 0.10%.

[0055] Meanwhile, the steel according to one embodiment of the present invention may additionally include at least one of Ti: 1.0% or less, Al: 5.0% or less, and Si: 5.0% or less.

[0056] Ti: 1.00% or less

[0057] Ti is an element that suppresses austenite grain growth by forming carbonitrides, which can help increase strength. Even if Ti is not included at all, it does not affect the achievement of the purpose of the present invention, and thus the present invention may not include the above-described Ti at all. On the other hand, if Ti is added excessively, the Ti may crystallize or coarsen, thereby deteriorating the quality of the cast steel. Therefore, the present invention may include Ti at 1.00% or less. In another embodiment, the present invention may include the above-described Ti at 0.005% to 0.90%, and in another embodiment, it may include 0.01% to 0.80%.

[0058] Al: 5.00% or less

[0059] Al can be incorporated into the matrix to increase the strength of the steel and increase the stacking fault energy, thereby increasing the strength that controls the deformation mode, and as a result, can cause the deformation mode to slip. Such Al can be included in the steel of the present invention to obtain the above-described effect, but even if Al is not included at all, there is no effect on achieving the purpose of the present invention, and thus the present invention may not include the above-described Al at all. On the other hand, if Al is added excessively, there may be a problem that coarse AlN crystallizes or precipitates, thereby deteriorating the quality of the cast steel, and therefore the present invention may include Al at 5.00% or less. In another embodiment, the present invention may include the above-described Al at 0.01% to 4.50%, and in another embodiment, it may include the above-described Al at 0.02% to 4.00%.

[0060] Si: 5.00% or less

[0061] Si is an element that improves the castability of molten steel and, especially, when added to austenitic steel, effectively increases the strength by being dissolved in the steel. It is also an element that affects the activity of carbon in the steel, effectively suppressing the formation of carbides and increasing the toughness. However, when added in excess of 5.00%, there is a problem that the stacking fault energy is reduced, promoting twinning, and a decrease in toughness may occur due to high strength. Therefore, it is desirable to limit the upper limit to 5.00%. As another example, the upper limit of the Si content may be 3.00%, and as another example, the upper limit of the Si content may be 2.50%. Furthermore, according to another example, the lower limit of the Si content may be 0.10%, and as another example, 0.30%.

[0062] In addition, according to another aspect of the present invention, the steel of the present invention can satisfy the following equation 1, which is a relationship between carbon (C) and manganese (Mn).

[0063] [Equation 1] 23.6[C]+[Mn] ≥ 28.000, 33.5[C]-[Mn] ≤ 23.00

[0064] (In the above formula 1, [C] and [Mn] represent the weight percentages of C and Mn contained in the steel, respectively.)

[0065] The present invention has conducted an in-depth study on the relative behavior between C and Mn contents in relation to carbide formation, and as a result, it has been found that determining the relative content relationship between C and Mn, as shown in Fig. 1, is an important factor in effectively controlling the amount of carbide precipitation while effectively promoting the stabilization of austenite.

[0066] In order to promote the stabilization of austenite, it is preferable to control the value of 23.6[C]+[Mn] to 28,000 or more, provided that other components meet the range specified in the present invention. If the value of 23.6[C]+[Mn] is less than 28,000, the stability of austenite decreases, causing a working-induced transformation due to deformation, which may lower the impact toughness of the steel. In addition, when a working-induced transformation occurs, alpha prime martensite with a BCC structure is formed in addition to non-magnetic austenite, which increases the permeability, so the non-magnetic properties of the steel may deteriorate.

[0067] In addition, in one embodiment of the present invention, the value of 33.5[C]-[Mn] may be 23.00 or less. This is to prevent deterioration of physical properties due to the formation of carbides caused by excessive carbon content. In another embodiment, the value of 33.5[C]-[Mn] may be 20.00 or less, and in another embodiment, it may be 18.00 or less.

[0068] The steel according to one aspect of the present invention may contain iron and other unavoidable impurities in addition to the aforementioned components. However, since unintended impurities may inevitably be mixed in from raw materials or the surrounding environment during the typical manufacturing process, they cannot be completely excluded. Since these impurities are readily apparent to anyone skilled in the art, their full contents are not specifically mentioned herein. Furthermore, the addition of additional effective components other than the aforementioned components is not completely excluded.

[0069] According to an example of the present invention, the microstructure of the steel may have a main structure of austenite.

[0070] This is to ensure that the steel according to one aspect of the present invention secures the desired properties. In one embodiment, the area fraction of the austenite may be 95 area% or more. In another embodiment, the area fraction of the austenite may be 97 area% or more. In particular, preferably, the area fraction of the austenite may be 100 area% for the purpose of ensuring non-magnetism, but is not necessarily limited thereto. The method for measuring the area fraction of the austenite in the present invention is not particularly limited, and can be easily confirmed through a measuring method commonly used by a person skilled in the art to which the present invention pertains for measuring microstructure and carbide.

[0071] In the case of a steel according to an example of the present invention, the area fraction of grain boundary carbides formed at the grain boundaries of the austenite may be 5.0 area% or less. In this case, as a non-limiting example, the area that serves as a reference when measuring the area fraction of the grain boundary carbides may be the entire measurement region.

[0072] That is, one example of the present invention can suppress the deterioration of low-temperature impact toughness and ultra-low-temperature toughness in the weld heat-affected zone by reducing the carbide formed at the austenite grain boundary by reducing the content of Cr, Mo, and Nb included in the steel. Therefore, the carbide formed at the austenite grain boundary described above may include at least one or more of Cr, Mo, and Nb carbides, and the area fraction of the grain boundary carbide may be 5.0 area% or less. In another embodiment, the area fraction of the grain boundary carbide may be 4.8 area% or less, and in another embodiment, it may be 4.5 area% or less. In addition, according to a non-limiting embodiment, the carbide formed at the austenite grain boundary described above may be measured using a scanning electron microscope or an optical microscope, and in consideration of this, the minimum value of the observable circle equivalent diameter of the grain boundary carbide may be 100 nm. The above equivalent diameter may refer to the diameter of a virtual circle having the same area as the grain boundary carbide exposed on the surface.

[0073] In addition, the steel according to an example of the present invention can form V-derived fine precipitates in the grains through aging treatment, and at this time, the V-derived fine precipitates may mean at least one of VCN fine precipitates and VC fine precipitates. In addition, the fine precipitates may mean precipitates having a diameter of 50.0 nm or less. At this time, the diameter may mean the equivalent diameter of a circle described above. As another example, the fine precipitates may mean precipitates having a diameter of 10.0 nm or less, and as another example, the fine precipitates may mean precipitates having a diameter of 5.0 nm or less. Meanwhile, since the diameter of the above-described fine precipitates is better the smaller it is, the lower limit thereof is not specifically limited. However, as a non-limiting example, when measuring the diameter using a transmission electron microscope, considering that the minimum diameter of the observable fine precipitates is 0.2 nm, the lower limit of the diameter of the fine precipitates may be 0.2 nm. The present invention can secure high strength properties of steel despite a decrease in Cr by forming such V-derived fine precipitates at a certain level or more. In particular, according to an example of the present invention, the present invention can form the V-derived fine precipitates at a certain level or more within the grains, and in this case, the high strength properties of the steel can be more efficiently secured.

[0074] More specifically, in order to secure the effect of the strength improvement described above, the steel according to one aspect of the present invention has at least one type of fine precipitate of VC and VCN having a diameter of 50.0 nm or less at 100 / mm. 2 This may include:

[0075] The method for measuring the number of fine precipitates per unit area of ​​at least one of the VC and VCN is not particularly limited because it can be easily adopted by a person skilled in the art according to the purpose, but as an example, it can be measured using a transmission electron microscope. In another embodiment, the present invention provides a method for measuring the number of fine precipitates per unit area of ​​at least one of the VC and VCN having a diameter of 50.0 nm or less at 120 / mm. 2 or more than 150 / mm 2 It may include, in another embodiment, 500 / mm 2 or more than 1000 / mm 2 The above may be included. According to one aspect of the present invention, the more fine precipitates of at least one kind of VC and VCN, the better, so the upper limit is not specifically limited, but considering that the number cannot be infinite in reality, the upper limit of the number of fine precipitates of at least one kind of VC and VCN having a diameter of 50.0 nm or less per unit area is 10,000 / mm. 2 It could be.

[0076] In addition, as described above, when reducing the grain boundary carbides, there is an advantage in that it is possible to prevent excessive carbides from precipitating in the heat affected zone (HAZ) of the weld after welding the steel, thereby deteriorating the ultra-low temperature toughness. More specifically, the steel according to one aspect of the present invention has a heat resistance of 3.0 kJ / mm 2 After performing submerged arc welding with a heat input of , the area fraction of grain boundary carbide in the weld heat affected zone can be 5.0 area% or less. As another example, the area fraction of grain boundary carbide in the weld heat affected zone can be 4.8 area% or less, and as another example, it can be 4.5 area% or less. With regard to the lower limit of the equivalent circle diameter of grain boundary carbide in the weld heat affected zone, the description with respect to carbide formed in the austenite grain boundary before welding is equally applicable and therefore is omitted.

[0077] As described above, the steel according to one aspect of the present invention can secure high strength characteristics and excellent impact toughness.

[0078] Specifically, the steel according to one aspect of the present invention may have a yield strength at room temperature of 550 MPa or more, and according to another example, may be 690 MPa or more, and according to another example, the lower limit of the yield strength at room temperature may be 750 MPa or 900 MPa. In addition, the Charpy impact energy value measured at -84°C may be 27 J or more.

[0079] Furthermore, the steel according to one example of the present invention can secure excellent non-magnetic properties. Specifically, the steel according to the present invention can have an investment rate of 1.200 or less after undergoing 20% ​​cold plastic deformation at room temperature.

[0080] As described above, the fact that the permeability is 1.200 or less after 20% cold plastic deformation at room temperature means that the austenite structure can be stably maintained even after cold forming. Specifically, since austenite is an unstable structure, even if austenite is obtained at room temperature, the austenite can be transformed into epsilon martensite or alpha martensite through additional cooling or processing. In this case, the impact toughness tends to be poor and the non-magnetic properties tend to deteriorate. Therefore, the steel according to an example of the present invention can lower the permeability by controlling the relative contents between C and Mn through the above-described Equation 1 as described above, thereby allowing the austenite structure to be stably maintained even after cold forming, thereby securing excellent non-magnetic properties.

[0081] As another example, the steel according to one embodiment of the present invention may have an investment ratio of 1.100 or less when the strain during cold plastic deformation at room temperature is at least 2%.

[0082] Lastly, since the steel according to one embodiment of the present invention has excellent hydrogen embrittlement resistance, a high level of safety can be secured when applied to a sour environment.

[0083] Specifically, the steel according to one embodiment of the present invention may have a crack length ratio (CLR) of 10% or less in a hydrogen induced cracking (HIC) test, a preferable crack length ratio (CLR) of 5% or less, and a more preferable crack length ratio (CLR) of 2% or less.

[0084] This hydrogen-induced cracking (HIC) test is performed by immersing steel in an acidic solution (5% NaCl + 0.5% CH3COOH) saturated with H2S gas, maintaining it for 96 hours, and then observing the number of cracks that occur. That is, the cross-section of a specimen with a constant width (W) and thickness (T) is observed, and the crack length (a) is measured at the point of crack occurrence, and the crack length ratio (CLR) is derived from the average value of these values. The formula for calculating the crack length ratio (CLR) is as follows: Equation 2.

[0085] [Formula 2]

[0086] CLR(Crack Length Ratio, %) = ∑(a / w) * 100

[0087] (In the above equation 2, a means the length of a single crack (㎛), and w means the width of the specimen (㎛).)

[0088] Hereinafter, a method for manufacturing steel according to one embodiment of the present invention will be described. However, the method for manufacturing steel described below is merely an example, and the steel of the present invention does not necessarily have to be manufactured using this manufacturing method. It should be noted that any manufacturing method that satisfies the claims of the present invention can be used to implement each embodiment of the present invention without any problems.

[0089] A method for manufacturing a steel according to one embodiment of the present invention may include the steps of: heating a slab comprising, in wt%, C: 0.050% to 1.70%, Mn: 15.0% to 40.0%, Cr: 3.00% or less, V: 1.00% to 3.00%, N: 1.000% or less (excluding 0%), Mo: 3.50% or less, and Nb: 1.00% or less, with the remainder being iron (Fe) and unavoidable impurities; finishing hot-rolling the slab to obtain a hot-rolled steel sheet; cooling the hot-rolled steel sheet to room temperature after solution treatment; and aging the hot-rolled steel sheet.

[0090] Below, each step is explained in detail.

[0091] Steps to heat the slab

[0092] A method for manufacturing steel according to one embodiment of the present invention first prepares a slab comprising C: 0.050% to 1.70%, Mn: 15.0 to 40.0%, Cr: 3.00% or less, V: 1.00 to 3.00%, N: 1.000% or less (excluding 0%), Mo: 3.50% or less, and Nb: 1.00% or less, with the remainder being iron (Fe) and unavoidable impurities, and then the slab can be heated.

[0093] In addition, the above slab may additionally include at least one of Ti: 1.00% or less, Al: 5.00% or less, and Si: 5.00% or less, and may satisfy the following formula 1. Since the composition of the slab is as described above, its description is omitted.

[0094] [Equation 1] 23.6[C]+[Mn] ≥ 28.000, 33.5[C]-[Mn] ≤ 23.00

[0095] (In the above formula 1, [C] and [Mn] represent the weight percentages of C and Mn contained in the steel, respectively.)

[0096] The above-described heating can be performed in a temperature range of 1000°C or more and 1300°C or less. If the slab heating temperature is less than 1000°C, there are disadvantages in that re-dissolution and homogenization of the alloy components may not occur or it may take a long time to reach the target temperature to the center of the slab. In another embodiment, the lower limit of the slab heating temperature may be 1050°C, and in another embodiment, the lower limit of the slab heating temperature may be 1100°C or 1150°C.

[0097] If the above slab heating temperature exceeds 1300°C, partial melting may occur in the slab alloy component segregation area or severe surface oxidation may occur. In addition, in another embodiment, the upper limit of the slab heating temperature may be 1250°C, and in another embodiment, the upper limit of the slab heating temperature may be 1230°C or 1200°C.

[0098] Steps for obtaining hot-rolled steel sheets

[0099] After the above-described heating, the method for manufacturing steel according to one aspect of the present invention can finish hot-roll the slab, thereby obtaining a hot-rolled steel sheet.

[0100] In addition, as a non-limiting example, the finishing hot rolling may be performed at a temperature of 700°C or higher and 1050°C or lower. When the finishing hot rolling temperature is lower than 700°C, rolling is not easy due to the high high-temperature strength of the material, and excessive rolling in the non-recrystallized region causes the strength of the material to increase excessively, which has the disadvantage of reducing impact toughness. In another embodiment, the lower limit of the finishing hot rolling temperature may be 750°C or 800°C.

[0101] If the above finishing hot rolling temperature exceeds 1050℃, there is a disadvantage in that austenite coarsens and the strength decreases. In addition, as another example, the upper limit of the above finishing hot rolling temperature of the slab may be 1000℃ or 950℃. Meanwhile, the reduction ratio during the above hot rolling may be applied within an appropriate range depending on the desired plate thickness, and as a non-limiting example, the final thickness of the hot-rolled hot-rolled steel plate may be 5 to 50 mm.

[0102] Step of cooling to room temperature after solution treatment

[0103] Next, in one example of the present invention, the hot-rolled steel sheet obtained by the above-described method can be subjected to a solution treatment and then cooled to room temperature.

[0104] In addition, as a non-limiting example, the solution treatment may be performed at a temperature of 900°C or higher and 1200°C or lower for a time of 30 minutes or higher and 2 hours or lower. The solution treatment is intended to dissolve the intergranular and intragranular coarse carbides generated during hot rolling into the parent material and to reduce the internal energy caused by dislocations excessively generated by rolling. In order to achieve the above-described object, the method for manufacturing a steel according to one aspect of the present invention may set the temperature of the solution treatment to 900°C or higher, and the time of the solution treatment to 30 minutes or longer. According to another embodiment, the lower limit of the solution treatment temperature may be 950°C or 1000°C.

[0105] On the other hand, if the solution treatment temperature exceeds 1200°C or the treatment time exceeds 2 hours, there is a disadvantage in that the austenite becomes excessively coarsened, resulting in a decrease in strength. In addition, according to another embodiment, the upper limit of the solution treatment temperature described above may be 1180°C, and according to another embodiment, it may be 1150°C.

[0106] statute of limitations processing stage

[0107] A method for manufacturing steel according to one embodiment of the present invention may include, after the solution treatment, a step of aging the hot-rolled steel sheet. This aging treatment is intended to enhance the strength of the austenitic steel by precipitating at least one type of fine precipitate, among VC and VCN, within the grains.

[0108] As a non-limiting example, the aging treatment may be performed at a temperature of 500°C or higher and 850°C or lower for a time of 30 minutes or higher and 5 hours or lower. If the aging treatment temperature is lower than 500°C, the diffusion of precipitated elements is not easy, which has the disadvantage of taking too long a time for precipitation. In another embodiment, the lower limit of the aging treatment temperature may be 550°C or 600°C. If the aging treatment temperature exceeds 850°C, there is the disadvantage of excessive coarsening of austenite, which reduces the strength. In addition, in another embodiment, the upper limit of the aging treatment temperature may be 830°C or 800°C.

[0109] If the aging treatment time is less than 30 minutes, there is a disadvantage in that sufficient time for precipitation cannot be secured. In another embodiment, the lower limit of the aging treatment time may be 36 minutes or 42 minutes. On the other hand, the upper limit of the aging treatment time is preferably 5 hours or less. If it exceeds 5 hours, there is a problem that the strength actually decreases due to overaging, which is uneconomical. In another embodiment, the upper limit of the aging treatment time may be 4.8 hours or 4.5 hours.

[0110] Hereinafter, a steel material and its manufacturing method according to one aspect of the present invention will be described in more detail through specific examples. It should be noted that the following examples are intended solely to facilitate understanding of the present invention and are not intended to define the scope of the invention. The scope of the invention can be determined by the matters set forth in the patent claims and matters reasonably inferred therefrom.

[0111] After preparing a 250 mm thick slab having the alloy composition described in Table 1 below, the slab was heated, hot rolled, and solution treated under the conditions described in Table 2 below, cooled to room temperature, and then aging treated to manufacture steel. However, in the case of Comparative Example 7, solution treatment and aging treatment were not performed. The microstructure and physical properties of the steel thus manufactured were measured, and the results are shown in Tables 3 and 4 below.

[0112] At this time, the area fraction (area %) of the austenite was measured five times at room temperature using an optical microscope at a point 1 / 4t from the specimen thickness after taking a specimen of approximately 2 cm in width and length at the midpoint of the steel plate width and length. At this time, the area of ​​the austenite was calculated through image processing, and the average of the measured values ​​is shown in Table 3 below. At this time, the measurement magnification of the optical microscope was 200 times.

[0113] In addition, the number of fine precipitates of at least one type among VC and VCN per unit area (number / mm) 2 ), the specimen was collected in the same manner as the above-mentioned method, and measured 5 times using a transmission electron microscope at a point 1 / 4t from the specimen thickness, and the number of precipitates per unit area was calculated through image processing, and then the average was taken for measurement. At this time, the measurement magnification of the transmission electron microscope was 200,000 times.

[0114] In addition, for the area fraction (area %) of austenite grain boundary carbides, specimens were collected in the same manner as the above-mentioned collection method, and measured 5 times using a scanning electron microscope at 1 / 4t of the specimen thickness, and the area fraction of the precipitates was calculated through image processing and then averaged to derive the result. The grain boundary carbides were at least one of Cr, Mo, and Nb carbides, and the measurement magnification of the scanning electron microscope was 2000 times.

[0115] The yield strength at room temperature in Table 3 below was measured using a uniaxial tensile test method, and the Charpy impact energy was measured using a Charpy impact tester after the specimen was maintained at -84℃ for more than 15 minutes. In addition, the crack length ratio (CLR) in the hydrogen-induced cracking (HIC) test was derived by observing the number of cracks that occurred after immersing the specimen in an acidic solution (5% NaCl + 0.5% CH3COOH) saturated with H2S gas and maintaining it for 96 hours, using Equation 2 below.

[0116] [Formula 2]

[0117] CLR(Crack Length Ratio, %) = ∑(a / W) * 100

[0118] (In Equation 2, a represents the length of a single crack (㎛), and w represents the width of the specimen (㎛).)

[0119] Afterwards, the manufactured steel was subjected to cold plastic deformation at a strain rate of 20% using a uniaxial tensile tester at room temperature, and then the investment rate of the steel was measured, which is shown in Table 3 below.

[0120] The investment rate of these steels was obtained by cutting cold-deformed specimens, preparing samples with a size of approximately 2 cm in width and length, and measuring them three times using an investment rate measuring device, and then deriving the average value.

[0121] Additionally, 3.0kJ / mm for manufactured steel 2 Welding was performed using the submerged arc welding method with a heat input of , and then the area fraction (area %) of grain boundary carbides in the weld heat-affected zone was measured and shown in Table 3 below. The grain boundary carbides were at least one of Cr, Mo, and Nb carbides. The area fraction of grain boundary carbides in the weld heat-affected zone was measured five times using the same method as the base metal, and the average of the measured values ​​was derived.

[0122] Classification Alloy composition (weight %) CMn23.6C+Mn33.5C-MnCrMoVNbTiAlSiN Example 10.58 26.94 0.58 8-7.47--2.07----0.030 Example 20.62 28.14 2.73 2-7.33 0.48 3.05 1.98 0.02---0.014 Example 30.523 0.242.472-12.782.010.042.31-0.02--0.016Example 40.7132.048.756-8.222.232.062.050.030.02-0.910.016Example 50.4825.837.128-9.721.521.021.28---1.120.017 Comparative Example 10.26 19.22 5.33 6-10.49 1.36-----0.37 0.015 Comparative Example 20.49 28.54 0.064-12.09 3.52-2.02----0.017 Comparative Example 31.35 18.35 0.16 26.93 2.15--0.02 0.04--0.018 Comparative Example 40.55 27.3 40.28-8.88--2.15----0.015Comparative Example 50.4213.523.4120.57-------0.009Comparative Example 60.5827.641.29-8.170.515.232.11----0.016Comparative Example 70.5527.039.98-8.58--2.11----0.018

[0123] Slab heating temperature (℃)Hot rolling finishing temperature (℃)Solution treatment temperature (℃)Solution treatment time (hour)Aging temperature (℃)Aging time (hour)Steel thickness (mm)Example 1 1 1 5 2 8 8 6 1 1 4 8 0.52 7 6 3 0.55 12Example 2 1 1 4 8 9 0.51 1 6 2 0.74 7 24 0.82 20Example 3 1 1 7 5 9 1 5 1 1 5 0 0.69 8 12 0.77 15Example 4 1 1 8 0 9 2 1 1 5 5 1.02 7 20 0.66 18Example 5 1 1 8 5 8 5 9 1 1 5 4 0.87 7 49 1.02 22Comparative example 1 1 1 8 2 8 9 5 1 1 5 2 0.63 7480.5612Comparison Example 2 114788911750.688641.3225Comparison Example 3 120591011821.138370.8830Comparison Example 4 114989411530.589506.735Comparison Example 5 118490811630.767381.4210Comparison Example 6 115589311650.997651.0324Comparison Example 7 1168892----12

[0124] Microstructure (area %) VC or VCN micro-precipitates number (units / mm) 2 ) Area fraction of grain boundary carbide (area%) Area fraction of grain boundary carbide (area%) in the weld heat-affected zone Example 195% or more γ89 21.0% or less 1.0% or less Example 295% or more γ10 121.0% or less 1.0% or less Example 395% or more γ110 91.0% or less 1.0% or less Example 495% or more γ150 51.0% or less 1.0% or less Example 595% or more γ54 81.0% or less 1.0% or less Comparative Example 195% or more γ-1.0% or less 1.0% or less Comparative Example 295% or more γ85 3.0% or less 3.0% or less Comparative Example 394% γ-5.0% or more 5.0% or more Comparative Example 495% or more γ-1.0% or less 1.0% Below Comparative Example 584%γ-1.0% or less1.0% or lessComparative Example 693%γ6205.0% or more5.0% or moreComparative Example 795% or moreγ301.0% or less1.0% or less

[0125] Relative permeability, room temperature yield strength (MPa), Charpy impact energy (-84 o C, J)CLR(%)Example 11.002762690Example 21.002911720Example 31.002915680Example 41.0021117580Example 51.002778890Comparative Example 11.52342150Comparative Example 21.005489120Comparative Example 31.00354650Comparative Example 41.0024231130Comparative Example 53.6349236Comparative Example 61.004986150Comparative Example 71.0024881120

[0126] Looking at Tables 1 to 3 above, Comparative Examples 1, 3 and 5, in which no V was added at all, were unable to secure high strength characteristics because at least one type of fine precipitate among VC and VCN was not formed at all.

[0127] In addition, it can be seen that in Comparative Examples 2 and 4, although V was added, the strength actually decreased due to overaging as the aging temperature was excessively high or the aging time was excessively long.

[0128] In Comparative Example 6, since Mo exceeded the range proposed in the present invention, carbides exceeded 5 area% at the austenite grain boundaries. As a result, excellent toughness could not be secured.

[0129] Lastly, in the case of Comparative Example 7, since the solution treatment and aging treatment were not performed, it was not possible to secure at least one type of fine VC or VCN precipitate to a certain level, and accordingly, the yield strength at room temperature was less than 550 MPa.

[0130] On the other hand, in the case of Examples 1 to 5 that satisfy the alloy composition and manufacturing conditions of the present invention, excellent physical properties could be secured by securing the components and microstructure desired by the present invention.

Claims

1. Contains, by weight%, C: 0.050% to 1.70%, Mn: 15.0% to 40.0%, Cr: 3.00% or less, V: 1.00% to 3.00%, N: 1.000% or less (excluding 0%), Mo: 3.50% or less, and Nb: 1.00% or less, with the remainder being iron (Fe) and unavoidable impurities. The microstructure is mainly composed of austenite. The area fraction of grain boundary carbides formed at the grain boundaries of the above austenite is 5.0 area% or less, At least one type of fine precipitate of VC or VCN with a diameter of 50.0 nm or less, 100 / mm 2 Steel containing the number per unit area above.

2. In paragraph 1, Steel additionally containing at least one of Ti: 1.00% or less, Al: 5.00% or less, and Si: 5.00% or less.

3. In paragraph 1, Steel satisfying the following equation 1. [Equation 1] 23.6[C]+[Mn] ≥ 28.000, 33.5[C]-[Mn] ≤ 23.00 (In the above formula 1, [C] and [Mn] represent the weight percentages of C and Mn contained in the steel, respectively.) 4. In paragraph 1, The above grain boundary carbide is a steel material containing at least one of Cr, Mo and Nb carbides.

5. In paragraph 1, Steel having an area fraction of the above austenite of 95 area% or more.

6. In paragraph 1, Steel with a yield strength of 550 MPa or more at room temperature.

7. In paragraph 1, Steel having a Charpy impact energy value of 27 J or more at -84℃.

8. In paragraph 1, Steel having an investment ratio of 1.2 or less after undergoing 20% ​​cold plastic deformation at room temperature.

9. In paragraph 1, Steel having an investment ratio of 1.100 or less when the strain during cold plastic deformation at room temperature is at least 2%.

10. In paragraph 1, 3.0kJ / mm 2 Steel having an area fraction of grain boundary carbides in the weld heat affected zone of 5.0 area% or less after submerged welding with a heat input of .

11. In paragraph 1, Steel having a crack length ratio (CLR) of 10% or less in a hydrogen-induced cracking (HIC) test defined by Equation 2 below. [Formula 2] CLR(Crack Length Ratio, %) = ∑(a / W) * 100 (In the above equation 2, a means the length of a single crack (㎛), and w means the width of the specimen (㎛).) 12. A step of heating a slab comprising C: 0.050% to 1.70%, Mn: 15.0% to 40.0%, Cr: 3.00% or less, V: 1.00% to 3.00%, N: 1.000% or less (excluding 0%), Mo: 3.50% or less, and Nb: 1.00% or less, with the remainder being iron (Fe) and unavoidable impurities; A step of obtaining a hot-rolled steel sheet by finishing hot rolling the above slab; A step of cooling the hot-rolled steel sheet to room temperature after solution treatment; and A method for manufacturing steel, comprising: a step of aging the hot-rolled steel sheet.

13. In paragraph 12, A method for manufacturing a steel material, wherein the above slab additionally contains at least one of Ti: 1.00% or less, Al: 5.00% or less, and Si: 5.00% or less.

14. In paragraph 12, The above slab is a method for manufacturing steel satisfying the following formula 1. [Equation 1] 23.6[C]+[Mn] ≥ 28.000, 33.5[C]-[Mn] ≤ 23.00 (In the above formula 1, [C] and [Mn] represent the weight percentages of C and Mn contained in the steel, respectively.) 15. In paragraph 12, The above heating is performed at a temperature of 1000℃ or higher and 1300℃ or lower, The above finishing hot rolling is performed at a temperature of 700℃ or higher and 1050℃ or lower. The above solution treatment is performed at a temperature of 900°C or higher and 1200°C or lower for 30 minutes or longer and 2 hours or less. A method for manufacturing steel in which the above aging treatment is performed at a temperature of 500℃ or higher and 850℃ or lower for 30 minutes or longer and 5 hours or shorter.

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