JP6484716B2 - Lean duplex stainless steel and manufacturing method thereof - Google Patents

Description

  The present invention relates to a lean duplex stainless steel having a ferrite and austenitic structure and a method for producing the same.

In general, austenitic stainless steel with good workability and corrosion resistance contains iron (Fe) as a base metal, chromium (Cr), nickel (Ni) as main raw materials, molybdenum (Mo) and copper (Cu). It has been developed in various steel types to suit various applications by adding other elements.
304 series stainless steel with excellent corrosion resistance and workability includes expensive raw materials such as Ni and Mo. As an alternative to this, 200 series and 400 series stainless steels have been studied. Each has the disadvantage that formability and corrosion resistance do not reach 300 series stainless steel.
On the other hand, since a duplex stainless steel in which an austenite phase and a ferrite phase are mixed has all the advantages of an austenite and a ferrite, various types of duplex stainless steels have been developed.

Recently, in order to compensate for the disadvantages of price competitiveness, expensive alloy elements such as Ni and Mo contained in the duplex stainless steel are eliminated, and low-cost alloy elements are added. There is a growing interest in lean duplex stainless steels with even greater increases.
However, such a lean duplex stainless steel has a disadvantage in that hot workability is fragile due to a difference in strength between the ferrite and austenite phases, and a large number of surface cracks and edge cracks are generated.
In addition, when an alloy element whose components are adjusted for stainless steel is processed by a conventional general continuous casting method, when solidifying from a liquid phase to a solid phase, pores (porosity) are formed inside the slab due to a difference in nitrogen solid solubility. ) Occur frequently. Such internal pores can cause many defects on the product surface during the subsequent reheating and hot rolling processes, as well as a large number of cracks at the edges of the hot rolled coil. Bring the disadvantage of letting.
The matters described as the above background art are for the purpose of deepening the understanding of the background of the present invention, and are accepted as applicable to the prior art already known to those having ordinary knowledge in the technical field. Must not.

The present invention can reduce the cost by adjusting the content of the alloy component, while ensuring excellent stretch ratio and corrosion resistance by controlling so as to satisfy the lamination failure energy value existing in the lean duplex stainless steel. A lean duplex stainless steel and a method for producing the same are provided.
Further, by controlling so as to satisfy the value of critical deformation (critical strain induced martensite formation) in which a firing-induced martensite phase is formed, a lean duplex stainless steel that can ensure excellent stretch ratio and corrosion resistance and its A manufacturing method is provided.
Further, while solidifying from a liquid to a solid phase during casting, a lean duplex stainless steel that can solve the problem of a large release of nitrogen gas due to a rapid decrease in nitrogen solid solubility and its manufacture Provide a method.

The lean duplex stainless steel according to an embodiment of the present invention is a ferritic-austenitic stainless steel, and has an austenitic phase lamination failure energy (SFE) value of 19 to 20 represented by the following [Formula 2]. 37, the range of the critical deformation value in which the firing-induced martensite phase is formed is 0.1 to 0.25.
SFE = 25.7 + 1.59 × Ni / [K (Ni) -K (Ni) × V (γ) + V (γ)]) + 0.795 × Cu / [K (Cu) -K (Cu) × V (γ) + V (γ)])-0.85 × Cr / [K (Cr) -K (Cr) × V (γ) + V (γ)] + 0.001 × (Cr / [K (Cr) -K (Cr) × V (γ) + V (γ)]) 2 + 38.2 × (N / [K (N) -K (N) × V (γ) + V (γ)]) 0.5-2.8 × Si / [K (Si ) -K (Si) × V (γ) + V (γ)]-1.34 × Mn / [K (Mn) -K (Mn) × V (γ) + V (γ)] + 0.06 × (Mn / [ K (Mn) -K (Mn) × V (γ) + V (γ)]) 2 ……………… [Formula 2]
In [Formula 2], Ni, Cu, Cr, N, Si, and Mn mean the content percentage (wt%) with respect to the total weight of each component element, and K (x) is the distribution coefficient of each component element (x) as follows: V (γ) is the austenite phase fraction (0.45 to 0.75 range).
K (x) = [content of x element in the ferrite phase] / [content of x element in the austenite phase] ............ [Equation 3]

At this time, in K (x), K (Cr) = 1.16, K (Ni) = 0.57, K (Mn) = 0.73, K (Cu) = 0.64, (N) and K (Si) can have the following values depending on the contents of N and Si (wt%).
N: In the case of 0.2 to 0.32%, K (N) = 0.15,
If N <0.2%, K (N) = 0.25,
When Si ≦ 1.5%, K (Si) = 2.76−0.96 × Si,
When Si> 1.5%, K (Si) = 1.4
The stretch ratio of stainless steel is preferably 45% or more.

Stainless steel is, by weight, C: 0.08% or less (excluding 0%), Si: 0.2-3.0%, Mn: 2-4%, Cr: 18-24%, Ni: 0.2% to 2.5%, N: 0.15 to 0.32%, Cu: 0.2 to 2.5%, and the balance preferably contains Fe and other inevitable impurities.
Stainless steel may further include at least one or more of W: 0.1 to 1.0% and Mo: 0.1 to 1.0% by weight.
The stainless steel further includes at least one of Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05%, and V: 0.001 to 0.15% by weight. be able to.

Meanwhile, a method for producing a lean duplex stainless steel according to an embodiment of the present invention is a method for producing a ferrite-austenite lean duplex stainless steel, the process of preparing molten steel, and the following formula: 2] The laminar fracture energy (SFE) value of the austenite phase represented by 2] is 19 to 37, and the critical deformation value range in which the firing-induced martensite phase is formed is 0.1 to 0.25. A process of treating with stainless steel may be included.
SFE = 25.7 + 1.59 × Ni / [K (Ni) -K (Ni) × V (γ) + V (γ)]) + 0.795 × Cu / [K (Cu) -K (Cu) × V (γ) + V (γ)])-0.85 × Cr / [K (Cr) -K (Cr) × V (γ) + V (γ)] + 0.001 × (Cr / [K (Cr) -K (Cr) × V (γ) + V (γ)]) 2 + 38.2 × (N / [K (N) -K (N) × V (γ) + V (γ)]) 0.5-2.8 × Si / [K (Si ) -K (Si) × V (γ) + V (γ)]-1.34 × Mn / [K (Mn) -K (Mn) × V (γ) + V (γ)] + 0.06 × (Mn / [ K (Mn) -K (Mn) × V (γ) + V (γ)]) 2 ……………… [Formula 2]
In [Formula 2], Ni, Cu, Cr, N, Si, and Mn mean the content percentage (wt%) with respect to the total weight of each component element, and K (x) is the distribution coefficient of each component element (x) as follows: It is represented by [Formula 3], and V (γ) is an austenite phase fraction (range of 0.45 to 0.75).
K (x) = [content of x element in ferrite phase] / [content of x element in austenite phase] ............ [Equation 3]

In particular, the process of treating the molten steel with stainless steel includes a stage in which the molten steel is temporarily stored in a turn dish while maintaining the molten steel at a temperature 10-50 ° C. higher than the theoretical solidification temperature, and the molten steel is injected with a mold in the turn dish at 500-1500 ° C. / including a step of performing primary cooling by passing the mold while maintaining a cooling rate of min, and a step of performing secondary cooling while drawing and passing the molten steel on which the solidified shell is formed by primary cooling through the segment. Features.
At this time, it is preferable to inject 0.25 to 0.35 L / kg of cooling water as molten steel in which a solidified shell is formed in the stage of secondary cooling.

Further, when the surface temperature of the slab drawn after the stage of secondary cooling is in the range of 1100 to 1200 ° C., 100 to 125 L / kg · min of cooling water is applied to the surface of the slab at a ratio of air to cooling water (air / The method further includes a step of mixing and spraying so that (cooling water) becomes 1.0 to 1.2 and performing third cooling.
On the other hand, the process of treating the molten steel with stainless steel includes a step of producing a strip by solidifying the molten steel while passing between a pair of casting rolls, and the nitrogen solids contained in the molten steel in the step of producing the strip. Nitrogen above the solubility limit can be discharged out of the solidified shell through a casting roll.

At this time, in the stage of manufacturing the strip, it is preferable to use a casting roll having a gas discharge channel formed on the outer peripheral surface in the circumferential direction as at least one of the pair of casting rolls.
In addition, a plurality of gas discharge channels formed in a casting roll used in the step of manufacturing a strip are formed with a width of 50 to 500 μm and a depth of 50 to 300 μm, and the interval between adjacent gas discharge channels is 100 to 100. It is preferable that unevenness of 15 to 25 μm is formed on the surface of the casting roll at 1000 μm.

On the other hand, in the process of preparing the molten steel, the molten steel is in% by weight, C: 0.08% or less (however, 0% is excluded), Si: 0.2-3.0%, Mn: 2-4%, Cr: 18-24%, Ni: 0.2% -2.5%, N: 0.15-0.32%, Cu: 0.2-2.5%, the balance being Fe and other inevitable impurities It is preferable to contain.
And in the process of preparing molten steel, molten steel is weight% and may further contain at least 1 or more types among W: 0.1-1.0% and Mo: 0.1-1.0%.
Further, in the process of preparing the molten steel, the molten steel is in% by weight, among Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05% and V: 0.001 to 0.15%, At least one kind or more can be further included.

According to the embodiment of the present invention, resource saving and raw material cost can be significantly reduced by adjusting the contents of Ni, Si and Cu alloy components which are expensive elements. In particular, it can be sufficiently used as an alternative application of the 200 series and 300 series (STS 304 and 316) used for molding by ensuring the corrosion resistance equal to or higher than that of the STS 304 and the excellent stretch ratio.
And when the alloy element by the Example of this invention is continuously cast, the pinhole which generate | occur | produces inside a slab can be suppressed by controlling molten steel temperature and a cooling rate.
Furthermore, when strip casting the alloy element according to the embodiment of the present invention, by improving the casting roll, discharge of nitrogen gas generated when solidifying from the liquid phase to the solid phase is smoothly induced to generate internal pores. In addition, the occurrence of surface defects can be prevented.

The drawing which shows the critical deformation value in which a baking induction martensite phase is formed in the stress-deformation curve of the invention steel by the Example of this invention, and a comparison steel. It is the photograph which shows the presence or absence of the formation of a firing induction martensite phase in the structure | tissue photograph of the typical transmission electron microscope of the comparative steel by an Example of this invention, and invention steel, (a) is comparative steel, (b) is invention steel. is there. The graph showing the relationship between a draw ratio and the critical deformation value in which a baking induction martensite phase is formed. BRIEF DESCRIPTION OF THE DRAWINGS FIG. 1 is a schematic view showing a continuous casting method manufacturing process of a lean duplex stainless steel according to an embodiment of the present invention. 1 is a schematic view illustrating a manufacturing process for a strip casting method of a lean duplex stainless steel according to an embodiment of the present invention. The schematic diagram of the nitrogen exhaust channel formed in the casting roll of this invention. It is a structure photograph of the comparative material I and the invention material A, (a) shows the comparison material I, (b) shows the invention material A. The surface defect photograph of the comparative material H. The surface defect photograph of the comparative material F.

Hereinafter, embodiments of the present invention will be described in more detail with reference to the accompanying drawings. However, the present invention is not limited to the embodiments disclosed below, and may be embodied in various different forms. This embodiment is provided in order to complete the disclosure of the present invention and to enable those having ordinary knowledge to fully understand the scope of the present invention.
The present invention relates to a lean duplex stainless steel having a ferrite-austenite structure, and the ferrite-austenite structure referred to in the present invention means that the ferrite phase and austenite phase occupy most of the structure. It does not mean that steel is formed solely of ferrite and austenite phases. For example, the fact that the ferrite phase and the austenite phase occupy most of the structure means that the sum of the ferrite phase and the austenite phase accounts for 90% or more in the structure forming the stainless steel. The remaining part can be occupied by the martensite phase in which the austenite phase is transformed.

FIG. 1 is a drawing showing critical deformation values at which a firing-induced martensite phase is formed by stress-deformation curves of inventive steels and comparative steels according to examples of the present invention.
In the present invention, the content is% by weight (hereinafter, unless otherwise specified, the content of the component is% by weight), C: 0.08% or less (however, 0% is excluded), Si: 0.2-3.0% , Mn: 2-4%, Cr: 18-24%, Ni: 0.2% -2.5%, N: 0.15-0.32%, Cu: 0.2-2.5%, balance Is a lean duplex stainless steel having a two-phase structure formed of a ferrite phase and an austenite phase and containing Fe and other inevitable impurities. Lean duplex stainless steel is W: 0.1-1.0%, Mo: 0.1-1.0%, Ti: 0.001-0.1%, Nb: 0.001-0.05 %, V: 0.001 to 0.15% can be further included.

C is an element forming an austenite phase, and is an element effective for increasing material strength by solid solution strengthening. However, if added excessively, Cr effective for corrosion resistance at the ferrite-austenite phase boundary easily binds to the carbide-forming elements, so the Cr content around the grain boundary decreases and resistance to corrosion resistance. In order to maximize the corrosion resistance, it is desirable to add C in the range of more than 0 to 0.08% or less.
Si is partly added for the deoxidation effect, and is an element that forms a ferrite phase and is concentrated in the ferrite phase during annealing heat treatment. Therefore, 0.2% or more must be added in order to ensure an appropriate ferrite phase fraction. However, excessive addition exceeding 3.0% abruptly increases the hardness of the ferrite phase and affects the decrease in the draw ratio of the duplex stainless steel, and further secures the austenite phase to ensure a sufficient draw ratio. Make it difficult. Moreover, excessive addition reduces the fluidity | liquidity of slag at the time of steelmaking, forms the inclusion couple | bonded with oxygen, and reduces corrosion resistance. Therefore, it is desirable to limit the Si content to 0.2 to 3.0%.

Mn is an element that increases the deoxidizer and nitrogen solid solubility, and may be used in place of expensive Ni as an austenite phase forming element. When a large amount of Mn is added, the solid solubility of nitrogen is effective, but the corrosion resistance is deteriorated by forming MnS by combining with S in steel. Therefore, when the content exceeds 4%, it becomes difficult to ensure the corrosion resistance of 304 steel level. On the other hand, when the content of Mn is less than 2%, it is difficult to secure an appropriate austenite phase fraction even if Ni, Cu, N, etc., which are austenite phase forming elements, are adjusted. Solubility becomes low and sufficient solid solution of nitrogen cannot be obtained at normal pressure. Therefore, it is desirable to limit the Mn content to 2 to 4%.
Cr not only plays a major role in securing the ferrite phase of the duplex stainless steel as a stabilizing element of the ferrite phase together with Si, but is an essential element for ensuring corrosion resistance. Increasing the content increases the corrosion resistance, but in order to maintain the phase fraction, the content of expensive Ni and other austenite phase forming elements must be increased. Accordingly, it is desirable to limit the Cr content to 18 to 24% in order to ensure the corrosion resistance equal to or higher than that of 304 steel while maintaining the phase fraction of the duplex stainless steel.

Ni, together with Mn, Cu and N, is a stabilizing element for the austenite phase and plays a major role in securing the austenite phase of the duplex stainless steel. Instead of maximally reducing the expensive Ni content in order to save costs, Mn and N, which are other austenite phase forming elements, can be increased to sufficiently maintain the phase fraction balance by reducing Ni. However, in order to suppress the formation of the firing-induced martensite phase generated during cold working, 0.2% or more must be added in order to ensure sufficient austenite phase stability. On the other hand, if a large amount of Ni is added, the austenite phase fraction increases and it is difficult to secure an appropriate austenite phase fraction, and it is difficult to secure competitiveness against 304 steel due to an increase in product manufacturing cost due to expensive Ni in particular. . Therefore, it is desirable to limit the Ni content to 0.2 to 2.5%.
N is an element that contributes greatly to the stabilization of the austenite phase together with Ni in the duplex stainless steel, and is an element that has a high diffusion rate in the solid phase during annealing heat treatment and that is almost concentrated in the austenite phase. One. Therefore, an increase in N content can induce an increase in corrosion resistance and an increase in strength. However, the solid solubility of N varies depending on the added Mn content. If the N content exceeds 0.32% in the Mn range of the present invention, the surface defects are induced by the occurrence of blow holes, pin holes, etc. during casting due to an increase in nitrogen solid solubility, and the steel is induced It becomes difficult to produce a stable product. On the other hand, in order to ensure the corrosion resistance of 304 steel level, it is necessary to add 0.15% or more of N. If the N content is too low, it is difficult to ensure an appropriate phase fraction. Therefore, it is desirable to limit the N content to 0.15 to 0.32%.

Cu is an austenite phase stabilizing element together with Mn, Ni and N, and it is desirable to reduce the Cu content which plays the same role as Ni in order to reduce cost. However, it is desirable to add 0.2% or more in order to suppress the formation of excessive firing-induced martensite generated during cold working and ensure sufficient austenite phase stability. On the other hand, if the Cu content exceeds 2.5%, product processing becomes difficult due to hot brittleness, so it is desirable to adjust the Cu content to 0.2 to 2.5%.
W and Mo are elements that form an austenite phase and improve corrosion resistance, and elements that cause deterioration of corrosion resistance and mechanical properties by promoting the formation of intermetallic compounds at 700 to 1000 ° C. during heat treatment. But there is. When the content of W and Mo exceeds 1%, the formation of intermetallic compounds causes corrosion resistance, and in particular, a drastic decrease in the draw ratio. Moreover, in order to show the improvement effect of corrosion resistance, it is good to add 0.1% or more. Therefore, the contents of W and Mo are desirably limited to 0.1 to 1.0%, and at least one or more of W and Mo are preferably contained.

  Ti, Nb, and V are elements that react with nitrogen to form nitrides, which are crystallized in the molten steel with TiN, NbN, VN, and the like, and act as a nucleation field for the ferrite phase during solidification. Even if the cooling rate is increased, solidification is sufficiently progressed to suppress slab breakage. In addition, these elements are sufficiently dissolved in the manufacturing process, that is, during reheating and hot rolling, and react with carbon and nitrogen during cooling to form carbonitrides, thereby suppressing the formation of Cr carbides. Helps improve corrosion resistance. In particular, it is an element that suppresses the formation of Cr carbide in the heat affected zone when welding. When these elements are added excessively, that is, when Ti exceeds 0.1%, Nb exceeds 0.05%, and V exceeds 0.15%, these crystals form large clusters during solidification. As a result, the clogging phenomenon of the casting nozzle is caused, and when these are present in the surface layer portion of the slab, they act as a cause of defects during rolling and breakage during processing. Also, most of them are expensive alloy elements, and when they are added in a large amount, the manufacturing cost increases. Therefore, it is preferable to limit to the range of Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05%, V: 0.001 to 0.15%, and Ti, Nb and It is preferable that at least one of V is contained.

On the other hand, the present invention maintains the stretch ratio and the corrosion resistance by controlling the lamination defect energy by adjusting the content of alloy elements, the distribution coefficient and the phase fraction.
For example, the following [Equation 1] is an equation for deriving the stacking failure energy (SFE) by utilizing the content of all components in the alloy.
SFE = 25.7 + 1.59 (Ni + 0.5Cu) -0.85Cr + 0.001 × Cr2 + 38.2N0.5-2.8Si-1.34Mn + 0.06Mn2 ---- [Formula 1]
Cr, Ni, Cu, Si, Mn, and N in [Formula 1] mean the total content (wt%) of each component element.
However, the present applicant measured and calculated the stacking failure energy of the steel of the present invention by various methods, and as a result, calculated the stacking failure energy using only the component content of the alloy composition as shown in [Formula 1]. Rather than doing so, it has been found that calculation utilizing the component content of the austenite structure can be approached more accurately in order to predict the physical properties of the alloy. For this reason, calculating the stacking fault energy considering the inter-structure distribution coefficient of the alloy element was actually measured rather than calculating the stacking fault energy using only the component content of the entire alloy composition. It was found that the stacking fault energy value can be approximated by the stacking fault energy value.

For this reason, the present applicant has supplemented [Equation 1] with the following [Equation 2] so that the stacking failure energy of the austenite phase can be inferred using the partition coefficient of the austenite phase.
SFE = 25.7 + 1.59 × Ni / [K (Ni) -K (Ni) × V (γ) + V (γ)]) + 0.795 × Cu / [K (Cu) -K (Cu) × V (γ) + V (γ)])-0.85 × Cr / [K (Cr) -K (Cr) × V (γ) + V (γ)] + 0.001 × (Cr / [K (Cr) -K (Cr) × V (γ) + V (γ)]) 2 + 38.2 × (N / [K (N) -K (N) × V (γ) + V (γ)]) 0.5-2.8 × Si / [K (Si ) -K (Si) × V (γ) + V (γ)]-1.34 × Mn / [K (Mn) -K (Mn) × V (γ) + V (γ)] + 0.06 × (Mn / [ K (Mn) -K (Mn) × V (γ) + V (γ)]) 2 ……………… [Formula 2]
In [Formula 2], Ni, Cu, Cr, N, Si, and Mn mean the total content (wt%) of each component element.
And in [Formula 2], K (x) is expressed by the following [Formula 3] as a distribution coefficient of each component element (x).
K (x) = [content of x element in ferrite phase] / [content of x element in austenite phase] ............ [Equation 3]

The present applicant measured the distribution coefficient for each alloy element in various annealing conditions and alloy systems by utilizing Fe-EPMA and FE-TEM, which are more accurate than EDAX analysis using a normal scanning electron microscope. At this time, it was confirmed that the measured distribution coefficient for almost all alloy elements was 900 to 1200 ° C., which is a temperature range of hot rolling annealing or cold rolling annealing, and that the value did not change due to temperature change.
That is, in K (x), K (Cr) = 1.16, K (Ni) = 0.57, K (Mn) = 0.73, K (Cu) = 0.64, and K (N) It was confirmed that K and Si (Si) change depending on the contents of N and Si (wt%). However, when N: 0.2 to 0.32%, K (N) = 0.15, and when N <0.2%, K (N) = 0.25, and Si ≦ 1.5% In this case, K (Si) = 2.76−0.96 × Si, and when Si> 1.5, K (Si) = 1.4. At this time, the alloy elements N and Si mean components of the entire stainless steel. However, in this embodiment, N is 0.15 to 0.32%, and therefore K (N) = 0.25 is applied when N is 0.15% or more and less than 0.2%, and N: In the case of 0.2 to 0.32, K (N) = 0.15 is applied, and Si: 0.2 to 3.0%. Therefore, in the case of Si: 0.2 to 1.5%, K (Si) = 2.76−0.96 × Si is applied, and when Si is more than 1.5% and not more than 3.0%, K (Si) = 1.4 is applied.

In [Expression 2], V (γ) is an austenite phase fraction, and the austenite phase fraction is defined by the following relational expression.
1 = V (α) + V (γ)
Here, V (α) is the ferrite phase fraction, and V (γ) has a value in the range of 0.45 to 0.75.

On the other hand, the reason why the stacking failure energy value of the austenite phase is limited to 19 to 37 will be described.
It is known that the stacking failure energy of the austenite phase controls the deformation mechanism of the austenite phase. Usually, the stacking failure energy of the austenite phase indicates the degree to which the firing deformation energy added from the outside contributes to the deformation of the austenite phase when it is a single-phase austenitic stainless steel. In general, the lower the lamination failure energy, the higher the degree of formation of a firing-induced martensite phase that contributes to work hardening of the steel after the epsilon martensite phase is formed in the austenite phase. When the lamination defect energy is intermediate, mechanical twins are formed in the austenite phase. In the case of an intermediate stacking failure energy, a firing-induced martensite phase is formed at the intersection of these twins, and the added firing deformation energy causes a phase change mechanically and transitions from the austenite phase to the martensite phase. Therefore, in the case of stainless steel, it is known that a firing-induced martensite phase is formed in a fairly wide range except for the difference of the intermediate phase (epsilon martensite phase or mechanical twin). Therefore, when the stacking defect energy is less than 50 mJ / m ² , the firing induced martensite phase is formed after the epsilon martensite phase is formed in the austenite phase, or the mechanical twin is formed in the austenite phase. Later, a firing-induced martensite phase is formed.
However, when the stacking failure energy is 50 mJ / m ² or more, mechanical twinning or epsilon martensite phase is not formed, and deformation is promoted by dislocation migration, so the transition from the austenite phase to the martensite phase often occurs. It is known not to break.

On the other hand, when [Formula 1] is applied to infer the stacking failure energy only by the alloy component, the formation of the firing induced martensite phase is easy at 11 or less, and rapid work hardening at the initial stage of deformation, that is, the firing induced martensite. It was confirmed that a phase was formed resulting in a rapid decrease in draw ratio. It was confirmed that the formation behavior of the firing-induced martensite phase was changed in some component systems by changing the alloy elements distributed to the austenite phase by heat treatment and manufacturing process.
Here, the applicant corrected the calculation formula as shown in [Formula 2] in consideration of the distribution coefficient of the alloy elements distributed in the austenite phase after various manufacturing processes and heat treatments. As a result, when the calculated stacking failure energy of the austenite phase was less than 19, an epsilon martensite phase was formed first as an intermediate phase, and a martensite phase was formed at the intersection of the formed epsilon martensite phase. However, these martensite phases were formed abruptly in the early stage of deformation, and a phenomenon that the stretch ratio decreased due to rapid work hardening appeared. On the other hand, when the stacking failure energy of the austenite phase calculated using the corrected calculation formula exceeds 37, as a result of investigation using a transmission electron microscope, formation of a martensite phase was observed after firing deformation. Confirmed not to be. Therefore, it was found that the range of the stacking failure energy of the austenite phase is 19 to 37.

On the other hand, the lean duplex stainless steel according to the present invention is preferably formed of an austenitic phase with a volume fraction of 45 to 75% and a ferrite phase of 25 to 55%.
This is because, when the austenite phase fraction is less than 45%, an excessive concentration phenomenon of elements forming the austenite phase occurs in the austenite phase during annealing. As a result, the austenite phase is sufficiently stable and the transition of the firing-induced martensite phase that occurs during deformation is suppressed, the amount of austenite phase due to sufficient solid solution of the alloy elements is increased, and the tensile strength of the material is also sufficient. Can be secured. However, the phenomenon of decreasing ductility occurs and the desired stretch ratio and strength cannot be obtained sufficiently. Therefore, from the viewpoint of high ductility, it is desirable that the austenite phase fraction is 45% or more.

However, when the austenite phase fraction exceeds 75%, surface cracks and the like occur during hot rolling, resulting in a decrease in hot workability and loss of properties as a dual phase steel. Therefore, the austenite phase fraction is preferably 75% or less.
On the other hand, in the case of the present invention, it is preferable that the range of the critical deformation value in which the firing induced martensite phase is formed during cold working or tensile deformation is maintained at 0.1 to 0.25.
The critical deformation amount at which the firing-induced martensite phase is formed is measured from the inflection point of the stress-deformation curve, which usually contributes to work hardening of the martensite phase in steels where the firing-induced martensite phase is formed. The deformation value at the time is shown.

In addition, a method for obtaining a critical deformation value at which a firing-induced martensite phase is formed will be described as follows.
First, after a cold-rolled annealed material is sampled and processed in parallel with the rolling direction in accordance with the ASTM sub-size standard, it is used at room temperature (for example, 20 to 25 ° C.) using a tensile tester. A tensile test was performed until the material was broken at a strain rate of 0 × 10 ⁻³ / s. The change in the slope of the true deformation-true stress curve obtained at this time is the work hardening rate. The change in work hardening rate is closely related to the formation of firing-induced martensite phase. In the case of work hardening speed, after yield occurs, it gradually decreases as tensile deformation progresses, and when the firing induced martensite phase occurs and work hardening begins to contribute, that is, an inflection point is formed at critical deformation. The And when tensile deformation is advanced by deformation above the inflection point, and at the same time, the formation of the firing-induced martensite phase increases, the work hardening rate increases again.
Accordingly, the critical deformation value is a deformation value at a point where firing-induced martensite is formed and starts to contribute to work hardening, and means a deformation value at a point corresponding to the inflection point in the stress-strain curve obtained by a tensile experiment. Mathematically, it means a point where the value obtained by second-order differentiation of the curve becomes “0”.
For this reason, when the critical deformation value is less than 0.1, the firing-induced martensite phase is formed too easily during deformation, and the ductility of the material rapidly decreases due to rapid work hardening in the early stage of deformation. In addition, when the firing-induced martensite phase is formed extremely late, that is, when the critical deformation value exceeds 0.25, the stretching ratio decreases due to the occurrence of necking, which is local stress concentration due to insufficient work hardening of the material. Occurs. Thus, there is an appropriate range of work hardening rates. Therefore, in the present invention, it is desirable that the critical deformation value range in which the firing-induced martensite phase is formed is 0.1 to 0.25.

In addition, in the lean duplex stainless steel according to the present invention, it is very important to control the stability of the austenite phase. In general, the firing-induced martensite phase is a light phase formed when an unstable austenite phase is deformed, and induces work hardening and contributes to an increase in the drawing ratio of steel. In the case of the present invention steel of a duplex stainless steel comprising an austenite phase and a ferrite phase, the stability of the austenite phase can be adjusted by utilizing an appropriate distribution of alloy elements. In the present invention, a rapid solidification method is used as a method for appropriately distributing alloy elements. In the case of rapid solidification, since there is not enough time for diffusion to occur in the solid phase, the formed austenite phase and ferrite phase solidify in a non-equilibrium state. When these non-equilibrium solidified phases are subjected to hot rolling annealing heat treatment in a short time, the stability of the austenite phase can be controlled within a desired range by utilizing the distribution of the generated alloy elements. As a method for this purpose, the alloy was designed so that the content of nitrogen having a high diffusion rate in the solid phase was maintained higher than usual, and most of the nitrogen was segregated into the austenite phase.
〔Example〕

Hereinafter, the stretch ratio and the corrosion resistance will be described in detail through various examples of the lean duplex stainless steel according to the present invention.
As shown in Table 1 below, after preparing a test piece using molten steel with adjusted component content, hot rolling, hot rolling annealing, cold rolling annealing after cold rolling, The fraction was adjusted and the stretch ratio and corrosion resistance were measured.
The tensile test piece was measured by adjusting the tensile deformation rate at a speed of 1.0 × 10 ⁻³ / s at room temperature after processing a test piece of ASTM-sub size parallel to the rolling direction. The following [Table 1] shows the alloy composition (% by weight) with respect to the experimental steel types.

Table 2 below shows the phase fraction of the ferrite phase and austenite phase measured after annealing heat treatment of some experimental steel types in Table 1 at 1100 ° C.
The following [Table 3] shows the stacking failure energy value, distribution coefficient and phase fraction calculated by [Equation 1] without considering the distribution coefficient for the comparative steel and invention steel used in the explanation of the present invention. The results for the stacking failure energy value, the Gibbs Free energy difference, the presence or absence of the firing-induced martensite phase, and the stretching ratio calculated by [Equation 2] are shown.
At this time, the reason for utilizing the Gibbs Free energy difference is that the thermodynamic Gibbs freeness when the crystal structure of the phase having the same component is the austenite phase of FCC and the martensite phase of BCC. This is because the energy difference is calculated and the condition of ΔG = GM−Gγ ≦ 0 (Gibbs energy of martensite phase−Gibbs energy of austenite phase) is satisfied, and a firing-induced martensite phase is formed. Thus, although the Gibbs free energy difference and the formation of the firing induced martensite phase are closely related, for example, when the Gibbs free energy difference (ΔG) is a positive number, the firing induced martensite phase is formed. If the Gibbs free energy difference (ΔG) is a negative number, it means that a firing-induced martensite phase is formed.

In this example, in order to calculate the Gibbs free energies of the austenite phase and the martensite phase, calculation was performed using FACTsage 6.4 (Thermfact and GTT-Technologies) of common software. In particular, in order to calculate Gibbs free energy, it is necessary to know the components of the alloy present in the austenite phase among the steels present in the two phases of the ferrite phase and the austenite phase. The amount of the component can be calculated using the distribution coefficient and the phase fraction presented in the present invention. For example, X component in austenite phase = X / [K (X) −K (X) × V (γ) + V (γ)] (X: all X components, K (X): partition coefficient, V (γ ): Austenite phase fraction).
Moreover, the presence or absence of the formation of the firing-induced martensite phase was measured by utilizing a ferrite scope (usual product) in the crack extension section before necking occurred during tensile deformation.

In the case of lean duplex stainless steel, the phase fraction varies depending on the alloy components and the heat treatment temperature.
For this reason, [Table 2] shows the ferrite and austenite phase fractions when heat-treating Comparative Steel 1 to Comparative Steel 5 and Inventive Steel 1 to Inventive Steel 10 at 1100 ° C., respectively. It can be seen that in the case of Invention Steel 1 to Invention Steel 10, the ferrite phase fraction is within the range of about 25 to 55%, and the austenite phase fraction is within the range of 45 to 75%. When the comparative steel 2 is heat-treated at 1100 ° C., the ferrite phase fraction and the austenite phase fraction appear as 83% and 17%, respectively. That is, it can be seen that the comparative steel 2 is not included in the range of the ferrite and austenite phase fractions of the present invention.
And it turns out that the comparative steel 4 is not contained in the appropriate range of the stacking fault energy (SFE) value of an austenite phase with the stacking fault energy of an austenite phase of 17.88 mJ / m < ² > which considered distribution.

On the other hand, FIG. 1 is a typical nominal deformation-nominal stress comparison curve obtained in the present invention, and is a result of a tensile experiment after heat-treating each material at 1100 ° C.
In the case of the comparative steel 1, the firing induced martensite phase was not formed during the uniform deformation. As a result, since there was no firing-induced martensite phase capable of suppressing local necking due to work hardening during firing deformation, a decrease in the stretching ratio was predicted. Actually, in the case of the comparative steel 1, the draw ratio was about 31%, which was very inferior.
And in the case of the comparative steel 3, it can confirm that the critical deformation value in which a baking induction martensite phase is formed at the time of baking deformation is 0.1 or less (inflection point; indicated by an arrow). It was predicted that the draw ratio would decrease due to rapid work hardening due to the rapid formation of the induced martensite phase. Actually, the comparative steel 3 is a duplex stainless steel composed of a ferrite phase and an austenite phase, but it was very inferior at a stretch ratio of about 35%.

  On the other hand, in the case of the steel of the present invention, in the stress-deformation curve, when the range of critical deformation in which the firing induced martensite phase is formed is 0.1 to 0.25, the firing induced martensite phase formed during processing It was found that various values of the stretching ratio can be obtained by appropriately controlling the formation rate of. That is, when the invention steel 8 and the invention steel 1 are compared, it can be seen that the stretch ratio increases as the critical deformation in which the firing-induced martensite phase is formed increases. This is because the austenite phase is controlled to be transferred to the firing-induced martensite phase during cold working, and “almost all of the stretch ratio is 45% or more. This shows a stretch ratio much better than that of the comparative steel 1 which is an ordinary lean two-phase steel to be replaced with the steel of the present invention, and is comparable to the stretch ratio of the 304 steel to be replaced with the present invention. An excellent stretch rate was shown.

More specifically, when the critical deformation value is less than 0.1, the firing-induced martensite phase is rapidly formed, and the stretch rate is rapidly decreased by the hardening of the material by the rapid work hardening. When the critical deformation value exceeds 0.25, the firing-induced martensite phase is formed very slowly, and local necking of the material due to deformation cannot be suppressed. On the other hand, in the case of a lean duplex stainless steel composed of the austenite phase-ferrite phase of the present alloy system, if the critical deformation value range in which the firing induced martensite phase is formed is 0.1 to 0.25, the conventional A draw ratio of 45% or more, which is far superior to 30% or less, which is the draw ratio of the duplex stainless steel, can be secured, and the draw ratio is comparable to 304 steel under some deformation conditions. % Or more can be secured. From this, it is desirable that the critical deformation value at which the firing-induced martensite phase is formed during cold working is 0.1 to 0.25.
And as shown in the comparative steel 1 and the comparative steel 5 of [Table 3], even if the stacking fault energy by [Formula 2] exists in the range of 19-37, the Gibbs which a martensite phase is formed in an austenite phase The free energy value represents a positive number, and it can be confirmed that the formation of firing-induced martensite phase may not be observed in the microstructure observed with a transmission electron microscope. Was observed to decrease.
And as shown in the comparative steel 2 and the comparative steel 3 of [Table 3], the Gibbs free energy value represents a negative number, and the critical deformation value is 0.1 even if the firing induced martensite phase is formed. It was observed that the stretching ratio also decreased when not in the range of ~ 0.25.

(A) and (b) of FIG. 2 show the microstructures of the transmission electron microscopes of comparative steel 1 and inventive steel 1, respectively. In the case of the comparative steel 1, as shown in FIG. 2A, it can be seen that deformation bands and mechanical twins due to deformation are observed, but no firing-induced martensite phase is observed. In the case of the invention steel 1, as shown in FIG. 2 (b), it can be seen that the firing-induced martensite phase is formed at the intersection of deformation bands and mechanical twins (the firing-induced martensite phase is indicated by an arrow). ).
On the other hand, FIG. 3 is a graph showing the relationship between the stretch ratio and the critical deformation value at which the firing-induced martensite phase is formed, and is a result of measuring the critical deformation at which the martensite phase is formed based on the stress-deformation curve of FIG. is there.
As shown in FIG. 3, when the critical deformation value is less than 0.1 and when it exceeds 0.25, the stretch ratio of 45% is not ensured, while the range of the critical deformation value is 0.1 to 0.25. If it has, it can be confirmed that a stretching ratio of 45% or more can be secured.

On the other hand, the lean duplex stainless steel according to the present invention solves the problem of generation or discharge of nitrogen gas due to the difference in nitrogen solid solubility when solidified from the liquid phase to the solid phase. It becomes possible to manufacture by casting method.
First, a method for producing a lean duplex stainless steel by a continuous casting method will be described.
FIG. 4 is a schematic view illustrating a manufacturing process for a continuous casting system of lean duplex stainless steel according to an embodiment of the present invention.

The lean duplex stainless steel according to an embodiment of the present invention is manufactured by a normal continuous casting equipment 100 in which a raidle 110, a turn dish 120, a mold 130, and a plurality of segments 140 are sequentially arranged. However, the rear end portion of the segment 140 is further provided with an injection unit 150 that mixes and injects air and cooling water.
In order to manufacture lean duplex stainless steel by a continuous casting method, first, the molten steel having the alloy components presented above is prepared and moved to the raidle 110, and then temporarily attached to the turn dish 120 using the shroud nozzle 111. save. At this time, it is desirable that the molten steel temporarily stored in the turn dish 120 is maintained at a temperature 10 to 50 ° C. higher than the theoretical solidification temperature.

More specifically, ΔT (° C.) of the difference between the molten steel temperature and the theoretical solidification temperature in the turn dish 120 has a lower limit of 10 ° C. and an upper limit of 50 ° C. When ΔT becomes lower than the lower limit of 10 ° C., the turn dish 120 Solidification of the molten steel M begins and problems occur in continuous casting. On the other hand, when ΔT exceeds the upper limit of 50 ° C., the solidification rate during solidification slows down and the solidification equipment becomes coarse, causing solidification cracks and It becomes easy to generate a linear defect during hot rolling.
And the molten steel M is inject | poured into the casting_mold | template 130 using the immersion nozzle 121 with the turn dish 120. FIG. At this time, the mold 130 is primarily cooled by passing the mold 130 while maintaining the cooling rate of the molten steel M at 500 to 1500 ° C./min.
At this time, if the cooling rate is less than 500 ° C./min, the nitrogen pinhole that has become coarse due to the problem that the nitrogen gas generated due to the difference in nitrogen solubility due to delta ferrite solidification formed in the initial stage is discharged through the solidified shell of the mold 130. As a result, a large amount of nitrogen pinholes are generated in the continuously cast slab. Further, the delta ferrite formed in the initial stage is coarsened and becomes weak against external stress. In addition, when the cooling rate is less than 500 ° C./min, the amount of cooling in the mold 130 (primary cooling) and the amount of cooling in the segment 140 (secondary cooling) during continuous casting is reduced. The heat transfer of the piece S is delayed, the strength of the slab solidified layer is lowered, and a phenomenon that the slab bulges occurs to cause deterioration in operation and quality.

If the cooling rate is controlled to exceed 1500 ° C./min, it is very advantageous from the viewpoint of nitrogen pinholes, but the current continuous casting equipment is limited, and continuous operation is impossible. During continuous casting, the time for the segregation of the remaining solute elements to diffuse between dendrites is reduced and cracks on the surface of the slab are generated. Due to such a phenomenon, there is a problem that an overlapping phenomenon occurs in which the slab shell (outer shape, shell) is temporarily broken inside the mold 130. Therefore, it is preferable to set the cooling rate at the time of primary cooling in the mold 130 at 500 to 1500 ° C./min.
The molten steel M in which the solidified shell is formed in the mold 130, that is, the slab S, is drawn into the segment 140 and subjected to secondary cooling. At this time, the slab S is 0.25 to 0.35 L / kg. It is desirable to inject cooling water. The reason for limiting the specific water amount in the segment 140 in this way is as follows.

If the specific water amount in the segment 140, that is, the secondary cooling zone is set relatively large, the solidified structure can be formed finely. However, if the specific water amount exceeds 0.35 L / kg, the continuous casting process can be performed. Since the time for diffusion of the segregated impurities between the solidified structures is reduced, the sigma is present and cracks are generated on the surface of the slab. Further, since not only cracks due to thermal stress but also residual stress is excessively generated on the surface, surface cracks are generated during grinding of the slab. On the other hand, if the specific water amount is less than 0.25 L / kg, the solidified structure becomes excessive and solidification cracks occur due to the sigma shape generated at the grain boundaries, and the slab solidified shell (shell) during continuous casting. ) Is reduced, and cracking occurs due to bulging of the slab.
Accordingly, the specific water amount range in the segment 140 is desirably 0.25 to 0.35 L / kg.

Then, tertiary cooling is performed on the slab S that has been secondarily cooled while being pulled out by the segment 140. In the third cooling, the surface temperature of the slab S is in the range of 1100 to 1200 ° C. while continuously pulling out at the segment 140, and cooling water is cooled with air by 100 to 125 L / kg · min with respect to the entire surface of the slab S. It is cooling by mixing and spraying so that the ratio of water (air / cooling water) is 1.0 to 1.2.
The tertiary cooling is controlled so as to ensure a uniform scale on the surface of the slab S. The reason for this is that in the case of lean duplex stainless steel, the amount of oxidation in the heating furnace is very small, and the lubrication effect by the scale during hot rolling is small, so it is very difficult to reduce surface cracks. Therefore, in order to prevent temperature reduction due to contact between the roll and the steel plate during rolling, and to reduce surface friction by reducing the frictional force between the roll and the steel plate, a dense and thick scale is formed on the steel plate surface, It should be ensured that thinning does not occur easily during rolling. The reason for limiting the surface temperature of the slab S, the specific amount of cooling water, and the ratio of cooling water to air (air / cooling water) as described above is desired on the surface of the slab S when the above conditions are not satisfied. This is because the scale is not formed at a standard thickness (approximately 35 μm ± 2 μm), and the generated scale is not uniformly formed.

Hereinafter, the slab was changed while changing the molten steel temperature in the turn dish, the cooling rate in the mold, and the specific water amount in the secondary cooling zone as shown in [Table 4]. As a result, the degree of occurrence of cracks on the surface of the pinhole and slab is shown in [Table 4]. At this time, whether or not pinholes were generated in the slab was observed by grinding the surface of the slab after grinding the surface of the slab by about 0.5 mm.

As can be seen from [Table 4], invention materials A to E, which are invention examples that satisfy all of the control conditions of the present invention, are all continuous cast slabs and no pinholes due to nitrogen are generated, and swells. In addition, no defects occurred on the surface of the hot rolled coil.
And since the comparative material F and the comparative material G had the cooling rate in a casting_mold | template in the range of this invention, the pinhole by nitrogen did not generate | occur | produce in the inside of a slab. However, in the case of the comparative material F, the specific water amount was larger than the range of the present invention, and no bulging occurred during casting, but the thermal stress acted on the slab surface to induce the generation of cracks. Moreover, the range of the specific water amount of the secondary cooling zone of the comparative material G was less than the range of the present invention, and cracks occurred on the surface of the slab due to the occurrence of swelling on the slab. As a result, linear defects occurred on the surface of the hot rolled coil due to the formation of excessive scale locally during hot rolling.
And the comparative material I and the comparative material J had the cooling rate in a casting_mold | template slower than the range of this invention, and the severe pinhole generate | occur | produced in the slab. However, the surface of the continuous cast slab is good when the specific water amount of the secondary cooling zone is within the range of the present invention, but a large amount of linear notches are generated during hot rolling due to pinholes existing in the slab.

On the other hand, FIG. 7 is a structural photograph of the comparative material I and the inventive material A manufactured by the continuous casting method of the present invention, and FIG. 8 is a surface defect photograph of the comparative material H manufactured by the continuous casting method of the present invention. 9 is a surface defect photograph of the comparative material F manufactured by the continuous casting method of the present invention. At this time, FIGS. 8 and 9 are photographs of surface defects on the surface of the hot-rolled coil discovered after hot rolling of the comparative material H and the comparative material F. FIG.
As can be seen from FIG. 7, pinholes were not found on the surface of the slab of Invention Material A, but it can be confirmed that Comparative Material I generates a large amount of pinholes. Further, as can be seen from FIG. 8, when the surface of the hot rolled coil was observed after hot rolling the comparative material H with relatively good pinhole generation, there were many pinhole defects that were stretched in the rolling direction. Observed. In FIG. 9, when the surface of the hot rolled coil was observed after hot rolling the comparative material F, a large amount of slab cracking surface defects were observed.
Therefore, by appropriately controlling the cooling rate in the mold during continuous casting and the specific water amount in the secondary cooling zone according to the present invention through various embodiments, the occurrence of pinholes and cracks and bulges during casting was suppressed. It was confirmed that not only excellent slab quality of lean duplex stainless steel composed of austenite phase and ferrite phase was obtained, but also stable continuous casting operation was possible.

発明材１〜発明材４のように、鋳片表面温度が１０００〜１２００℃の地点で空気／冷却水の割合を１．０〜１．２と維持しながら、冷却数量１００〜１２０ Ｌ／ｋｇ・分で２０〜３０分間冷却水を噴射する場合、スケールが非常に均一で厚くなることが分かる。
しかし、比較材１及び比較材２のように、冷却数量を５０Ｌ／ｋｇ・分または８０Ｌ／ｋｇ・分に噴射した場合は、冷却水の量が十分ではないため酸化スケールの生成を促進させることができず、均一な酸化スケールも得ることができなかった。 And the quantity of cooling, the injection time, the ratio of air / cooling water, the surface temperature of the slab during the tertiary cooling of the lean duplex stainless steel having the composition according to the present invention and subjected to the primary cooling and the secondary cooling. As shown in [Table 5], slabs were produced, and as a result, scale thickness and uniformity were shown in [Table 5].

Cooling quantity 100-120 L / kg while maintaining the ratio of air / cooling water at 1.0-1.2 at the point where the slab surface temperature is 1000-1200 ° C. It can be seen that when the cooling water is sprayed for 20 to 30 minutes per minute, the scale becomes very uniform and thick.
However, when the cooling quantity is injected to 50 L / kg · min or 80 L / kg · min as in Comparative Material 1 and Comparative Material 2, the amount of cooling water is not sufficient, so that the generation of oxide scale is promoted. Thus, a uniform oxide scale could not be obtained.

Then, when the oxide scale was examined by changing the ratio of air to cooling water (air / cooling water) as in the comparative material 3 to the comparative material 5, the thickness of the scale layer was increased as the amount of air was sufficient. It was confirmed that it increased. Accordingly, in order to form a scale layer having a desired thickness, it is preferable to maintain the ratio of air to cooling water (air / cooling water) at 1.0 or more. However, when the ratio of air exceeds the upper limit 1.2, a sufficient scale layer can be obtained, but there is a possibility that the entire cooling water system may be hindered.
When the cooling water injection time is 15 minutes and 10 minutes as in the comparative material 6 and the comparative material 7, the temperature of the cooling water injection point and the ratio of the cooling water / air are close to the conditions of the invention material. However, a scale layer having a uniform and sufficient thickness could not be obtained due to lack of spraying time. Therefore, in order to obtain a uniform and thick scale layer, it was confirmed that sufficient injection time was required for the slab to react with air. However, if the injection time of the cooling water exceeds a certain time, the slab may stagnate and the production amount may decrease.

Further, when the cooling water is sprayed at the points where the surface temperature of the slab is 932 ° C. and 1063 ° C. as in the comparative material 8 and the comparative material 9, the thickness of the oxide scale is 15 μm and 26 μm, and the non-uniform oxide scale. showed that. However, it was predicted that the higher the temperature of the slab, the more the oxide scale was formed and the uniform scale could be obtained. Moreover, it can be predicted that a uniform scale layer can be obtained as the surface temperature of the cast slab increases with continuous casting.
Thus, after the continuous casting process is completed, the formation of scale can be optimized by injecting the optimal cooling water and air ratio and the optimal quantity of cooling water at the optimal injection position in the cooling injection process. , The surface quality can be improved. In addition, the cost for the scratch removal process can be minimized and the added value can be improved.

Next, a method for producing a lean duplex stainless steel by strip casting will be described.
FIG. 5 is a schematic view illustrating a manufacturing process of a strip casting method of a lean duplex stainless steel according to an embodiment of the present invention, and FIG. 6 is a schematic view of a nitrogen exhaust channel formed on a casting roll according to the present invention. is there.
The lean duplex stainless steel according to an embodiment of the present invention is manufactured by a normal strip casting equipment 200 in which a ladle 210, a turn dish 220, a pair of casting rolls 230, an in-line roller 260, and a take-up roll 270 are sequentially arranged. Is done. However, a gas discharge channel 231 is formed on the surface of the casting roll 230.
In order to manufacture lean duplex stainless steel by the strip casting method, first, the molten steel M having the above-described alloy components is prepared and moved to the cradle 210, and then turn dish using the shroud nozzle 211. 220 is temporarily stored. Then, the strip S is produced by solidification while passing between the pair of casting rolls 230 through the injection nozzle 221, and the produced strip S is rolled by an in-line roller 260 continuously arranged with the casting roll 230. Taken by the take-up roll 270.

On the other hand, a meniscus shield 250 is mounted on the upper part of the casting roll 230 in order to prevent the molten metal surface from being oxidized by contact with air, and an appropriate gas is injected into the meniscus shield 250 to appropriately prevent oxidation. An atmosphere is formed.
As described above, the molten steel M passes through the roll nip where the pair of casting rolls 230 come into contact, is rolled through the in-line roller 260, and then manufactured with a strip S of 10 mm or less through processes such as a heat treatment process and cold rolling.

One of the most important technical elements of the twin roll strip caster that directly manufactures the above-described strip S of 10 mm or less is an internal water-cooled casting roll (twin-drum) that rotates in the opposite direction through the injection nozzle 221 at a high speed. rolls 230) and the side dam 240, side dam, supplying molten steel M, releasing a large amount of heat through the surface of the water-cooled casting roll 230, rapidly cooling the molten steel, breaking the thin plate of the desired thickness Without increasing the actual yield.
The high ductility lean duplex stainless steel manufacturing method of the present invention solves the problem of nitrogen exceeding the solid solution limit contained in molten steel, which is the cause of edge cracks and surface cracks, and the problem of reduced hot workability due to nitrogen content. did.
That is, when the molten steel M is solidified through the casting roll 230, the rapid casting is completed while discharging the nitrogen exceeding the solid solution limit, and the rapid casting is completed by using the inline roller 260 continuously advanced after casting. The above-mentioned problems were solved by manufacturing a thin strip S of about 2 to 5 mm.

Various means may be provided for removing nitrogen above the solid solution limit contained in the molten steel M in the process of strip casting. In the method for producing a high ductility lean duplex stainless steel according to the present invention, As an example, by forming a nitrogen discharge channel on the surface of the casting roll 230, nitrogen exceeding the solid solution limit was discharged during solidification of the molten steel.
The problem of internal pores due to nitrogen occurs in the process of rapidly cooling while the molten steel M passes between a pair of casting rolls 230.
Therefore, nitrogen exceeding the solid solution limit in the molten steel M must be advanced at the same time as the molten steel M passes through the casting roll 230. For this reason, it is desirable to form a gas discharge channel 231 on the surface of the casting roll 230 so that nitrogen can be discharged during casting.

The gas discharge channel 231 is a fine channel that does not allow the molten steel M to pass through and discharges only nitrogen gas. The gas discharge channel 231 may be formed on the casting roll 230 in various ways, and may be formed on the surface of the casting roll 230 in the circumferential direction. Nitrogen gas can be guided and discharged.
The gas discharge channel 231 is a fine channel having a width of 50 to 500 μm and a depth of 50 to 300 μm, and a plurality of gas discharge channels 231 are formed in the circumferential direction of the casting roll 230. It is desirable to form the gap of about 100 to 1000 μm.
The shape, structure, and application position of the gas exhaust channel 231 can be variously modified as long as the function can be achieved.
On the other hand, when a plurality of such gas discharge channels 231 are formed, the contact area between the casting roll 230 and the molten steel M passing through the casting roll 230 may be reduced. It is preferable that irregularities protrude from the surface of the casting roll. Such irregularities have an average size of 15 to 25 μm.

Hereinafter, in order to confirm the influence of nitrogen above the solid solution limit on the thin plate in the molten steel, a lean duplex stainless steel was manufactured while changing the composition and casting method of the molten steel as shown in [Table 6] below. . At this time, Comparative Example 1 is obtained by casting a molten steel having a specific composition using a general continuous casting method, and Comparative Example 2 is a general strip casting (rapid casting) method using a molten steel having a specific composition. Examples 1 to 5 are cast in a strip casting process using a casting roll according to the present invention while discharging nitrogen above the solid solution limit in molten steel.

As shown in [Table 6], in Comparative Example 1, it was confirmed that nitrogen was not discharged during the continuous casting process and pores were generated inside the slab.
Further, in Comparative Example 2, it was confirmed that nitrogen was not discharged during the strip casting process of a general method, and pores were generated inside the strip.
This is due to the difference in nitrogen solubility that the molten steel generates while passing through the mold or casting roll and solidifying.
The nitrogen composition of the highly ductile lean duplex stainless steel of the present invention is in the range of 1500-3200 ppm.
On the other hand, the process of solidification of the molten steel from the liquid phase to the solid phase proceeds in the order of liquid phase-> liquid phase + delta->delta-> delta + austenite, but when the liquid phase changes to delta phase, The solid solubility is about 1164 ppm, and a solid solubility difference of about 836 to 1836 ppm occurs. For this reason, some of the nitrogen supersaturated in the liquid phase is gasified when it solidifies, forming various pores inside the solidified material, as well as a solidified shell formed on the surface of the material. In addition, a large number of pores are formed.
Thus, there are many pores inside the solidified material, and some of these pores are squeezed during hot rolling, but uncompressed pores become internal defects and heat the heating furnace. When it is exposed to the outside, it develops into various surface defect forms.
On the other hand, Examples 1 to 5 were strip casting processes according to the present invention, and it was confirmed that nitrogen was discharged during the process and no pores were generated inside the strip.

  As mentioned above, although this invention was demonstrated based on an accompanying drawing and a preferable Example, this invention is not limited to it, It is limited by the claim which is mentioned later. Therefore, a person having ordinary knowledge in the technical field can variously modify and modify the present invention without departing from the technical idea of the claims.

100: Continuous casting equipment 110: Ladle 111: Shroud nozzle 120: Turn dish 121: Immersion nozzle 130: Mold 140: Segment 150: Injection means 200: Strip casting equipment 210: Ladle 211: Shroud nozzle 220: Turn dish 221: Injection nozzle 230: Casting roll 231: Gas discharge channel 240: Side dam 250: Meniscus shield 260: Inline roller 270: Rolling roll M: Molten steel S: Cast slab, strip

Claims (13)

In mass%, C: 0.08% or less (however, 0% is excluded), Si: 0.2-3.0%, Mn: 2-4%, Cr: 18-24%, Ni: 0.00. 2% to 2.5%, N: 0.15 to 0.32%, Cu: 0.2 to 2.5%, the balance was formed of ferrite and austenite phases composed of Fe and other inevitable impurities A lean duplex stainless steel having a two-phase structure,
Critical deformation, which is a deformation value at a point where a lamination induced energy (SFE) value of an austenite phase represented by the following [Equation 2] is 19 to 37 and a work- induced martensite phase is formed and starts to contribute to work hardening A lean duplex stainless steel characterized by a value range of 0.1 to 0.25.
SFE = 25.7 + 1.59 × Ni / [K (Ni) −K (Ni) × V (γ) + V (γ)]) + 0.795 × Cu / [K (Cu) −K (Cu) × V ( γ) + V (γ)]) − 0.85 × Cr / [K (Cr) −K (Cr) × V (γ) + V (γ)] + 0.001 × (Cr / [K (Cr) −K ( Cr) × V (γ) + V (γ)]) 2 + 38.2 × (N / [K (N) −K (N) × V (γ) + V (γ)]) 0.5-2.8 × Si /[K(Si)−K(Si)×V(γ)+V(γ)]−1.34×Mn/[K(Mn)−K(Mn)×V(γ)+V(γ)]+0. 06 * (Mn / [K (Mn) -K (Mn) * V ([gamma]) + V ([gamma])]) 2 [Equation 2]
In [Formula 2], Ni, Cu, Cr, N, Si, and Mn mean the total content (wt%) of each component element, and K (x) is the following [Formula 2] as a distribution coefficient of each component element (x). 3], and V (γ) is an austenite phase fraction (0.45 to 0.75 range).
K (x) = [content of x element in the ferrite phase] / [content of x element in the austenite phase] ............ [Equation 3] Among the above K (x), K (Cr) = 1.16, K (Ni) = 0.57, K (Mn) = 0.73, K (Cu) = 0.64, K (N) 2. The lean duplex stainless steel according to claim 1, wherein K and Si (Si) have the following values depending on the contents of N and Si (wt%).
N: 0.2 to 0.32%, K (N) = 0.15,
If N <0.2%, K (N) = 0.25,
When Si ≦ 1.5%, K (Si) = 2.76−0.96 × Si,
When Si> 1.5%, K (Si) = 1.4   The lean duplex stainless steel according to claim 1 or 2, wherein the stretch ratio of the stainless steel is 45% or more. 2. The stainless steel according to claim 1 , further comprising at least one of W: 0.1 to 1.0% and Mo: 0.1 to 1.0% in mass%. Lean duplex stainless steel. The stainless steel may be at least one of Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05% and V: 0.001 to 0.15% by mass. The lean duplex stainless steel according to claim 1 , comprising: A method of producing a ferrite-austenite-based lean duplex stainless steel,
In mass%, C: 0.08% or less (however, 0% is excluded), Si: 0.2-3.0%, Mn: 2-4%, Cr: 18-24%, Ni: 0.00. 2% to 2.5%, N: 0.15 to 0.32%, Cu: 0.2 to 2.5%, the process of preparing a molten steel composed of Fe and other inevitable impurities ,
Critical deformation, which is a deformation value at a point where a lamination induced energy (SFE) value of an austenite phase represented by the following [Equation 2] is 19 to 37 and a work- induced martensite phase is formed and starts to contribute to work hardening Temporarily storing the molten steel in a turn dish while maintaining the molten steel 10-50 ° C. higher than the theoretical solidification temperature so that the value range is 0.1-0.25 ;
Injecting molten steel with the mold in the turn dish and maintaining the cooling rate of 500-1500 ° C./min to pass through the mold for primary cooling;
A lean duplex stainless steel comprising a step of treating the molten steel with stainless steel including a step of secondary cooling while drawing and passing the molten steel formed with a solidified shell by primary cooling. Manufacturing method.
SFE = 25.7 + 1.59 × Ni / [K (Ni) −K (Ni) × V (γ) + V (γ)]) + 0.795 × Cu / [K (Cu) −K (Cu) × V ( γ) + V (γ)]) − 0.85 × Cr / [K (Cr) −K (Cr) × V (γ) + V (γ)] + 0.001 × (Cr / [K (Cr) −K ( Cr) × V (γ) + V (γ)]) 2 + 38.2 × (N / [K (N) −K (N) × V (γ) + V (γ)]) 0.5-2.8 × Si /[K(Si)−K(Si)×V(γ)+V(γ)]−1.34×Mn/[K(Mn)−K(Mn)×V(γ)+V(γ)]+0. 06 * (Mn / [K (Mn) -K (Mn) * V ([gamma]) + V ([gamma])]) 2 [Equation 2]
In [Formula 2], Ni, Cu, Cr, N, Si, and Mn mean the total content (wt%) of each component element, and K (x) is the distribution coefficient of each component element (x) as shown below [Formula 3 V (γ) is the austenite phase fraction (range of 0.45 to 0.75).
K (x) = [content of x element in ferrite phase] / [content of x element in austenite phase] ............ [Equation 3] The method for producing a lean duplex stainless steel according to claim 6 , wherein 0.25 to 0.35 L / kg of cooling water is injected as molten steel in which a solidified shell is formed in the secondary cooling step. . When the surface temperature of the slab drawn after the secondary cooling is in the range of 1100 to 1200 ° C., 100 to 125 L / kg · min of cooling water is applied to the surface of the slab at a ratio of air to cooling water (air / cooling). The method for producing a lean duplex stainless steel according to claim 6 , further comprising a step of mixing and spraying so that water is 1.0 to 1.2 and then performing tertiary cooling. A method of producing a ferrite-austenite-based lean duplex stainless steel,
In mass%, C: 0.08% or less (however, 0% is excluded), Si: 0.2-3.0%, Mn: 2-4%, Cr: 18-24%, Ni: 0.00. 2% to 2.5%, N: 0.15 to 0.32%, Cu: 0.2 to 2.5%, the process of preparing a molten steel composed of Fe and other inevitable impurities,
Critical deformation, which is a deformation value at a point where a lamination induced energy (SFE) value of an austenite phase represented by the following [Equation 2] is 19 to 37 and a work-induced martensite phase is formed and starts to contribute to work hardening as a range of values is 0.1 to 0.25, comprising the step of processing the step of producing a strip solidifying while passing between the molten steel a pair of casting rolls including the molten steel in the stainless steel ,
A method for producing a lean duplex stainless steel, wherein nitrogen exceeding the nitrogen solid solution limit contained in the molten steel in the step of producing the strip is discharged to the outside of the solidified shell through the casting roll. 10. The manufacturing method according to claim 9, wherein in the step of manufacturing the strip, a casting roll having a gas discharge channel formed in a circumferential direction on an outer peripheral surface is used as at least one of the pair of casting rolls. 2. A method for producing a lean duplex stainless steel according to 1. A plurality of gas discharge channels formed on a casting roll used in manufacturing the strip have a width of 50 to 500 μm and a depth of 50 to 300 μm, and a gap between adjacent gas discharge channels is 100 to 1000 μm. The method for producing a lean duplex stainless steel according to claim 10 , wherein irregularities of 15 to 25 μm are formed on the surface of the casting roll. In the process of preparing the molten steel, the molten steel is mass%, and further includes at least one of W: 0.1 to 1.0% and Mo: 0.1 to 1.0%. The method for producing a lean duplex stainless steel according to claim 6 or 9 . In the process of preparing the molten steel, the molten steel is in mass%, Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05% and V: 0.001 to 0.15%, The method for producing a lean duplex stainless steel according to claim 6 or 9 , further comprising at least one kind.

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