EP3040434B1 - Rostfreier duplexstahl und struktur aus dem rostfreien duplexstahl, meeresstruktur, erdöl/gas-umgebungsstruktur, pumpenlaufrad, pumpengehäuse und ventilkörper zur durchflusseinstellung damit - Google Patents

Rostfreier duplexstahl und struktur aus dem rostfreien duplexstahl, meeresstruktur, erdöl/gas-umgebungsstruktur, pumpenlaufrad, pumpengehäuse und ventilkörper zur durchflusseinstellung damit Download PDF

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EP3040434B1
EP3040434B1 EP13892680.3A EP13892680A EP3040434B1 EP 3040434 B1 EP3040434 B1 EP 3040434B1 EP 13892680 A EP13892680 A EP 13892680A EP 3040434 B1 EP3040434 B1 EP 3040434B1
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stainless steel
duplex stainless
content
phase
steel
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French (fr)
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EP3040434A4 (de
EP3040434A1 (de
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Naoya Okizaki
Yusaku Maruno
Masafumi Noujima
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Hitachi Ltd
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Hitachi Ltd
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    • 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
    • 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
    • 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
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • 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
    • C22C38/004—Very low carbon steels, i.e. having a carbon content of less than 0,01%
    • 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
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • 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
    • C22C38/48—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
    • C22C38/58—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 duplex stainless steels and structures using the same.
  • Duplex stainless steels mainly have a two-phase metal structure including a ferrite phase (alpha phase) and an austenite phase (gamma phase).
  • the duplex stainless steels have a high strength and excel in pitting corrosion resistance and crevice corrosion resistance in a chloride/sulfide environment. Using the properties, the duplex stainless steels are widely used as materials for marine structures and for petrochemical industries.
  • the duplex stainless steels are known to have inferior toughness when exposed to a high temperature under some manufacturing conditions or working conditions. This is because of the formation of hard and fragile intermetallic compounds (sigma phase, chi phase, and Laves phase) mainly containing Cr, Mo, or another element as a principal component; and an embrittlement phase of nitrides/carbides.
  • duplex stainless steels suffer from the precipitation of the intermetallic compounds with increasing contents of Cr, Mo, and W.
  • the duplex stainless steels often suffer from defects caused by blowhole formation during manufacturing because nitrides are precipitated in larger amounts with an increasing nitrogen content, if the duplex stainless steels have an excessively high nitrogen content.
  • a work is subjected to a solution heat treatment at 950°C to 1200°C to give an appropriate phase ratio between the ferrite phase and austenite phase.
  • the work is then subjected to rapid cooling from the solution heat treatment temperature down to room temperature typically by water cooling.
  • This process is performed to avoid the precipitation of the embrittlement phase and to avoid 475°C embrittlement.
  • the process disadvantageously impedes stable manufacturing of large-sized structures, particularly of thick-walled structures prepared by casting or forging. This is because the embrittlement phase is precipitated inside the material steel due to the difference in cooling rate between the surface and the inside of the steel.
  • the disadvantage is, however, trivial in thin-walled materials such as thin sheets and pipes.
  • duplex stainless steel also disadvantageously suffers from reduction in toughness due to the embrittlement phase precipitation in a weld heat affected zone or as a result of annealing performed so as typically to reduce the residual stress.
  • Patent Literature 1 discloses a super duplex stainless steel in order to suppress the formation of sigma phase, chi phase, and other intermetallic compounds that adversely affect corrosion resistance and mechanical properties.
  • the super duplex stainless steel contains, in weight percent, Cr in a content of 21.0% to 38.0%, Ni in a content of 3.0% to 12.0%, Mo in a content of 1.5% to 6.5%, W in a content of 0% to 6.5%, Si in a content of 3.0% or less, Al in a content of 1.0% or less, Mn in a content of 8.0% or less, N in a content of 0.2% to 0.7%, C in a content of 0.1% or less; and at least one element selected from the group consisting of B in a content of 0.1% or less, Cu in a content of 3.0% or less, and Co in a content of 3.0% or less.
  • the super duplex stainless steel desirably further contains at least one element selected from the group consisting of Ca in a content of 0.5% or less, Mg in a content of 0.5% or less, Ta in a content of 0.5% or less, Nb in a content of 0.5% or less, Ti in a content of 1.5% or less, Zr in a content of 1.0% or less, Sn in a content of 1.0% or less, and In in a content of 1.0% or less, as described in the literature.
  • Steel compositions containing Nb are disclosed in EP 0 818 552 A2 and US 4 055 448 A .
  • the super duplex stainless steel in PTL 1 may suffer from embrittlement progression, because the steel has a high nitrogen content and thereby often suffers from nitride formation; and this may impede appropriate dissolution of added elements in the alloy.
  • An object of the present invention is to suppress the formation of intermetallic compounds (sigma phase, chi phase, and Laves phase) and nitrides in a duplex stainless steel and to help the duplex stainless steel to have better corrosion resistance, embrittlement resistance, productivity, weldability, and thermal processability.
  • the invention provides a duplex stainless steel, a structure using the steel and a method for manufacturing the steel in accordance with the attached claims.
  • the present invention in the aspect, adapts a tantalum-containing duplex stainless steel to contain nitrogen in a low content and helps the duplex stainless steel to less suffer from nitride formation.
  • the aspect also helps the duplex stainless steel to have better corrosion resistance, embrittlement resistance, productivity, weldability, and thermal processability because metallic tantalum not forming nitrides inhibits the diffusion or migration of elements that form intermetallic compounds.
  • the present invention relates to duplex stainless steels and structures using them. More specifically, the present invention relates to duplex stainless steels actually providing still better embrittlement resistance and productivity while maintaining good corrosion resistance by suppressing the formation of embrittlement phases; and products using the duplex stainless steels.
  • the embrittlement phases are formed upon manufacturing (upon casting, forging, hot rolling, or welding), upon welding, and upon a heat treatment of highly corrosion-resistant duplex stainless steels and are exemplified by precipitates such as nitrides and carbides; and intermetallic compounds such as sigma phase and chi phase.
  • tantalum is positively added so as to suppress the intermetallic compound formation in a duplex stainless steel, because tantalum inhibits the diffusion or migration of intermetallic-compound-forming elements.
  • the present invention provides a steel as set forth in claim 1.
  • the stainless steel has a nitrogen content of 0.05% to 0.25% and a carbon content of 0.02% or less to suppress the formation of nitrides and carbides.
  • the nitrogen content is preferably 0.05% to 0.19%.
  • the present invention also provides, in an embodiment, a super duplex stainless steel containing, in mass percent, N in a content of 0.05% to 0.25%, C in a content of 0.02% or less, P in a content of 0.02% or less, Si in a content of 0.5% or less, Mn in a content of 1.2% or less, Ni in a content of 6.0% to 8.0%, Cr in a content of 24.0% to 26.0%, Mo in a content of 3.0% to 5.0%, W in a content of 4.0% or less, and Ta in a content of 0.2% to 0.5% and has a pitting resistance equivalent number (PREW) of 40 or more.
  • PREW pitting resistance equivalent number
  • the present invention provides a duplex stainless steel structure obtained by preparing a structure of an alloy having the chemical composition by forging or casting; and subjecting the formed structure to a solution heat treatment at a temperature of 950°C to 1200°C for a time period of 30 minutes to 2 hours so as to have a ratio of austenite phase to ferrite phase of 0.2 to 0.8.
  • the resulting duplex stainless steel structure less suffers from embrittlement phase formation inside of the structure and can provide a product having good toughness.
  • Exemplary particularly useful structures formed from the alloy having the chemical compositions include marine structures; oil & gas structures (structures used in oil & gas environments); and pump impellers, pump casings, and flow rate adjusting devices for use in chemical plant structures.
  • the present inventors intended to help thick-walled cast products, forged products, and hot-worked products to have better productivity and embrittlement resistance while maintaining satisfactory corrosion resistance. Accordingly, they made investigations on technologies for suppressing embrittlement phase precipitation caused by intermetallic compounds and carbonitride. As a result, they had following findings.
  • Figure 1A is a conceptual diagram illustrating an embrittlement phase formation mechanism in a customary duplex stainless steel.
  • the duplex stainless steel in Fig. 1A includes a ferrite phase 1, an austenite phase 2, and a grain boundary 3 formed between the two phases.
  • an element that forms an intermetallic compound such as Cr, Mo, or W diffuses or migrates via a vacancy 4 and moves toward the grain boundary 3.
  • An intermetallic compound 6 and a carbide/nitride 7 are formed in a grain boundary region including the grain boundary 3. These are also called embrittlement phases.
  • a steel, if containing the embrittlement phases in large amounts, may become brittle and may often suffer from deterioration in corrosion resistance, embrittlement resistance, productivity, weldability, and thermal processability.
  • Figure 1B is a conceptual diagram illustrating the embrittlement phase formation suppressing mechanism in the duplex stainless steel according to the present invention.
  • a tantalum atom 11 occupies the vacancy 4 more readily than the intermetallic-compound-forming element 5 does and thereby inhibits the diffusion of the intermetallic-compound-forming element 5. This can prevent the formation typically of a nitride of the intermetallic-compound-forming element 5 in a grain boundary region 12.
  • the intermetallic compound 6 includes, for example, the sigma phase and chi phase and is known to be readily precipitated at the alpha phase-gamma phase interface as an origin and to grow toward the alpha phase.
  • Cr, Mo, Si, and W each acting as an element forming the intermetallic compound 6 (intermetallic-compound-forming element 5) migrate from the metal matrix, are enriched at the grain boundary of the alpha phase/gamma phase interface and are precipitated as the intermetallic compound 6.
  • the precipitation of the intermetallic compound 6 can be retarded probably by reducing the diffusion rate of the intermetallic-compound-forming elements 5.
  • Cr, Mo, and W are over-size elements each having an atomic radius larger than the average atomic radius of elements forming the stainless steel and are considered to intensively interact with the atomic vacancy (vacancy 4) in the metal matrix and to move via the vacancy 4 as a preferential diffusion path.
  • a specific element is added so as to allow the vacancy 4 to trap the added element, where the added element has an atomic radius larger than those of the intermetallic-compound-forming elements and interacts with the vacancy 4 more intensively than the elements do. This can reduce the diffusion rate of the intermetallic-compound-forming elements 5 particularly in a temperature range of 650°C to 950°C where embrittlement phase precipitation is an issue.
  • Embrittlement in the temperature range can be avoided by rapid cooling in the case of a small-sized stainless steel (steel), but becomes an issue in the case of a large-sized steel because it is difficult to rapidly cool the inside of such large-sized steel.
  • the present invention in an embodiment, solves the problem by adapting the steel chemical composition.
  • a metal element having a large atomic radius has an extremely low free energy to form a nitride or carbide.
  • a thermal equilibrium calculation demonstrates the following phenomena: A nitride is formed in a liquid phase during manufacturing and is hardly dissolved in the matrix when Zr, Ti, Hf, or another element having a high nitride/carbide forming capability is added, particularly when the steel is a super duplex stainless steel added with nitrogen so as to have better corrosion resistance.
  • niobium (Nb) is considered to be an element that is readily taken in the sigma phase that acts as an intermetallic compound by itself.
  • tantalum (Ta) is selected as an element to be added, because Ta can be relatively easily dissolved in the metal matrix upon manufacturing and is resistant to precipitation as an intermetallic compound.
  • Chromium (Cr) content 20.0% to 40.0%
  • Chromium element is a basic element and is most important in helping the stainless steel to maintain corrosion resistance at certain level.
  • the stainless steel herein has to have a two-phase structure including both austenite and ferrite, as being a duplex stainless steel.
  • the chromium content is specified to be 20% or more in consideration of a chromium equivalent (Cr eq ) and a nickel equivalent (Ni eq ) as defined by expressions as follows; and a percentage (fraction) of the ferrite phase determined by these equivalents.
  • the upper limit of the chromium content is specified to be 40% in consideration of economic efficiency, because increase in Cr eq requires increase in Ni eq .
  • the chromium content is more preferably in the range of 24% to 26%.
  • Nickel element stabilizes austenite and is useful in increasing general corrosion resistance in relation to corrosion resistance.
  • the nickel content is therefore specified to be 3% or more.
  • the nickel content is specified to be 12% or less in terms of upper limit in consideration of the relationship between the chromium equivalent and the nickel equivalent, the phase fraction, and the economic efficiency.
  • the nickel content is more preferably in the range of 6% to 8%.
  • Molybdenum (Mo) content 7.0% or less
  • Molybdenum element is important in helping the stainless steel to maintain corrosion resistance as with chromium, functionally stabilizes the ferrite phase, but may accelerate the intermetallic compound formation, when it is added. To prevent this, the molybdenum content is controlled to 7.0% or less. The molybdenum content is more preferably in the range of 3.0% to 5.0%.
  • Tungsten (W) content 6.5% or less
  • Tungsten element improves the corrosion resistance.
  • a precipitation rate of intermetallic compounds is reduced by replacing tungsten with Mo in a half amount, and thereby improves the corrosion resistance and mechanical properties.
  • Tungsten is an expensive alloy element.
  • tungsten may accelerate the intermetallic compound formation and adversely affect the corrosion resistance of a weld bead when it is added in a large amount. To prevent this, the tungsten content is controlled to 6.5% or less. The tungsten content is more preferably in the range of 4.0% or less.
  • Silicon (Si) content 3.0% or less
  • Silicon element stabilizes the ferrite phase and is effective for deoxidation during manufacturing.
  • the element also increases the fluidity of a molten steel upon manufacturing and welding and thereby reduces surface defects.
  • the element increases the precipitation rate of intermetallic compounds and reduces the steel ductility.
  • the silicon content is preferably 3.0% or less and more preferably 0.5% or less.
  • Manganese (Mn) content 8.0% or less
  • the element has a deoxidation effect upon melting and refining, but may cause the stainless steel to have inferior corrosion resistance and may accelerate the formation of intermetallic compounds, when it is added in an excessively large amount.
  • the manganese content is preferably controlled to 8% or less in terms of its upper limit. The manganese content is more preferably in the range of 1.2% or less.
  • Nitrogen (N) content 0.7% or less
  • Nitrogen element is useful for improving the resistance to pitting corrosion and is one of most important elements in relation to corrosion resistance, because nitrogen has the effect in a magnitude of about 30 times that of chromium.
  • nitrogen is added to compensate the strength of the steel.
  • nitrogen may cause cracking due to blowhole generation during manufacturing if it is added in a content greater than 0.7%.
  • the nitrogen content is preferably controlled to 0.7% or less. Nitrogen forms Ta-containing nitrides and impairs the effects of Ta addition when nitrogen is present in combination with Ta as added.
  • the nitrogen content is more preferably 0.3% or less, furthermore preferably 0.05% to 0.25%, and particularly preferably in the range of 0.05% to 0.19%.
  • Carbon (C) content 0.1% or less
  • Carbon element forms carbides and induces grain boundary sensitization upon welding. Carbon forms Ta-containing carbides and impairs the effects of Ta addition when carbon is present in combination with Ta as added. To prevent this, the carbon content is preferably minimized. However, reduction in the carbon content may invite increase in production cost. The carbon content is therefore specified to 0.1% or less. The carbon content is more preferably in the range of 0.02% or less.
  • Tantalum (Ta) content 0.05% to 1.0%
  • Tantalum element is one of elements featuring the present invention. Tantalum has an atomic radius larger than the average atomic radius of elements forming the duplex stainless steel, advantageously prevents the diffusion of major intermetallic-compound-forming elements, and effectively reduces the precipitation rate of intermetallic compounds, as is described above.
  • tantalum invites poor economy, and further disturbs the balance in ratio between the ferrite phase and austenite phase, if tantalum is added in an excessively high content.
  • the tantalum content is preferably controlled to 1.0% in terms of its upper limit. In contrast, if tantalum is added in a content less than 0.05%, its effects may be not expected.
  • the tantalum content is more preferably in the range of 0.2% to 0.5% in view of the balance in the amount of Ta to form a solid solution in the nitride phase and the ferrite phase.
  • Phosphorus (P) content 0.1% or less
  • Phosphorus element is an impurity inevitably contaminated into the steel.
  • the element impairs the corrosion resistance, segregates at the grain boundary, and thereby accelerates the embrittlement phase precipitation.
  • the phosphorus content is preferably minimized and is controlled preferably to 0.1% or less, more preferably to 0.02% or less, and particularly preferably to 0.005% or less. Excessive reduction in phosphorus content, however, may invite increase in production cost.
  • the phosphorus content may be determined also in this view.
  • Table 1 indicates chemical compositions (in mass percent) of duplex stainless steels according to Example 1 example useful to understand the invention (Sample Steel C)) and Comparative Examples 1 and 2 (comparative steels (Sample Steels A and B)).
  • Sample Steel A had a chemical composition equivalent to that of a standardized steel S32750.
  • Sample Steel B had a chemical composition having low N, C, and Si contents.
  • Sample Steel C had a chemical composition of an alloy equivalent to Sample Steel B, except for being added with a very small amount of Ta.
  • the ingots were heated to 1250°C, forged, and yielded steel plates of 20 by 50 by 150 (mm).
  • the forged steel plates were each subjected to a solution heat treatment at 1100°C for one hour so as to give an appropriate ratio of austenite phase to ferrite phase; and then rapidly cooled by water cooling so as to avoid embrittlement phase precipitation.
  • Figures 2A, 2B and 2C depict external view images of Sample Steels A, B and C prepared by forging, respectively.
  • the images demonstrate that the sample steels could be manufactured without suffering from cracking and defects due to forging.
  • Figures 3A, 3B and 3C depict metal structure observation results of Sample Steels A, B and C after manufacturing, respectively.
  • sample steels were each subjected to a heat treatment at 800°C, where the temperature is within a temperature range in which an embrittlement phase is readily precipitated. This was performed so as to evaluate the embrittlement phase precipitation under conditions of cooling during manufacturing and of reheating by welding.
  • Figure 4 is a graph illustrating how the amount of residual ferrite varies depending on the heat treatment time at 800°C.
  • the graph is plotted with the abscissa indicating the heat treatment time and the ordinate indicating the ferrite amount.
  • the ferrite amount was measured with a ferrite scope using magnetic induction.
  • the tendency of intermetallic compound precipitation can be evaluated by evaluating the amount of residual ferrite. This is because the intermetallic compound precipitation which is one of embrittlement phase precipitations proceeds as a result of decomposition of the ferrite phase into an intermetallic compound phase and an austenite phase under a precipitation temperature condition.
  • Figure 4 demonstrates that the stainless steel of Example 1 had a lower rate of ferrite phase decrease and less suffered from the decomposition of the ferrite phase than those of Comparative Examples 1 and 2.
  • Figure 5A, 5B and 5C depict metal structure images of Sample Steel A, B and C, respectively, after a heat treatment at 800°C for 30 minutes.
  • the figures demonstrate that an embrittlement phase 53 increased in addition to a ferrite phase 51 and an austenite phase 52.
  • Sample Steel C included the embrittlement phase 53 as precipitated in a smaller amount than those of Sample Steels A and B as the comparative steels (Comparative Examples) and less suffered from the precipitation of the embrittlement phase 53.
  • Sample Steel B included the embrittlement phase 53 in a larger amount.
  • Figure 6 depicts Charpy impact value measurement results after a heat treatment at 800°C for 5 minutes.
  • the Charpy impact value was measured according to Japanese Industrial Standard (JIS) Z2242 (2005). The measurement was performed by a procedure schematically illustrated as follows.
  • Charpy test specimens having a size of 10 mm by 10 mm by 55 mm and having a 2-mm V-notch were sampled from each sample steel plate before and after the heat treatment from the longitudinal direction of the plate so that the central part should be the notched portion, and impact values of the test specimens were measured.
  • Figure 6 demonstrates that Sample Steel C had a Charpy impact value after the heat treatment of higher than those of Sample Steels A and B as the comparative steels. This indicates that the suppression of intermetallic compound formation helped the steel to have better toughness.
  • Figures 7A, 7B , 7C and 7D depict energy-dispersive X-ray (EDX) analyses results at a grain boundary (alpha/gamma interface) after a heat treatment at 800°C for one minute.
  • EDX energy-dispersive X-ray
  • Figures 7A and 7B depict an electron photomicrograph and measurement results of concentration distributions of respective elements at the analysis position (along the analysis line) in the arrow direction in Fig. 7A , respectively, of Sample Steel A as the comparative steel.
  • Figures 7C and 7D depict an electron photomicrograph and measurement results of concentration distributions of respective elements at the analysis position (along the analysis line) in the arrow direction in Fig. 7C , respectively, of Sample Steel C.
  • Figures 7A and 7C depict fine structures of a ferrite phase 71 and an austenite phase 72 clearly, with a grain boundary indicated by dashed lines.
  • the concentrations of respective elements were measured at the analysis position 73 indicated by the line in the arrow direction (from the austenite phase 72 to the ferrite phase 71).
  • Figures 7B and 7D are plotted with the abscissa indicating the distance and the ordinate indicating the concentration.
  • the comparative steel had high Mo and Cr concentrations in the vicinity of the grain boundary facing the ferrite phase.
  • the sample steel C had a Ta concentration peak in the vicinity of the grain boundary facing the ferrite phase and exhibited lower Mo and Cr concentrations than those of the comparative steel in Fig. 7B .
  • tantalum (Ta) preferably diffuses at the ferrite-austenite grain boundary and thereby inhibits the diffusion of Mo and Cr acting as intermetallic-compound-forming elements.
  • sample steel C and comparative steels were each subjected to a heat treatment simulating a post weld heat treatment (PWHT) for residual stress relaxation, and how the heat treatment affects the residual stress and impact value was evaluated.
  • PWHT post weld heat treatment
  • a tensile residual stress was applied to each of Sample Steels A, B and C by subjecting them to surface grinding of the steel plate using a grindstone with a grain size of #46 at a rotation speed of 1440 rpm to a depth of cut of 0.01 mm and thereby imparting a highly deformed layer to the surface.
  • the test samples each applied with the residual stress in the surface by surface grinding were subjected to a heat treatment at 650°C for 30 minutes simulating the PWHT, and how the heat treatment conditions affect the residual stress and mechanical properties was evaluated.
  • Figure 8 depicts results of comparisons in residual stress between before and after the heat treatment.
  • the residual stress values were measured for the ferrite phase and for the austenite phase, respectively.
  • the measured values were multiplied by the volume ratio of each phase and averaged, and the average was defined as a macro-stress and evaluated.
  • test samples were imparted with a tensile stress of about 900 to about 1100 MPa by the surface processing (surface grinding), but had a lower tensile stress of about 200 MPa after the heat treatment at 650°C for 30 minutes, indicating that the heat treatment gave a stress relaxation effect of about eight-tenths the initial stress.
  • Figure 9 depicts comparison results in the Charpy impact test between before and after the heat treatment.
  • Figure 9 demonstrates that the sample steel C had a better impact value than those of the comparative steels and maintained an impact value of about 100 J/cm 2 even after the heat treatment. Specifically, the sample steel C maintained an impact value of 100 J/cm 2 or more even after the heat treatment at 650°C for 30 minutes, where the heat treatment relaxed the residual stress by eight-tenths the initial stress.
  • a pitting potentials was measured before and after the heat treatment (at 650°C for 30 minutes), and the results are indicated below.
  • the pitting potential was measured according to JIS G0577 (2005).
  • Figure 10 depicts a comparison in pitting potential among the sample steel C and the comparative steels.
  • Figure 10 demonstrates that the order of pitting corrosion resistance (after the heat treatment) of the respective sample steels is as follows.
  • Sample Steel C > Sample Steel B (comparative steel) > Sample Steel A (comparative steel; corresponding to customary steel S32750).
  • the sample steel C had a pitting potential higher than that of the customary steel.
  • Figure 11 is a cross-sectional view of a vertical mixed flow seawater pump.
  • the vertical mixed flow seawater pump in Fig. 11 includes components such as a bell mouth 117 that reduces the drag of seawater coming in from a feed channel; a shaft 111 that transfers the rotative power of a driving motor; an impeller hub 115 that is fixed to the shaft 111; impeller vanes 113 that impart the rotative power of the driving motor to the seawater efficiently; a casing liner 114 that has a spherical inner wall so as to allow the impeller vanes 113 to have an always constant outer clearance; a casing 112 that converts the velocity energy of the seawater given by the impeller vanes 113 into pressure energy; a column pipe 119 that allows the pressurized seawater to pass therethrough; an impeller cap 116; and a cone 118.
  • a bell mouth 117 that reduces the drag of seawater coming in from a feed channel
  • a shaft 111 that transfers the rotative power of a driving motor
  • an impeller hub 115 that is fixed to the shaft 111
  • the casing liner 114 and the casing 112 were each formed from the steel of Example 1 by casting; whereas the impeller hub 115 and the impeller vanes 113 were each formed from the steel of Example 1 by forging. These steels after casting or forging were subjected to a solution heat treatment at 1100°C for one hour and then subjected to water cooling so as to have a two-phase composition including ferrite in an amount of 40% to 50% to give the structures.
  • a junction between the casing liner 114 and the casing 112 and junctions between the impeller hub 115 and the impeller vanes 113 were bonded by metal inert gas arc welding (MIG welding). Weld heat affected zones were wrapped around with a band heater, heated up to 650°C, and subjected to a heat treatment at that temperature (650°C) for 30 minutes, followed by rapid cooling.
  • MIG welding metal inert gas arc welding
  • the residual stress of the weld heat affected zones was measured by X-ray residual stress measurement to find that the tensile stress was reduced down to 80 MPa.
  • the steel of Example 1, as used, enabled the manufacturing of a seawater pump that less suffered from reduction in weld toughness, had a higher fatigue strength (better fatigue resistance), and had a longer working life.
  • Figure 12 is a cross-sectional view of a flow rate adjusting device.
  • the flow rate adjusting device in Fig. 12 includes components such as a casing 121 that supports the entire device; a valve element 122 that controls a flow rate; a valve seat 123 into which the valve element 122 is fit; a handle 125; and a shaft 124 that controls the position of the valve element 122 by the rotation of the handle 125.
  • the casing 121 was formed from the steel of Example 1 by casting.
  • the steel of Example 1 enabled the manufacturing of a large-sized flow rate adjusting device having satisfactory corrosion resistance, as used.
  • the flow rate adjusting device is usable in environments typically of seawater, petroleum, and chemical plants.
  • impeller vane 114...casing liner, 115...impeller hub, 116...impeller cap, 117...bell mouth, 118...cone, 119...column pipe, 121...casing, 122...valve element, 123...valve seat, 124...shaft, 125...handle.

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  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Steel (AREA)

Claims (9)

  1. Duplex-Edelstahl, in Massenprozent bestehend aus:
    N in einem Gehalt von 0,05% bis 0,25%;
    C in einem Gehalt von 0,02% oder weniger;
    P in einem Gehalt von 0,02% oder weniger;
    Si in einem Gehalt von 0,5% oder weniger;
    Mn in einem Gehalt von 1,2% oder weniger;
    Ni in einem Gehalt von 6,0% bis 8,0%;
    Cr in einem Gehalt von 24,0% bis 26,0%;
    Mo in einem Gehalt von 3,0% bis 5,0%;
    W in einem Gehalt von 6,5% oder weniger; und
    Ta in einem Gehalt von 0,2% bis 0,5%,
    wobei der Rest Fe und unvermeidbare Verunreinigungen ist.
  2. Duplex-Edelstahl nach Anspruch 1, der eine durch den im Folgenden wiedergegebenen Ausdruck festgelegte Lochfraß-Widerstandsfähigkeits-Äquivalenzzahl (PREW) von 40 oder mehr aufweist: PREW = % ⁢ Cr + 3,3 × % Mo + 0,5 × % W + 30 × % ⁢ N
    Figure imgb0007
     wobei %Cr, %Mo, %W und %N entsprechenderweise die Gehalte von Cr, Mo, W and N in Massenprozent sind.
  3. Duplex-Edelstahlstruktur unter Verwendung des Duplex-Edelstahls nach einem der Ansprüche 1 und 2.
  4. Verfahren zur Gewinnung einer Duplex-Edelstahlstruktur, indem:
    durch Schmieden oder Gießen eine Struktur aus einem Duplex-Edelstahl nach einem der Ansprüche 1 und 2 gebildet wird; und
    die gebildete Struktur einer Lösungs-Wärmebehandlung bei einer Temperatur von 950°C bis 1200°C über eine Zeitspanne von 30 Minuten bis 2 Stunden ausgesetzt wird, sodass sie ein Verhältnis der Austenit-Phase zur Ferrit-Phase von 0,2 bis 0,8 aufweist.
  5. Seetechnische Struktur, die eine Duplex-Edelstahlstruktur nach Anspruch 3 darstellt.
  6. Öl- und gastechnische Struktur, die eine Duplex-Edelstahlstruktur nach Anspruch 3 darstellt.
  7. Pumpenflügelrad, das eine Duplex-Edelstahlstruktur nach Anspruch 3 darstellt.
  8. Pumpengehäuse, das eine Duplex-Edelstahlstruktur nach Anspruch 3 darstellt.
  9. Ventilelement einer Strömungsraten-Einstellvorrichtung, das eine Duplex-Edelstahlstruktur nach Anspruch 3 darstellt.
EP13892680.3A 2013-08-28 2013-08-28 Rostfreier duplexstahl und struktur aus dem rostfreien duplexstahl, meeresstruktur, erdöl/gas-umgebungsstruktur, pumpenlaufrad, pumpengehäuse und ventilkörper zur durchflusseinstellung damit Not-in-force EP3040434B1 (de)

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PCT/JP2013/073038 WO2015029167A1 (ja) 2013-08-28 2013-08-28 二相ステンレス鋼並びにこれを用いた二相ステンレス鋼製構造物、海洋構造物、石油・ガス環境構造物、ポンプインペラ、ポンプケーシング及び流量調節弁の弁体

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JP6482074B2 (ja) * 2014-09-02 2019-03-13 日本冶金工業株式会社 二相ステンレス鋼板とその製造方法
JP6686320B2 (ja) * 2015-08-05 2020-04-22 日本製鉄株式会社 ステンレス鋼管の製造方法
CN107312951A (zh) * 2016-04-26 2017-11-03 天津碧宇舟机械制造有限公司 一种均质机用高强度转子及其制备方法
CN107312979A (zh) * 2016-04-26 2017-11-03 天津碧宇舟机械制造有限公司 一种大功率耐腐蚀泥浆泵叶轮及其制造方法
CN105755397B (zh) * 2016-05-24 2017-07-07 江苏金基特钢有限公司 一种耐腐蚀易成型特种钢的加工方法
SE1950909A1 (en) 2019-07-31 2021-02-01 Ferritico Ab Duplex steel with improved embrittlement properties and method of producing such

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US4055448A (en) * 1973-04-10 1977-10-25 Daido Seiko Kabushiki Kaisha Ferrite-austenite stainless steel
JPS5441214A (en) * 1977-09-08 1979-04-02 Nippon Yakin Kogyo Co Ltd Twoophase highhstrength stainless steel
DE19628350B4 (de) * 1996-07-13 2004-04-15 Schmidt & Clemens Gmbh & Co Verwendung einer rostfreien ferritisch-austenitischen Stahllegierung
JP4031992B2 (ja) * 2001-04-27 2008-01-09 リサーチ インスティチュート オブ インダストリアル サイエンス アンド テクノロジー 優れた熱間加工性を持つ高マンガン二相ステンレス鋼及びその製造方法
KR100460346B1 (ko) * 2002-03-25 2004-12-08 이인성 금속간상의 형성이 억제된 내식성, 내취화성, 주조성 및열간가공성이 우수한 슈퍼 듀플렉스 스테인리스강
EP2677054B1 (de) * 2011-02-14 2020-03-25 Nippon Steel Corporation Duplex-edelstahl-blech oder -rohr und herstellungsverfahren dafür
JP5890330B2 (ja) * 2013-01-15 2016-03-22 株式会社神戸製鋼所 二相ステンレス鋼材および二相ステンレス鋼管

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WO2015029167A1 (ja) 2015-03-05
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