US20170326628A1 - Lean duplex stainless steel and method for producing the same - Google Patents

Lean duplex stainless steel and method for producing the same Download PDF

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US20170326628A1
US20170326628A1 US15/536,391 US201515536391A US2017326628A1 US 20170326628 A1 US20170326628 A1 US 20170326628A1 US 201515536391 A US201515536391 A US 201515536391A US 2017326628 A1 US2017326628 A1 US 2017326628A1
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stainless steel
duplex stainless
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molten steel
lean duplex
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Jeom Yong CHOI
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Posco Holdings Inc
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    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/16—Controlling or regulating processes or operations
    • B22D11/22—Controlling or regulating processes or operations for cooling cast stock or mould
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/04—Continuous casting of metals, i.e. casting in indefinite lengths into open-ended moulds
    • B22D11/055—Cooling the moulds
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/12—Accessories for subsequent treating or working cast stock in situ
    • B22D11/124—Accessories for subsequent treating or working cast stock in situ for cooling
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/16—Controlling or regulating processes or operations
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite

Definitions

  • the present invention relates to lean duplex stainless steel having a ferritic-austenitic structure and a method for producing the same.
  • the austenitic stainless steel with good workability and corrosion resistance contains iron (Fe) as a base metal and chromium (Cr) and nickel (Ni) as main raw materials, and has been developed as a variety of steel to meet various applications by adding other elements such as molybdenum (Mo) and copper (Cu).
  • 300 series stainless steel which is excellent in corrosion resistance and workability, contains expensive raw materials such as Ni and Mo, and thus, as an alternative to this, 200- and 400-series stainless steels have been discussed.
  • 200-series and 400-series stainless steels have disadvantages respectively in that their formability and corrosion resistance do not reach to 300-series stainless steel.
  • duplex stainless steel in which an austenite phase and a ferrite phase are mixed has all the advantages of the austenitic stainless steel and the ferritic stainless steel, and thus, to date, various types of duplex stainless steels have been developed.
  • the present invention provides lean duplex stainless steel capable of ensuring excellent elongation and corrosion resistance and a method of manufacturing the same, by controlling the content of alloy components to reduce costs and also satisfactorily control the stacking fault energy value present in the lean-duplex stainless steel,
  • the present invention provides a lean duplex stainless steel capable of ensuring excellent elongation and corrosion resistance and a method of manufacturing the same, by satisfactorily controlling the value of the critical strain for strain induced martensite formation.
  • the present invention provides a lean duplex stainless steel capable of solving the problem of massive release of nitrogen gas due to a sharp reduction in nitrogen solubility during solidification from a liquid phase to a solid phase in casting, and a method of manufacturing the same.
  • the lean duplex stainless steel according to one embodiment of the present invention is a ferrite-austenitic stainless steel wherein it is preferable that the stacking fault energy (SFE) value of the austenite phase represented by the following formula 2 is 19 to 37 and the range of the value of the critical strain for strain induced martensite formation is 0.1 to 0.25.
  • SFE stacking fault energy
  • Ni, Cu, Cr, N, Si and Mn refer to the overall content (wt. %) of respective constituent element respectively
  • K(x) is represented by the following formula 3 as the distribution coefficient of respective constituent element (x)
  • V( ⁇ ) is the fraction of the austenite phase (the range of 0.45 to 0.75).
  • the elongation of the stainless steel is 45% or more.
  • the stainless steel comprises, by weight, C: 0.08% or less (excluding 0%), Si: 0.2 to 3.0%, Mn: 2 to 4%, Cr: 18 to 24%, Ni: 0.2 to 2.5%, Cu: 0.2 to 2.5%, balance Fe and the other unavoidable impurities.
  • the stainless steel may further comprise, by weight, at least one of W: 0.1 to 1.0% and Mo: 0.1 to 1.0%.
  • the stainless steel may further include, by weight, at least one of Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05%, and V: 0.001 to 0.15%.
  • the method of manufacturing the lean duplex stainless steel according to an embodiment of the present invention is a method of manufacturing a ferritic-austenitic lean duplex stainless steel, comprising preparing a molten steel; and treating the molten steel to form the stainless steel wherein the molten steel is treated so that the stacking fault energy (SFE) value of the austenite phase represented by the following formula 2 is 19 to 37 and the range of the value of the critical strain for strain induced martensite formation is 0.1 to 0.25:
  • SFE stacking fault energy
  • Ni, Cu, Cr, N, Si and Mn refer to the overall content (wt. %) of respective constituent element respectively
  • K(x) is represented by the following formula 3 as the distribution coefficient of respective constituent element (x)
  • V( ⁇ ) is the fraction of the austenite phase (the range of 0.45 to 0.75):
  • K( x ) [content of element x in ferrite phase]/[content of element x in austenite phase] Formula 3.
  • the process of treating the molten steel to form the stainless steel may comprise temporarily storing the molten steel in the tundish while maintaining the temperature of the molten steel at the temperature higher than the theoretical solidification temperature by 10 to 50° C.; primarily cooling the molten steel by injecting the molten steel in the tundish into the mold and passing the molten steel through the mold while maintaining a cooling rate of 500 to 1500° C./min; and secondarily cooling the molten steel having the solidified shell formed by the primary cooling process while drawing it into a segment and passing through.
  • the cooling water of 0.25 to 0.35 L/kg on the molten steel having the formed solidified shell.
  • the method may further comprise, tertiarily cooling, after the secondary cooling process, by spraying the cooling water of 100 to 125 L/kg ⁇ min on the surface of the cast-slab in the range of the surface temperature of the cast-slab being drawn of 1100 to 1200° C. wherein the cooling water is mixed with air such that the ratio of air to cooling water (air/cooling water) is 1.0 to 1.2
  • the process of treating the molten steel to form the stainless steel may comprise producing a strip by solidifying the molten steel while passing it between a pair of casting rolls wherein nitrogen contained in the molten steel in the process of producing the strip and exceeding a nitrogen solubility limit can be discharged through the casting roll to the outside of the solidified shell.
  • the gas discharge channel formed in the casting roll used in the process of producing the strip has a width of 50 to 500 ⁇ m and a depth of 50 to 300 ⁇ m, a plurality of gas discharge channels is formed in the casting roll the gap between adjacent gas discharge channels is 100 to 1000 ⁇ m and unevennesses of 15 to 25 ⁇ m are formed on the surface of the casting roll.
  • the molten steel comprises, by weight, C: 0.08% or less (excluding 0%), Si: 0.2 to 3.0%, Mn: 2 to 4%, Cr: 18 to 24%, Ni: 0.2 to 2.5%, N: 0.15 to 0.32%, Cu: 0.2 to 2.5%, balance Fe and the other unavoidable impurities.
  • the molten steel may further comprise, by weight, at least one of W: 0.1 to 1.0% and Mo: 0.1 to 1.0%.
  • the molten steel may further comprise, by weight, at least one of Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05%, and V: 0.001 to 0.15%.
  • the embodiment of the present invention it is possible to save resources and significantly reduce the cost of raw materials by controlling the contents of the alloy components of Ni, Si and Cu which are high-priced elements, and in particular, it is sufficiently usable as an alternative to the 200 and 300 series (STS 304, 316) used for molding, by ensuring corrosion resistance and excellent elongation at the same level or higher compared to STS 304.
  • the temperature of the molten steel and the cooling rate can be controlled to suppress the pinholes generated inside the cast-slab.
  • the generation of the internal pores and the occurrence of surface defects can be prevented by smoothly inducing the discharge of the nitrogen gas generated when solidified from a liquid phase to a solid phase, through the improvement of the casting roll.
  • FIG. 1 is a view showing a value of critical strain for strain induced martensite phase formation as stress-strain curves of an inventive steel according to an embodiment of the present invention and a comparative steel.
  • FIGS. 2 a and 2 b are photographs showing the formation of strain induced martensite phases as representative transmission electron microscope photographs of a inventive steel according to an embodiment of the present invention and a comparative steel.
  • FIG. 3 is a graph showing the relationship between the elongation and the value of critical strain for strain induced martensite phase formation.
  • FIG. 4 is a schematic view illustrating a manufacturing process of continuous casting method of the lean duplex stainless steel according to an embodiment of the present invention.
  • FIG. 5 is a schematic view illustrating a manufacturing process of strip casting method of the lean duplex stainless steel according to an embodiment of the present invention.
  • FIG. 6 is a schematic view illustrating a casting roll required in a manufacturing process of strip casting method of the lean-duplex stainless steel according to an embodiment of the present invention.
  • FIG. 7 is a photograph of the structure of the comparative material I and the inventive material A.
  • FIG. 8 is a photograph of surface defects of the comparative material H.
  • FIG. 9 is a photograph of surface defects of the comparative material F.
  • the present invention relates to lean duplex stainless steel having ferritic-austenitic structures wherein the ferritic-austenitic structures referred to in the present invention means that ferrite phase and austenite phase occupy most of the structures, and does not means that the stainless steel is formed only in ferrite phase and austenite phase.
  • the description that the ferrite phase and the austenite phase occupy most of the structures means that the sum of the ferrite phase and the austenite phase in the structures for forming stainless steel accounts for 90% or more, and the remainder, except for the ferrite phase and the austenite phase, may be occupied by the austenite phase transformed martensite phase.
  • FIG. 1 is a view showing a value of critical strain for strain induced martensite formation as stress-strain curves of the inventive steel according to an embodiment of the present invention and the comparative steel.
  • the present invention relates to lean duplex stainless steel having two-phase structures formed of a ferrite phase and an austenite phase comprising, by weight (unless otherwise specified below, the content of the ingredients is % by weight), C: 0.08% or less (excluding 0%), Si: 0.2 to 3.0%, Mn: 2 to 4%, Cr: 18 to 24%, Ni: 0.2 to 2.5%, N: 0.15 to 0.32%, Cu: 0.2 to 2.5%, balance Fe, and the other unavoidable impurities.
  • the lean duplex stainless steel may further comprise by weight at least one of W: 0.1 to 1.0%, Mo: 0.1 to 1.0%, Ti: 0.001 to 0.1%, Nb: 0.001 to 0.05%, and V: 0.001 to 0.15%.
  • C is an element for forming an austenite phase and is an effective element for increasing the strength of a material by solid solution strengthening.
  • C when added excessively, easily bonds with carbide-forming elements such as Cr, which is effective for corrosion resistance at the ferrite-austenite phase boundary, thereby reducing the Cr content around the grain boundary and thus reducing the corrosion resistance, C is preferably added in a range of more than 0% to 0.08% or less in order to maximize the corrosion resistance.
  • Si is partially added for the deoxidation effect, is an element for forming a ferrite phase and is an element which is concentrated on the ferrite phase during the annealing heat treatment. Therefore, 0.2% or more should be added to ensure a proper ferrite phase fraction.
  • the excessive addition of Si exceeding 3.0% increases the hardness of the ferrite phase drastically and thus affects the lowering of the elongation of the two-phase steel, and also makes it difficult to ensure the austenite phase for ensuring sufficient elongation.
  • Si when being added excessively, lowers the slag fluidity at the time of steelmaking and bonds with oxygen to form inclusions, thereby deteriorating the corrosion resistance. Therefore, the content of Si is preferably limited to 0.2 to 3.0%.
  • Mn is an element that increases the deoxidizing agent and nitrogen solubility, and is an austenite phase forming element and can be used for replacing expensive Ni.
  • Mn is added in a large amount, excessive Mn is effective for solubility of nitrogen, but combines with S in the steel to form MnS, thereby deteriorating the corrosion resistance. Therefore, when its content exceeds 4%, it becomes difficult to ensure the corrosion resistance of the level of 304 steel.
  • the content of Mn is less than 2%, even if Ni, Cu, N and the like which are the austenite phase forming elements are controlled, it is difficult to ensure a proper austenite phase fraction, and the solubility of N to be added is low and thus sufficient solution of nitrogen at normal pressure cannot be obtained. Therefore, the content of Mn is preferably limited to 2 to 4%.
  • Cr is an element for stabilizing the ferrite phase together with Si and plays a major role in ensuring the ferrite phase of the two-phase stainless steel, as well as it is an essential element for ensuring corrosion resistance.
  • the Cr content is increased, the corrosion resistance increases, but in order to maintain the phase fraction, the content of expensive Ni or other austenite phase forming elements must be increased. Accordingly, in order to ensure corrosion resistance equal to or higher than that of 304 steel while maintaining the phase fraction of the two-phase stainless steel, it is preferable to limit the content of Cr to 18 to 24%.
  • Ni is an element for stabilizing the austenite phase together with Mn, Cu and N, and plays a main role in ensuring the austenite phase of the duplex stainless steel.
  • Mn and N which are other austenite phase forming elements, instead of maximally decreasing the content of expensive Ni for cost reduction, it is possible to sufficiently maintain the phase fraction balance that can be influenced by the reduction of Ni.
  • the content of Ni is preferably limited 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 one of the elements which is mostly concentrated on the austenite phase due to the high diffusion rate on the solid phase during the annealing heat treatment. Therefore, the increase of the N content can incidentally induce an increase in corrosion resistance and an increase in strength.
  • the solubility of N varies depending on the content of Mn added. When the N content exceeds 0.32% in the range of Mn of the present invention, it is difficult to stably produce steel due to generation of surface defects such as blow holes or pinholes during casting due to excessive nitrogen solubility.
  • N is added in the amount of 0.15% or more in order to ensure the corrosion resistance of the level of 304 steel. If the content of N is too low, it becomes difficult to ensure a proper phase fraction. Therefore, the content of N is preferably limited to 0.15 to 0.32%.
  • Cu is an element for stabilizing the austenite phase together with Mn, Ni and N, and it is desirable that the content of Cu, which plays the same role as Ni, should be minimized for cost reduction. However, it is preferable to add at least 0.2% in order to ensure the stability of the austenite phase sufficient to suppress the excessive formation of the strain induced martensite phase occurring during the cold processing. On the other hand, if the content of Cu exceeds 2.5%, it becomes difficult to process the product due to hot brittleness. Therefore, the content of Cu is preferably adjusted to 0.2 to 2.5%.
  • W and Mo are elements for forming the austenite phase and elements for improving corrosion resistance, and are elements that promote the formation of an intermetallic compound at 700 to 1000 ⁇ during heat treatment, resulting in deterioration of corrosion resistance and mechanical properties.
  • the content of W and Mo exceeds 1% respectively, an intermetallic compound is formed, which may lead to a rapid decrease in corrosion resistance and particularly in elongation.
  • 0.1% or more can be added in order to exhibit the effect of improving the corrosion resistance. Therefore, the content of W and Mo is preferably limited to 0.1 to 1.0% respectively, and at least one of W and Mo may be contained.
  • Ti, Nb, and V are elements that react with nitrogen to form nitrides, and they are crystallized as TiN, NbN, and VN respectively in the molten steel and act as nucleation sites on the ferrite phase during solidification, so that sufficient solidification can proceed even when the cooling rate is increased, thereby suppressing the breakage of the slab.
  • these elements are sufficiently dissolved during the manufacturing process, i.e., reheating or hot rolling, and react with carbon and nitrogen during cooling to form carbonitride and thus inhibit the formation of Cr carbide, thereby contributing to improvement of corrosion resistance.
  • they are elements that inhibit the formation of Cr carbide in the heat affected zone during welding.
  • the present invention maintains elongation and corrosion resistance excellently by controlling the content of the alloy elements, the distribution coefficient and the phase fraction to control the stacking fault energy
  • the following formula 1 is a formula for deriving the stacking fault energy by utilizing the content of all components in the alloy.
  • the applicants of the present invention have found that as a result of measuring and calculating the stacking fault energies of the inventive steel by various methods, it is more accurate to predict the properties of the alloy by calculating the stacking fault energy while utilizing the content of the components of the austenite structure rather than calculating it using only the content of the components of the overall alloy composition as in formula 1.
  • the applicants of the present invention have found that for this purpose, calculating the stacking fault energy taking into account the interstitial distribution coefficient of the alloying elements rather than calculating the stacking fault energy by using only the component content of the overall alloy composition can obtain a more approximate approximation of the actually measured stacking fault energy value.
  • K(x) is represented by the following Formula 3 as the distribution coefficient of respective constituent element (x).
  • K( x ) [content of element x in ferrite phase]/[content of element x in austenite phase] Formula 3.
  • the applicants of the present invention measured the distribution coefficients for respective alloying elements in various annealing conditions and alloy systems by using Fe-EPMA and FE-TEM which are more accurate than EDAX analysis by conventional scanning electron microscopy. At this time, it was confirmed that the distribution coefficients for most of the measured alloying elements are not changed depending on the change of temperature at 900 to 1200° C., which is the temperature range of hot rolling annealing or cold rolling annealing.
  • N 0.2 to 0.32%
  • the alloying elements N and Si mean the entire components of the stainless steel.
  • N 0.15 to 0.32%
  • V( ⁇ ) is the austenite phase fraction
  • austenite phase fraction is defined by the following relation function
  • V( ⁇ ) is the ferrite phase fraction, and V( ⁇ ) has the range of 0.45 to 0.75.
  • the stacking fault energy of the austenite phase is known to control the transformation mechanism of the austenite phase.
  • the stacking fault energy of the austenite phase in the case of single-phase austenitic stainless steel represents the extent to which the externally applied plastic strain energy contributes to the strain of the austenite phase.
  • the lower the stacking fault energy the greater the degree of formation of the strain induced martensite phase that contributes to the work hardening of the steel after formation of the epsilon martensite phase on the austenite phase. If the stacking fault energy is moderate, mechanical twin is formed in the austenite phase.
  • the strain induced martensite phase is formed at the intersection of these twinning, so that the applied plastic strain energy causes mechanical phase change and thus transformation from the austenite phase to martensite phase. Therefore, the stainless steel is known to form the strain induced martensite phase in a very broad range, except for difference in the mesophase (the epsilon martensite phase or the mechanical twin). Accordingly, when the stacking fault energy is less than 50 mJ/m 2 , the strain induced martensite phase is formed after the epsilon martensite phase is formed in the austenite phase, or the strain induced martensite phase is formed after the mechanical twin is formed in the austenite phase.
  • the applicants of the present invention corrected the formula as in formula 2 above, taking into account the distribution coefficients of the alloying elements distributed to the austenite phase after various manufacturing processes and heat treatments.
  • the epsilon martensite phase is first formed as the medium phase and the martensite phase is formed at the intersection of the formed epsilon martensite phases.
  • these martensite phases are rapidly formed at the beginning of the strain, and thus a phenomenon in that the elongation is lowered due to the rapid work hardening is observed.
  • the stacking fault energy of the austenite phases calculated by using the corrected formula is more than 37, as a result of investigation with a transmission electron microscope, formation of the martensite phase is not observed after plastic strain. Therefore, it can be seen that the preferable range of the stacking fault energy of the austenite phases is 19 to 37.
  • the lean duplex stainless steel according to the present invention is preferably formed as austenite phases of 45 to 75% and ferrite phases of 25 to 55% by volume fraction.
  • the fraction of the austenite phases is less than 45%, an excessive concentration phenomenon of the austenite phase forming elements in the austenite phases occurs during the annealing.
  • the austenite phases are sufficiently stabilized to suppress transformation of the strain induced martensite phase which is generated during the strain, and the amount of the austenite phase due to sufficient solid solubility of the alloying elements is increased and thus the tensile strength of the material can be also sufficiently ensured.
  • the fraction of the austenite phase is preferably 45% or more.
  • the fraction of the austenite phase exceeds 75%, surface cracks, etc. occur during hot rolling, thereby resulting in deterioration of the hot workability and loss of properties as two-phase structure steel. Therefore, the fraction of the austenite phase is preferably 75% or less.
  • the range of the critical strain value for strain induced martensite phase formatin during the cold processing or the tensile strain is maintained at 0.1 to 0.25.
  • the amount of critical strain for strain induced martensite phase formation was measured from the inflection point of the stress-strain curve, and this usually refers to the strain value at the time of contributing to the work hardening of the martensite phases in the steel in which the strain induced martensite phase is formed.
  • a tensile test is performed at room temperature (for example, 20 to 25° C.) in a strain rate of 1.0 ⁇ 10 ⁇ 3 /s using the tensile tester until the material is broken.
  • the slope change of the true strain-true stress curve obtained at this time is the work hardening rate.
  • the change in the work hardening rate is closely related to the formation of the strain induced martensite phase.
  • the work hardening rate it gradually decreases as the tensile strain progresses after yielding, and then an inflection point is formed at the time point at which the strain induced martensite phase is generated and begin to contribute to the work hardening, i.e. at the critical strain. Then, when the tensile strain progresses due to the strain beyond the inflection point, and at the same time, when the formation of the strain induced martensite phase is increased, the work hardening rate is increased again.
  • the critical strain value is a strain value at a point where the strain induced martensite phase is formed and begins to contribute to work hardening, and refers to the strain value at the point corresponding to the inflection point in the stress-strain curve obtained by the tensile test and mathematically refers to the point at which the second derivative value of the curve becomes zero.
  • the range of the critical strain value for strain induced martensite phases formation is preferably 0.1 to 0.25.
  • the strain induced martensite phase is a mild phase formed when the unstable austenite phase is strained, and induces work hardening and thus contributes to the increase in the elongation of the steel.
  • the stability of the austenite phase can be controlled using the proper distribution of alloying elements.
  • a rapid solidification method is utilized as a method for enabling proper distribution of alloying elements.
  • the austenite phase and the ferrite phase to be formed are solidified in a non-equilibrium state.
  • these non-equilibrium solidification phases are subjected to hot-rolling annealing for a short period of time, it is possible to control the stability of the austenite phases sufficiently within a desired range by utilizing the distribution of the generated alloying elements.
  • the alloy was designed so that most of nitrogen is segregated on the austenite phases by keeping the content of nitrogen, which has a fast diffusion rate in the solid phase, at the content higher than the normal content.
  • the specimens were prepared using a molten steel whose component contents were adjusted as shown in the following table 1. Thereafter, the phase fraction of the material was controlled by proceeding with hot rolling, hot rolling, cold rolling and then cold rolling annealing, and then the elongation and corrosion resistance were measured.
  • the tensile test specimens were measured by adjusting the tensile strain rate to a rate of 1.0 ⁇ 10 ⁇ 3 /s at room temperature after processing the specimens of ASTM-sub size in parallel to the rolling direction.
  • the following table 1 shows the alloy composition (wt. %) of the test steel.
  • table 2 shows the phase fractions of the ferrite phases and the austenite phases of some experimental steels in table 1 above measured after annealing heat treatment thereof at 1100° C.
  • table 3 shows the results of the stacking fault energy values, the differences in Gibbs free energies, the presence or absence of the strain induced martensite phase, the critical strain values, and the elongations for the inventive steels used in the description of the present invention and comparative steels, which were calculated by formula 2 while taking into account the stacking fault energy values, distribution coefficients and phase fraction calculated by formula 1 without considering the distribution coefficient
  • the Gibbs free energy difference and formation of strain induced martensite phase are closely related. For example, it means that when the Gibbs free energy difference ( ⁇ G) is positive, a strain induced martensite phase is not formed, and when the Gibbs free energy difference ( ⁇ G) is negative, a strain induced martensite phase is formed.
  • the Gibbs free energies of the austenite phases and the martensite phases were calculated using commercial software FACTsage 6.4 (Thermfact and GTT-Technologies).
  • FACTsage 6.4 Thermfact and GTT-Technologies.
  • the amount of the alloy component present in the austenite phase can be calculated by using the distribution coefficient and the phase fraction shown in the present invention.
  • the phase fraction varies depending on the alloy components and the heat treatment temperature.
  • table 2 shows the fractions of ferrite phases and austenite phases when heat-treating comparative steels 1 to 5 and inventive steels 1 to 10 at 1100° C., respectively. It can be seen that in inventive steels 1 to 10, the fractions of the ferrite phases are in the range of about 25 to 55% and the fractions of the austenite phase are in the range of 45 to 75%. Meanwhile, in comparative steel 2, when heat-treating at 1100° C., the fraction of ferrite phases was 83%, and at this time, the fraction of the austenite phases was 17%. That is, it can be seen that comparative steel 2 is not included in the range of the fractions of the ferrite phase and the austenite phase of the present invention.
  • comparative steel 4 has the stacking fault energy of the austenite phase of 17.88 mJ/m 2 when considering distribution, and thus is not included in the appropriate range of the stacking fault energy (SFE) value of the austenite phase.
  • SFE stacking fault energy
  • FIG. 1 is a representative nominal strain-nominal stress comparison curve obtained by the present invention wherein the results of the tensile test after annealing each material at 1100° C. are shown.
  • comparative steel 3 it was confirmed that the value of the critical strain for strain induced martensite phase formation at the time of plastic strain is 0.1 or less (inflection point; indicated by an arrow). As a result, it was predicted that the elongation would be lowered due to rapid work hardening in accordance with rapid formation of strain induced martensite phases. In fact, comparative steel 3 is two-phase stainless steel composed of ferrite phase and austenite phase, but the elongation was about 35% which is very inferior.
  • the critical strain value is less than 0.1, a sharp reduction of the elongation is caused by the hardening of the material due to the rapid work hardening while the strain induced martensite phase is rapidly formed.
  • the critical strain value is more than 0.25, the strain induced martensite phase is formed too late and the local necking of the material due to the strain cannot be suppressed.
  • the value of the critical strain for strain induced martensite phase formation during the cold processing is preferably 0.1 to 0.25.
  • FIGS. 2 a and 2 b show transmission electron microscopic microstructures of comparative steel 1 and inventive steel 1, respectively.
  • comparative steel 1 as shown in FIG. 2 a , it can be seen that strain bands or mechanical twins due to the strain are observed, but the strain induced martensite phase is not observed.
  • invention steel 1 as shown in FIG. 2 a , it can be seen that the strain induced martensite phase is formed at the intersection of the strain bands or the mechanical twins (the strain induced martensite phase is indicated by an arrow).
  • FIG. 3 is a graph showing the relationship between the elongation and the value of the critical strain for strain induced martensite phase formation. The results of measuring the critical strain for martensite phase formation referring to the stress-strain curve of FIG. 1 are shown in FIG. 3 .
  • the lean duplex stainless steels according to the present invention can be manufactured by both the continuous casting method and the strip casting method by solving the problem of nitrogen gas generation or emission depending on the difference in nitrogen solubility when solidified from a liquid phase to a solid phase.
  • FIG. 4 is a schematic view illustrating the manufacturing process of the continuous casting method of the lean duplex stainless steel according to an embodiment of the present invention.
  • the lean duplex stainless steel according to one embodiment of the present invention is manufactured in a conventional continuous facility 100 in which ladle 110 , tundish 120 , mold 130 , and a plurality of segments 140 are sequentially disposed. Additionally, the rear end of the segment 140 is further provided with a spraying mean 150 for spraying air and cooling water mixed with each other.
  • a molten steel having the above-described alloy components is prepared and moved to the ladle 110 , and then temporarily stored in the tundish 120 using the shrouding nozzle 111 .
  • the molten steel temporarily stored in the tundish 120 is maintained at a temperature higher than the theoretical solidification temperature by 10 to 50° C.
  • ⁇ T (° C.), which is the difference between the temperature of the molten steel and the theoretical solidification temperature in the tundish 120 , has a lower limit of 10° C. and an upper limit of 50° C.
  • ⁇ T is lower than the lower limit of 10° C.
  • solidification of the molten steel M can proceed in the tundish 120 , thereby causing a problem in continuous casting
  • ⁇ T exceeds the upper limit of 50° C., the solidification rate is lowered during solidification and the solidification structure is coarsened, so that solidification cracks in the continuous casting cast slab and linear defects during hot rolling are liable to occur.
  • the molten steel M is injected into the mold 130 using the immersion nozzle 121 .
  • the molten steel M is passed through the mold 130 to be primary cooled while maintaining the cooling rate of the molten steel M at 500 to 1500° C./min.
  • the amount of cooling (primary cooling) in the mold 130 and the amount of cooling (secondary cooling) in the segment 140 during the continuous casting are reduced and thereby, during casting, the heat transfer of the cast slab S is delayed, the strength of the cast slab solidified layer is lowered, and the cast slab is bulged, resulting in the deterioration of operation and quality.
  • the cooling rate is controlled to exceed 1500° C./min, although it is very advantageous from the viewpoint of nitrogen pinhole, continuous operation is not possible due to the current limitations of the continuous casting facility, and the time for diffusing the segregation of solute elements remaining between dendrites during continuous casting is shortened, thereby generating surface cracks of the cast slab.
  • the cooling rate during the primary cooling in the mold 130 it is preferable to set the cooling rate during the primary cooling in the mold 130 to 500 to 1500° C./min.
  • the molten steel M that is, the cast slab S having the solidified shell formed in the casting mold 130 is drawn into the segment 140 and thus cooled secondarily, and at this time, it is preferable to spray the cooling water of 0.25 to 0.35 L/kg to the cast slab S.
  • the reason for limiting the spraying amount in the segment 140 in this way is as follows.
  • the solidified structure can be finely formed, but if the amount of the sprayed water exceeds 0.35 L/kg since the period of time for the segregated impurities to diffuse between the solidified structures in the continuous casting process is reduced, the cast slab is present in a sigma phase, and cracks are generated on the surface of the cast slab. In addition, since not only cracks due to thermal stress but also residual stress is generated on the surface excessively, surface cracks occur during grinding of the cast slab.
  • the range of the amount of the sprayed water in the segment 140 is preferably 0.25 to 0.35 L/kg.
  • a tertiary cooling is performed on the secondarily cooled cast slab S.
  • the tertiary cooling is carried out by spraying the cooling water of 100 to 125 L/kg-min on the entire surface of the cast slab S in the surface temperature range of the cast slab S at 1100 to 1200° C. wherein the air and cooling water are mixed so that the ratio thereof is 1.0 to 1.2 (air/cooling water), while continuing to draw into the segment 140 .
  • the tertiary cooling is controlled so as to ensure a uniform scale on the surface of the cast slab S.
  • the reason for this is that in the case of the lean duplex stainless steel, since the oxidation amount in the heating furnace is very small, the lubrication effect by the scale during the hot rolling is small and thus it is very difficult to reduce surface cracks. Therefore, in order to prevent the reduction of temperature due to the contact between the roll and the steel sheet during rolling and to reduce the frictional force between the roll and the steel sheet to prevent surface cracking, a dense and thick scale should be formed on the surface of the steel sheet and a peeling should not occur easily during rolling.
  • the reason for limiting the surface temperature of the cast slab S, the amount of the cooling water, and the ratio of the cooling water to air (air/cooling water) is that if the above conditions are not satisfied, a scale having a desired level of thickness (approximately 35 ⁇ m ⁇ 2 ⁇ m) is not formed on the surface of the slab S, and the generated scale is not uniformly formed.
  • the lean duplex stainless steels having the composition according to the present invention were produced by producing a cast slab while changing the molten steel temperature in the tundish, the cooling rate in the casting mold, and the amount of the sprayed water in the secondary cooling zone as shown in table 4, and the degrees of occurrence of pinholes and cracks on the surface of the cast slab generated therefrom are shown together in table 4.
  • the occurrence or not of pinholes in the cast slab was measured by grinding the surface of the slab by about 0.5 mm and observing the grinded surface.
  • inventive materials A to E which satisfy all of the control conditions of the present invention did not cause pinholes in continuous casting of the cast-slab due to nitrogen and did not cause bulging and defects on the surfaces of the hot-rolled coils.
  • the cooling rate in the mold was lower than the range of the present invention, and thus serious pinholes were generated in cast-slab.
  • the amount of the sprayed water in the secondary cooling zone is within the scope of the present invention, and thus the surface of the continuous casting cast-slab is good, but a large number of linear defects occurred during the hot rolling due to the pinholes present in the cast-slab.
  • FIG. 7 is a photograph of the structure of comparative material I and inventive material A produced according to the continuous casting method of the present invention
  • FIG. 8 is a photograph of surface defects of the comparative material H produced according to the continuous casting method of the present invention
  • FIG. 9 is a photograph of surface defects of comparative material F produced according to the continuous casting method of the present invention.
  • FIGS. 8 and 9 are photographs of surface defects of the surface of the hot-rolled coils found after the hot rolling of comparative materials H and F.
  • FIG. 7 it is confirmed that no pinholes were found on the surface of the inventive cast A-slab, but a large number of pinholes were found in comparative material I.
  • FIG. 8 when the surface of the hot-rolled coil after the hot rolling of comparative material H having relatively good occurrence of pinholes are observed, a large number of pinhole-like defects drawn in the rolling direction are observed.
  • FIG. 9 shows that when the surface of the hot-rolled coil after hot rolling of comparative material F is observed, a large number of crack-like surface defects of cast-slab are observed.
  • cast-slab was produced using the lean duplex stainless steel having the composition according to the present invention and subjecting to primary cooling and secondary cooling, while changing the amount of cooling water, the period of spraying time, the air/cooling water ratio, and the surface temperature of cast-slab as shown in table 5, and the thicknesses and the degree of uniformity of the scale obtained therefrom are shown together in Table 5.
  • inventive materials 1 to 4 it can be seen that when spraying the cooling water for 20 to 30 minutes at a cooling rate of 100 to 120 L/kg ⁇ min while keeping the ratio of air/cooling water at 1.0-1.2 at the surface temperature point of cast-slab of 1000 to 1200° C., the scale becomes very uniform and thick.
  • the thicknesses of the oxidized scales were 15 and 26 ⁇ m and the oxidized scales were non-uniform.
  • the scale formation can be optimized, the surface quality can be improved, the cost for process for removing defects can be minimized, and thus the added value can be improved.
  • FIG. 5 is a schematic view showing a manufacturing process of the strip casting method for the lean-duplex stainless steel according to an embodiment of the present invention
  • FIG. 6 is a schematic view of a nitrogen discharge channel formed in the casting roll of the present invention.
  • the lean duplex stainless steel according to an embodiment of the present invention is manufactured in a conventional strip casting facility 200 in which the ladle 210 , the tundish 220 , a pair of casting rolls 230 , the inline rollers 260 , and the winding rolls 270 are sequentially disposed. Additionally, the gas discharge channel 231 is formed on the surface of the casting roll 230
  • a molten steel M having the above-described alloy components is prepared and moved to the ladle 210 , and then temporarily stored in the tundish 220 using the shrouding nozzle 211 . Then, the molten steel M is solidified while passing between a pair of casting rolls 230 through the injection nozzle 221 to produce the strip S, and the manufactured strip S is rolled in an inline roller 260 disposed continuously with a casting roll 230 and wound around a winding roll 270 .
  • a manifold shield 250 is mounted to prevent the surface of the molten metal from being oxidized by contact with air, and an appropriate gas is added to the inside of the manifold shield 250 to appropriately form an anti-oxidizing atmosphere.
  • the molten steel M is rolled through the inline roller 260 while exiting the roll nip where a pair of casting rolls 230 meet, and then passes through the process such as a heat treatment process and a cold rolling process to form a strip S of 10 mm or less.
  • the molten steel M is provided between the side dam 240 and the inner water-cooled twin-drum rolls 230 rotating at a high speed through the injection nozzle 250 in opposite directions wherein the molten steel M is provided so that the molten steel is rapidly cooled by releasing a large amount of heat through the surface of the water-cooled casting roll 230 and the actual yield is improved without cracking the thin plate of desired thickness.
  • the nitrogen discharge channel 231 is formed on the surface of the casting roll 230 , thereby discharging nitrogen exceeding the solubility limit in solidifying the molten steel.
  • the problem of internal porosity due to nitrogen is mostly generated in the process of rapidly cooling the molten steel M while passing between a pair of casting rolls 230 .
  • the discharge of nitrogen exceeding the solid solubility limit in the molten steel M must be performed simultaneously with the passage of the molten steel M through the casting roll 230 .
  • the gas discharge channel 231 is formed 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 to which only the nitrogen gas can be discharged while the molten steel M cannot pass.
  • These gas discharge channels 231 may be formed in the casting roll 230 in various ways and may be formed in the circumferential direction on the surface of the casting roll 230 to guide and discharge the nitrogen gas toward the outside of the casting roll 230 in accordance with the rotation of the casting roll 230 .
  • the gas discharge channel 231 corresponds to a fine channel having a width of 50 to 500 ⁇ m and a depth of 50 to 300 ⁇ m, and a plurality of gas discharge channels 231 are formed in the circumferential direction of the casting roll 230 wherein the distance between adjacent gas discharge channels 231 is about 100 to 1000 ⁇ m.
  • the shape, structure, and application position of the gas discharge channel 231 may be variously modified as long as the function thereof can be achieved.
  • the contact area between the casting roll 230 and the molten steel M passing through the casting roll 230 can be reduced and thus in order to prevent this, it is preferable that unevennesses are preferably formed on the surface of the casting roll. These unevennesses have an average size of 15 to 25 ⁇ m.
  • comparative example 1 is an example in which molten steel having a specific composition is cast by using a general continuous casting method
  • comparative example 2 is an example in which molten steel having a specific composition is cast by using a general strip casting (rapid casting) method
  • examples 1 to 5 are examples in which casting is performed by a strip casting process while discharging nitrogen exceeding the solubility limit in molten steel by using the casting roll according to the present invention.
  • the nitrogen composition of the high ductility duplex stainless steel of the present invention ranges from 1500 to 3200 ppm.
  • 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 phase->delta phase->delta phase+austenite phase, and when the liquid phase changes into the delta phase, the solubility of nitrogen is about 1164 ppm and thus the solubility difference of about 836 to 1836 ppm occurs. Therefore, some of the supersaturated nitrogen in the liquid phase is gasified during solidification to form various pores in the solidified material and also to form many pores in the solidified shell formed on the surface of the material.
  • examples 1 to 5 are strip casting processes according to the present invention in which nitrogen is discharged during the process and pores are not generated in the strip.

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CN114619007A (zh) * 2022-02-14 2022-06-14 包头钢铁（集团）有限责任公司 一种低合金高氮钢连铸坯的生产方法

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