KR20250065516A - Austenitic steel material and manufacturing method for the same - Google Patents

Abstract

The present invention relates to an austenitic steel that can be applied to various fields such as the energy industry and shipbuilding, and a method for manufacturing the same, and more specifically, to an austenitic steel that can be used in structures that require high strength and excellent non-magnetic properties, such as a ship hull, or a steel pipe or facility for transporting crude oil with a high hydrogen sulfide content, and a method for manufacturing the same.

Description

{AUSTENITIC STEEL MATERIAL AND MANUFACTURING METHOD FOR THE SAME}

The present invention relates to an austenitic steel that can be applied to various fields such as energy industry, shipbuilding, and marine structures, and a method for manufacturing the same. More specifically, the present invention relates to an austenitic steel that can be applied to ship hulls requiring excellent non-magnetic properties, or steel pipes or facilities for transporting crude oil with a high hydrogen sulfide content, and a method for manufacturing the same.

As the mining environment for crude oil becomes increasingly sour, containing large amounts of hydrogen sulfide, etc., the steel pipes and structures used to transport it require materials with excellent hydrogen embrittlement resistance. Therefore, in order to transport these crude oils, materials with excellent hydrogen embrittlement resistance and sufficient toughness and strength are required to resist external pressure during the transport process.

Existing high-strength carbon steels have limitations in that they are difficult to produce and have limited use due to the presence of hard spots inside the steel that are vulnerable to hydrogen embrittlement. To overcome these limitations, various manufacturing methods and detection methods for hard spots have been proposed, but fundamentally, hard spots are unavoidable in order to increase strength, making it difficult to produce steel that has both excellent hydrogen embrittlement resistance and high strength characteristics.

Also, in the case of structures that use a large amount of steel materials, such as military vessels such as submarines, the location of the other party can be tracked by detecting changes in the magnetic field caused by the interaction between the hull and the Earth's magnetic field. If such structures are made of 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, a process called demagnetization is performed to periodically remove magnetism, and this reduces the operating time of the vessel and also causes economic losses. Therefore, the demand for steel with excellent non-magnetic properties is increasing.

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 permeability even after deformation, and a method for manufacturing the same.

The tasks of the present invention are not limited to the above-described contents. Those with ordinary knowledge in the technical field to which the present invention belongs will have no difficulty in understanding additional tasks of the present invention from the overall contents of this specification.

According to one aspect of the present invention, a steel material contains, in wt%, C: 0.050% to 1.7%, Mn: 15 to 40%, Cr: 3.0% or less, V: 1.0 to 3.0%, N: 1.000% or less (excluding 0%), Mo: 3.5% or less, and Nb: 1.0% or less, with the remainder being iron (Fe) and unavoidable impurities, and a microstructure whose 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 may include at least one or more kinds of fine precipitates of VC and VCN having a diameter of 50 nm or less in a number per unit area of 100/ ^mm2 or more.

The above-described steel may additionally contain at least one of Ti: 1.0% or less, Al: 5.0% or less, and Si: 5.0% or less.

The steel material described above can satisfy the following equation 1.

[Equation 1] 23.6[C]+[Mn] ≥ 28, 33.5[C]-[Mn] ≤ 23

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

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

The area fraction of the austenite described above can be 95 area% or more.

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 investment rate after 20% cold plastic deformation at room temperature may be 1.2 or less.

In addition, the steel described above may have an investment ratio of 1.1 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 Equation 2 below may be 10% or less.

[Formula 2]

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

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

In addition, the steel 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.

A method for manufacturing steel according to another aspect of the present invention may include the steps of: heating a slab comprising, in wt%, C: 0.050% to 1.7%, Mn: 15 to 40%, Cr: 3.0% or less, V: 1.0 to 3.0%, N: 1.000% or less (excluding 0%), Mo: 3.5% or less, and Nb: 1.0% or less, with the remainder being iron (Fe) and unavoidable impurities; subjecting the slab to a finish hot rolling process to obtain a hot-rolled steel sheet; subjecting the hot-rolled steel sheet to a solution treatment and then cooling it to room temperature; and subjecting the hot-rolled steel sheet to an aging treatment.

The above-described slab may additionally include at least one of Ti: 1.0% or less, Al: 5.0% or less, and Si: 5.0% or less, and may satisfy the following equation 1.

[Equation 1] 23.6[C]+[Mn] ≥ 28, 33.5[C]-[Mn] ≤ 23

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

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

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

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 process of explaining specific embodiments of the present invention.

Figure 1 is a graph showing the range of carbon and manganese according to one aspect of the present invention.
FIG. 2 is a transmission electron microscope photograph of Example 1, a steel material according to one aspect of the present invention.

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.

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.

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

In the past, austenitic steels with high manganese content were added with high chromium content to improve corrosion resistance and strength, and molybdenum and niobium, which play a similar role to the aforementioned chromium, also tended to be added to improve the strength of the steel.

However, in this case, there was a problem in securing excellent impact toughness of the steel because chromium, molybdenum or niobium combined with carbon to form carbides at the austenite grain boundaries.

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.

From this point of view, a steel according to one embodiment of the present invention may include, in wt%, C: 0.050% to 1.7%, Mn: 15% to 40%, Cr: 3.0% or less, V: 1.0% to 3.0%, N: 1.000% or less (excluding 0%), Mo: 3.5% or less, and Nb: 1.0% or less, with the remainder being iron (Fe) and unavoidable impurities.

Below, each component is described in detail.

C: 0.050%~1.7%

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, and also easily causing working-induced transformation into epsilon or alpha martensite by external stress, thereby reducing the 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 0.10% or more. On the other hand, if the content of C is excessive, the toughness may rapidly deteriorate due to carbide precipitation, and the workability may deteriorate due to an excessive increase in strength, so the present invention may include C in an amount of 1.7% or less. In another embodiment, the present invention can set the upper limit of C to 1.5% or less, and in another embodiment, to 1.3%.

Mn: 15~40%

Mn is an element that plays an important role in stabilizing austenite. In one embodiment of the present steel, it is preferable that 15% or more of Mn is included to stabilize austenite. If the Mn content is lower than this, epsilon martensite, which is a metastable phase, is formed and easily transformed into alpha martensite by processing-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% or more. In another embodiment, the Mn content may be 18% or more, and in another embodiment, the Mn content may be 20% or more.

If the Mn content is excessive, it may not only reduce the corrosion rate of the steel, but is also undesirable from an economical perspective. Therefore, the Mn content included in the steel according to one aspect of the present invention may be 40% or less. In another embodiment, the Mn content may be 35% or less, and in another example, it may be 30% or less.

Cr: 3.0% or less

Cr is an austenite stabilizing element, and can increase the strength of steel or contribute to improving corrosion resistance within an appropriate addition amount. However, as described above, Cr is a carbide forming element, and if it is excessively added to steel, it can form carbides at austenite grain boundaries, thereby reducing the low-temperature impact toughness of the steel. In addition, if the addition amount of Cr exceeds a certain level, excessive carbides may be precipitated in the heat-affected zone (HAZ) of the weld, which may deteriorate the ultra-low-temperature toughness. Therefore, the upper limit of Cr in the present invention may be 3.0%. In another embodiment, the upper limit of the Cr content may be 2.8%, and in yet another embodiment, the upper limit of the Cr content may be 2.5%.

V: 1.0~3.0%

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, the lower limit of the V content can be set to 1.0% for the purpose of achieving the above-described effect. In another embodiment, the present invention can include V at 1.1% or more, and in another embodiment, it can include V at 1.2% or more.

On the other hand, if the content of V is excessively added, it is practically impossible to completely dissolve the coarse carbide formed in 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 excessive amounts. Therefore, the upper limit of the V content in the present invention may be 3.0%, and in another embodiment, the upper limit of the V content may be 2.8%.

N: 1.000% or less (excluding 0%)

N is an element that stabilizes austenite along with carbon and improves toughness, and is particularly advantageous in improving strength through solid solution strengthening, like carbon. In particular, it is well known as an element that effectively increases stacking fault energy and promotes slip.

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 is limited to 1.000%. The upper limit of the preferable nitrogen (N) content may be 0.5000%, and the upper limit of the more preferable nitrogen (N) content may be 0.2000%. The lower limit of the preferable nitrogen (N) content may be 0.005000%, and the lower limit of the more preferable nitrogen (N) content may be 0.007000%.

Mo: 3.5% or less

Mo is an element that is employed in a matrix to increase strength. However, similar to Cr, if Mo is excessively added 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 be precipitated in the heat-affected zone (HAZ) of the weld, thereby deteriorating the ultra-low-temperature toughness. Therefore, the present invention may include Mo at 3.5% or less. In another embodiment, the present invention may set the upper limit of Mo to 3.4%, and in another embodiment, it may set it to 3.2%.

Nb: 1.0% or less

Nb is an element that forms NbC precipitates by combining with C, and it is an element that suppresses grain growth of austenite structure, raises the recrystallization temperature, increases the amount of non-recrystallized rolling in the steel manufacturing process, and contributes to improving the strength. In addition, even if Nb is not added at all, there is no problem in achieving the purpose of the present invention, but if a certain fraction or more of fine NbC precipitates are formed, it can contribute more effectively to improving the strength of the steel. Therefore, the present invention may limit the preferable lower limit of the Nb content to 0.010% in consideration of the precipitation strengthening phenomenon due to the formation of precipitates.

If the content of Nb is excessively added, coarse carbides formed in the steelmaking process may act vulnerablely to external force in 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 lower the impact toughness, which is not desirable. Therefore, the upper limit of the Nb content of the present invention may be limited to 1.0%, and the preferred upper limit of the Nb content may be 0.10%.

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.

Ti: 1.0% or less

Ti is an element that suppresses austenite grain growth by forming carbonitride, and can help increase strength. Even if Ti is not included at all, it does not have any effect on achieving the purpose of the present invention, so the lower limit of the content of Ti described above can be 0%. On the other hand, if Ti is excessively added, the Ti may be crystallized or coarsened, thereby deteriorating the quality of the cast steel, so the present invention can include Ti at 1.0% or less. In another embodiment, the present invention can include Ti described above at 0.0050% to 0.90%, and in another embodiment, it can include Ti described above at 0.010% to 0.80%.

Al: 5.0% or less

Al can be employed in a matrix to increase the strength of the steel and increase the stacking fault energy to increase the strength that controls the deformation mode, and as a result, can slip the deformation mode. 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 therefore the lower limit of the content of Al described above can be 0%. On the other hand, when Al is added excessively, there may be a problem that coarse AlN is crystallized or precipitated, deteriorating the quality of the cast steel, and therefore the present invention can include Al to be 5.0% or less. In another embodiment, the present invention can include the above-described Al to be 0.010% to 4.5%, and in still another embodiment, it can include the above-described Al to be 0.020% to 4.0%.

Si: 5.0% or less

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. In addition, it is an element that affects the activity of carbon in the steel, effectively suppresses the formation of carbides, and increases toughness. However, when added in excess of 5.0%, it reduces the stacking fault energy, promotes twinning, and there is a problem that the toughness may be lowered due to high strength, so it is desirable to limit the upper limit to 5.0%. As another example, the upper limit of the Si content may be 3.0%, and as another example, the upper limit of the Si content may be 2.5%. In addition, the lower limit of the Si content according to another example may be 0.10%, and as another example, 0.30%.

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).

[Equation 1] 23.6[C]+[Mn] ≥ 28, 33.5[C]-[Mn] ≤ 23

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

The present invention has conducted an in-depth study on the relative behavior between C and Mn contents with respect 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 promoting austenite stabilization while also effectively controlling the amount of carbide precipitation.

In order to promote stabilization of austenite, it is preferable to control the value of 23.6[C]+[Mn] to 28 or more, provided that other components satisfy the range specified in the present invention. When the value of 23.6[C]+[Mn] is less than 28, 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, in addition to non-magnetic austenite, alpha prime martensite with a BCC structure is formed, which increases the permeability, so the non-magnetic properties of the steel may deteriorate.

In addition, in one embodiment of the present invention, the value of 33.5[C]-[Mn] may be 23 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 or less, and in another embodiment, it may be 18 or less.

The steel according to one aspect of the present invention may contain the remaining Fe 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 a normal manufacturing process, they cannot be completely excluded. Since these impurities are known to anyone with ordinary knowledge in the art, not all of the contents are specifically mentioned in this specification. In addition, additional addition of effective components other than the aforementioned components is not completely excluded.

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

This is to secure the desired properties of the steel according to one aspect of the present invention. According to 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, for the purpose of securing non-magnetism, the area fraction of the austenite may be 100 area%, 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 belongs for measuring microstructure and carbide.

In the case of a steel according to an example of the present invention, the area fraction of grain boundary carbide formed at the grain boundary 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 carbide may be the entire measurement area.

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 contents of Cr, Mo, and Nb included in the steel, thereby reducing the carbides formed at the austenite grain boundaries. Therefore, the carbides formed at the austenite grain boundaries described above may include at least one of Cr, Mo, and Nb carbides, and the area fraction of the grain boundary carbides may be 5.0 area% or less. In another embodiment, the area fraction of the grain boundary carbides 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 carbides formed at the austenite grain boundaries 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 carbides may be 100 nm. The above equivalent circle diameter may refer to the diameter of a virtual circle having the same area as the grain boundary carbide exposed on the surface.

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 nm or less. At this time, the diameter may mean the equivalent circular diameter described above. As another example, the fine precipitates may mean precipitates having a diameter of 10 nm or less, and as another example, the precipitates having a diameter of 5.0 nm or less. Meanwhile, since the diameter of the fine precipitates described above is the smaller the better, the lower limit thereof is not specifically limited, but according to a non-limiting example, when measuring the diameter using a transmission electron microscope, considering that the minimum diameter of the fine precipitates that can be observed 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 to a certain level or higher. In particular, according to one example of the present invention, the present invention can form the V-derived fine precipitates to a certain level or higher within grains, and in this case, the high strength properties of the steel can be secured more efficiently.

More specifically, in order to secure the effect of strength improvement described above, the steel according to one aspect of the present invention may include at least 100/ ^mm2 or more of fine precipitates of at least one type of VC or VCN having a diameter of 50 nm or less.

The method for measuring the number of fine precipitates of at least one kind of VC, VCN per unit area is not particularly limited because it can be easily adopted by a person skilled in the art according to the purpose. As an example, the measurement can be performed using a transmission electron microscope. In another embodiment, the present invention can include at least one kind of fine precipitate of VC, VCN having a diameter of 50 nm or less, at 120/ ^mm2 or more or 150/ ^mm2 or more, and in another embodiment, at 500/ ^mm2 or more or 1000/ ^mm2 or more. According to one aspect of the present invention, the more fine precipitates of at least one of the VC and VCN, the better, and therefore the upper limit thereof is not specifically limited. However, considering that the number cannot be infinite in reality, the upper limit of the number per unit area of at least one of the VC and VCN having a diameter of 50 nm or less may be 10,000/ ^mm2 .

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

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

Specifically, the steel according to one aspect of the present invention may have a room temperature yield strength 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 room temperature yield strength may be 750 MPa or 900 MPa. In addition, the Charpy impact energy value measured at -84°C may be 27 J or more.

In addition, 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 1.2 or less after undergoing 20% cold plastic deformation at room temperature.

As described above, the fact that the investment ratio is 1.2 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 deteriorate and the non-magnetic properties tend to deteriorate. Therefore, the steel according to an example of the present invention can lower the investment ratio by controlling the relative content 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.

As another example, the steel according to one embodiment of the present invention may have an investment ratio of 1.1 or less when the strain is at least 2% during cold plastic deformation at room temperature.

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.

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.

This hydrogen-induced cracking (HIC) test is performed by immersing steel in an acid solution (5% NaCl + 0.5% _CH3 COOH) 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 equation for calculating the crack length ratio (CLR) is as shown in Equation 2 below.

[Formula 2]

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

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

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 only an example, and it is not necessary for the steel of the present invention to be manufactured by this manufacturing method. It should be noted that any manufacturing method that satisfies the scope of the claims of the present invention can be used to implement each embodiment of the present invention without any problem.

A method for manufacturing steel according to one embodiment of the present invention may include the steps of: heating a slab comprising, in wt%, C: 0.05 to 1.7%, Mn: 15 to 40%, Cr: 3.0% or less, V: 1.0 to 3.0%, N: 1.000% or less (excluding 0%), Mo: 3.5% or less, and Nb: 1.0% or less, with the remainder being iron (Fe) and unavoidable impurities; subjecting the slab to a finish hot rolling process to obtain a hot-rolled steel sheet; subjecting the hot-rolled steel sheet to a solution treatment and then cooling it to room temperature; and subjecting the hot-rolled steel sheet to an aging treatment.

Below, each step is explained in detail.

Steps to heating the slab

According to one embodiment of the present invention, a method for manufacturing steel includes first preparing a slab comprising C: 0.05 to 1.7%, Mn: 15 to 40%, Cr: 3.0% or less, V: 1.0 to 3.0%, N: 1.000% or less (excluding 0%), Mo: 3.5% or less, and Nb: 1.0% or less, with the remainder being iron (Fe) and unavoidable impurities, and then heating the slab.

In addition, the slab may additionally include at least one of Ti: 1.0% or less, Al: 5.0% or less, and Si: 5.0% or less, and may satisfy the following Equation 1. Since the composition of the slab is as described above, its description is omitted.

[Equation 1] 23.6[C]+[Mn] ≥ 28, 33.5[C]-[Mn] ≤ 23

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

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 are not achieved or it takes 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 can be 1050°C, and in another embodiment, the lower limit of the slab heating temperature can be 1100°C or 1150°C.

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

Steps for obtaining hot rolled steel sheets

After the heating described above, 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.

In addition, as a non-limiting example, the finishing hot rolling can be performed at 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 since excessive rolling in the non-recrystallized region occurs, the strength of the material increases excessively, which has the disadvantage of reducing the impact toughness. As another example, the lower limit of the finishing hot rolling temperature can be 750°C or 800°C.

If the above finishing hot rolling temperature exceeds 1050℃, there is a disadvantage that austenite coarsens and the strength decreases. In addition, as another example, the upper limit of the above slab finishing hot rolling temperature may be 1000℃ or 950℃. Meanwhile, the reduction ratio during the above hot rolling may be applied within an appropriate range depending on the intended 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.

Step of cooling to room temperature after solution treatment

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

In addition, as a non-limiting example, the solution treatment can be performed at 900°C or higher and 1200°C or lower for 30 minutes or higher and 2 hours or lower. The solution treatment is to dissolve grain boundary and intragranular coarse carbides generated during hot rolling into the parent material and to reduce internal energy caused by dislocations excessively generated by rolling. In order to achieve the above-described purpose, the method for manufacturing a steel according to one aspect of the present invention can 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 described above can be 950°C or 1000°C.

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 and the strength decreases. 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.

Steps to handle statute of limitations

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

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, there is a disadvantage that the time required for precipitation becomes excessively long because diffusion of precipitated elements is not easy. 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 a disadvantage that austenite becomes excessively coarsened and the strength decreases. In addition, in another embodiment, the upper limit of the aging treatment temperature may be 830°C or 800°C.

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 example, the lower limit of the aging treatment time may be 36 minutes or 42 minutes. On the other hand, it is preferable that the upper limit of the aging treatment time is 5 hours or less. If it exceeds 5 hours, there is a problem that the strength actually decreases due to overaging and it is uneconomical. In another example, the upper limit of the aging treatment time may be 4.8 hours or 4.5 hours.

Hereinafter, the steel 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 examples below are only for understanding the present invention and are not intended to specify the scope of the rights of the present invention. The scope of the rights of the present invention can be determined by the matters described in the patent claims and matters reasonably inferred therefrom.

(Example)

After preparing a 250 mm thick slab having the alloy composition described in Table 1 below, the slab was heated and hot rolled, 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-treating and aging-treating 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.

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

In addition, for the number of micro precipitates per unit area (number/ ^mm2 ) of at least one type of VC, VCN, a specimen was collected in the same manner as the above-mentioned collection method, and the measurement was performed five times using a transmission electron microscope at 1/4t of the specimen thickness, and the number of precipitates per unit area was calculated through image processing, and then the average was measured. At this time, the measurement magnification of the transmission electron microscope was 200,000 times.

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 five 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 the average was derived. 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.

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 maintaining the specimen 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.

[Formula 2]

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

(a in Equation 2 represents the length of a single crack (㎛), and w represents the width of the specimen (㎛).)

Thereafter, the manufactured steel was subjected to cold plastic deformation at a strain 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.

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

In addition, welding was performed on the manufactured steel by the submerged arc welding method with a heat input of 3.0 kJ/ ^mm2 , and then the area fraction (area %) of grain boundary carbides in the weld heat-affected zone was measured, which is 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 in the same manner as the base metal, and the average of the measured values was derived.

division Alloy composition (weight%) C Mn 23.6C+Mn 33.5C-Mn Cr Mo V Nb Ti Al Si N Example 1 0.58 26.9 40.588 -7.47 2.07 0.030 Example 2 0.62 28.1 42.732 -7.33 0.48 3.05 1.98 0.02 0.014 Example 3 0.52 30.2 42.472 -12.78 2.01 0.04 2.31 0.02 0.016 Example 4 0.71 32.0 48.756 -8.22 2.23 2.06 2.05 0.03 0.02 0.91 0.016 Example 5 0.48 25.8 37.128 -9.72 1.52 1.02 1.28 1.12 0.017 Comparative Example 1 0.26 19.2 25.336 -10.49 1.36 0.37 0.015 Comparative Example 2 0.49 28.5 40.064 -12.09 3.52 2.02 0.017 Comparative Example 3 1.35 18.3 50.16 26.93 2.15 0.02 0.04 0.018 Comparative Example 4 0.55 27.3 40.28 -8.88 2.15 0.015 Comparative Example 5 0.42 13.5 23.412 0.57 0.009 Comparative Example 6 0.58 27.6 41.29 -8.17 0.51 5.23 2.11 0.016 Comparative Example 7 0.55 27.0 39.98 -8.58 2.11 0.018

division Slavic heating temperature (℃) Hot rolling finishing temperature (℃) Solution treatment temperature (℃) Solvent treatment time (hours) Aging temperature (℃) Statute of limitations processing time (hours) Steel thickness (mm) Example 1 1152 886 1148 0.52 763 0.55 12 Example 2 1148 905 1162 0.74 724 0.82 20 Example 3 1175 915 1150 0.69 812 0.77 15 Example 4 1180 922 1155 1.02 720 0.66 18 Example 5 1185 859 1154 0.87 749 1.02 22 Comparative Example 1 1182 895 1152 0.63 748 0.56 12 Comparative Example 2 1147 889 1175 0.68 864 1.32 25 Comparative Example 3 1205 910 1182 1.13 837 0.88 30 Comparative Example 4 1149 894 1153 0.58 950 6.7 35 Comparative Example 5 1184 908 1163 0.76 738 1.42 10 Comparative Example 6 1155 893 1165 0.99 765 1.03 24 Comparative Example 7 1168 892 - - - - 12

division Microstructure (area%) Number of VC or VCN fine precipitates (units/mm ² ) Area fraction of grain boundary carbide (area%) Area fraction (area%) of grain boundary carbides in the weld heat affected zone Example 1 95% or more γ 892 less than 1% less than 1% Example 2 95% or more γ 1012 less than 1% less than 1% Example 3 95% or more γ 1109 less than 1% less than 1% Example 4 95% or more γ 1505 less than 1% less than 1% Example 5 95% or more γ 548 less than 1% less than 1% Comparative Example 1 95% or more γ less than 1% less than 1% Comparative Example 2 95% or more γ 85 3% or less 3% or less Comparative Example 3 94%γ Over 5% Over 5% Comparative Example 4 95% or more γ less than 1% less than 1% Comparative Example 5 84%γ less than 1% less than 1% Comparative Example 6 93%γ 620 Over 5% Over 5% Comparative Example 7 95% or more γ 30 less than 1% less than 1%

division Relative investment rate Yield strength at room temperature (MPa) Charpy impact energy (-84 ^o C, J) CLR(%) Example 1 1.002 762 69 0 Example 2 1.002 911 72 0 Example 3 1.002 915 68 0 Example 4 1.002 1117 58 0 Example 5 1.002 778 89 0 Comparative Example 1 1.52 342 15 0 Comparative Example 2 1.005 489 12 0 Comparative Example 3 1.003 546 5 0 Comparative Example 4 1.002 423 113 0 Comparative Example 5 3.6 349 2 36 Comparative Example 6 1.004 986 15 0 Comparative Example 7 1.002 488 112 0

Looking at Tables 1 to 3 above, Comparative Examples 1, 3 and 5, to 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.

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

In the case of Comparative Example 6, since Mo exceeded the range suggested by the present invention, carbides exceeded 5 area% at the austenite grain boundaries. As a result, excellent toughness could not be secured.

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

On the other hand, in the case of Examples 1 to 5 satisfying 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 (15)

Contains, in wt%, C: 0.050% to 1.7%, Mn: 15% to 40%, Cr: 3.0% or less, V: 1.0% to 3.0%, N: 1.000% or less (excluding 0%), Mo: 3.5% or less, and Nb: 1.0% 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,
Steel containing at least one type of fine precipitate of VC or VCN with a diameter of 50 nm or less in a number per unit area of 100/ ^mm2 or more.
In the first paragraph,
Steel additionally containing at least one of Ti: 1.0% or less, Al: 5.0% or less, and Si: 5.0% or less.
In the first paragraph,
Steel satisfying the following equation 1.

[Equation 1] 23.6[C]+[Mn] ≥ 28, 33.5[C]-[Mn] ≤ 23
(In the above formula 1, [C] and [Mn] represent the weight% of C and Mn contained in the steel, respectively.)
In the first paragraph,
The steel material wherein the above grain boundary carbide comprises at least one of Cr, Mo and Nb carbides.
In the first paragraph,
Steel having an area fraction of the above austenite of 95 area% or more.
In the first paragraph,
Steel with a yield strength at room temperature of 550 MPa or more.
In the first paragraph,
Steel having a Charpy impact energy value of 27 J or more at -84℃.
In the first paragraph,
Steel having an investment ratio of 1.2 or less after undergoing 20% cold plastic deformation at room temperature.
In the first paragraph,
Steel having an investment ratio of 1.1 or less when the strain during cold plastic deformation at room temperature is at least 2%.
In the first paragraph,
Steel having an area fraction of grain boundary carbides of 5.0 area% or less in the weld heat affected zone after submerged welding with a heat input of 3.0 kJ/ ^mm2 .
In the first paragraph,
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 represents the length of a single crack (㎛), and w represents the width of the specimen (㎛).)
A step of heating a slab comprising, in weight%, C: 0.050% to 1.7%, Mn: 15 to 40%, Cr: 3.0% or less, V: 1.0 to 3.0%, N: 1.000% or less (excluding 0%), Mo: 3.5% or less, and Nb: 1.0% or less, with the remainder being iron (Fe) and unavoidable impurities;
A step of obtaining a hot-rolled steel sheet by final 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.
In Article 12,
A method for manufacturing a steel material, wherein the above slab additionally contains at least one of Ti: 1.0% or less, Al: 5.0% or less, and Si: 5.0% or less.
In Article 12,
The above slab is a method for manufacturing steel satisfying the following equation 1.

[Equation 1] 23.6[C]+[Mn] ≥ 28, 33.5[C]-[Mn] ≤ 23
(In the above formula 1, [C] and [Mn] represent the weight% of C and Mn contained in the steel, respectively.)
In Article 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℃ or higher and 1200℃ 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.