EP2287346A1 - Bainitstähle mit Bor - Google Patents

Bainitstähle mit Bor Download PDF

Info

Publication number
EP2287346A1
EP2287346A1 EP10166261A EP10166261A EP2287346A1 EP 2287346 A1 EP2287346 A1 EP 2287346A1 EP 10166261 A EP10166261 A EP 10166261A EP 10166261 A EP10166261 A EP 10166261A EP 2287346 A1 EP2287346 A1 EP 2287346A1
Authority
EP
European Patent Office
Prior art keywords
composition
less
steel
equal
steel composition
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Granted
Application number
EP10166261A
Other languages
English (en)
French (fr)
Other versions
EP2287346B1 (de
Inventor
Perez Teresa Estela
Gomez Gonzalo Roberto
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Tenaris Connections BV
Original Assignee
Tenaris Connections Ltd
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Tenaris Connections Ltd filed Critical Tenaris Connections Ltd
Publication of EP2287346A1 publication Critical patent/EP2287346A1/de
Application granted granted Critical
Publication of EP2287346B1 publication Critical patent/EP2287346B1/de
Active legal-status Critical Current
Anticipated expiration legal-status Critical

Links

Images

Classifications

    • 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
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/10—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies
    • 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
    • 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
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • 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
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/10—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies
    • C21D8/105—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies 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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • 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
    • C21D2211/008—Martensite
    • 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • C21D9/085—Cooling or quenching

Definitions

  • Embodiments of the present disclosure pertain to seamless pipes formed from steels containing micro-alloying additions of boron and titanium, with yield strengths of at least 100 ksi (690 MPa), excellent toughness, and good weldability.
  • Such pipes are suitable for use as high strength line pipes, for example X100 in API 5L standard, and other possible applications.
  • Micro-alloying additions of boron to steel are desirable, as such additions may improve the mechanical properties of the steel.
  • boron additions may increase hardenability, the ability of steel to be hardened by heat treatment. By migrating to grain boundaries, where they inhibit austenite to ferrite phase transformation, boron additions may improve the ease with which martensite may be formed.
  • boron is effective at very low concentrations, providing significant improvements in hardenability at relatively low cost.
  • boron should remain in its free, elemental state.
  • boron reacts easily with impurities present in the steel, such as nitrogen.
  • impurities present in the steel such as nitrogen.
  • the positive effect on hardenability provided by boron may be reduced, owing to the decrease in free boron.
  • strong nitride formers such as titanium
  • titanium may be added to the steel composition in order to inhibit boron nitrides from forming.
  • relatively coarse titanium nitride particles may be formed during solidification. These particles, which may further grow during reheating prior to hot rolling, can lead to poor toughness in the steel and overshadow the property improvements yielded by the boron addition.
  • a method of making a boron-titanium steel with yield strength of at least 100 ksi (690 MPa), excellent toughness, and good weldability comprises providing a composition comprising carbon, titanium, and boron.
  • the method may additionally comprise providing one or more of manganese, silicon, nickel, chromium, molybdenum, vanadium, and niobium to the composition.
  • the method may also comprise cooling the composition from casting at a cooling rate sufficiently high to inhibit coarsening of titanium nitride (TiN) precipitates within the composition and to limit the size of the TiN precipitates to less than about 50 nm.
  • TiN titanium nitride
  • the method may further comprise hot rolling the composition so as to refine the microstructure and achieve grain sizes of about 20 to 50 ⁇ m, prior to transformation.
  • the method may further include cooling the composition in air after hot rolling and subjecting the composition to austenization and quenching; cooling the composition in air after hot rolling and subjecting the composition to austenization, quenching and tempering; or forced cooling the composition immediately after hot rolling at rates between about 5 to 50°C/sec without any subsequent heat treatment.
  • the steel composition may be formed into a steel pipe, for example, a seamless pipe.
  • a method of making a steel pipe comprises providing a steel composition comprising:
  • a method of making a steel composition comprises providing a steel composition comprising:
  • the final microstructure of the steel composition following cooling, without any tempering after cooling may comprise a mixture of bainite and martensite, with no more than about 30% of martensite. In certain embodiments, the microstructure may comprise no more than about 5% of martensite.
  • a method of making a steel composition comprises providing a steel composition comprising:
  • the final microstructure of the steel composition may comprise a mixture of bainite and martensite.
  • the microstructure comprises no more than about 30% of martensite. In further embodiments, the microstructure comprises no more than about 20% of martensite.
  • a method of making a steel composition comprises providing a steel composition comprising:
  • the final microstructure of the air cooled composition after tempering, may comprise a mixture of tempered bainite and martensite with no less than about 30% martensite. In certain embodiments, the air cooled composition may comprise no less than about 50% of martensite.
  • Figure 1 is a schematic flow diagram of one embodiment of a method of producing boron-titanium (B/Ti) steel pipes;
  • FIG. 1 is a continuous cooling transformation (CCT) plot of one embodiment of steel composition 1;
  • Figure 3 illustrates scanning electron micrographs of the microstructure of embodiments of steel composition 1 cooled from the austenitic range at rates of about 2°C/sec, 5°C/sec, 10°C/sec, and 20°C/sec;
  • Figures 4A and 4B are plots of impact energy (CVN) for an embodiment of steel composition 1 subjected to accelerated cooling; (A) impact energy as a function of cooling rate; (B) impact energy as a function of temperature;
  • Figure 5 is a plot of hardness as a function of tempering temperature for an embodiment of steel composition I in the quenched and tempered condition
  • Figure 6 illustrates a scanning electron micrograph of the microstructure of an embodiment of steel composition 1 that is quenched and tempered at about 410°C;
  • FIG. 7 is a continuous cooling transformation (CCT) plot of one embodiment of steel composition 2;
  • Figure 8 illustrates scanning electron micrographs of the microstructure of embodiments of steel composition 2 cooled from the austenitic range at rates of about 0.2°C/sec, 0.5°C/sec, 1°C/sec, 10°C/sec, 30°C/sec, and 50°C/sec;
  • FIG. 9 is a continuous cooling transformation (CCT) plot of one embodiment of steel composition 3;
  • Figure 10 illustrates scanning electron micrographs of the microstructure of embodiments of steel composition 3 cooled from the austenitic range at rates of about 0.2°C/sec, 0.5°C/sec, 1°C/sec, 10°C/sec, 30°C/sec, and 50°C/sec;
  • Figures 11A-11B are plots of hardness as a function of cooling rate from hot rolling for embodiments of steel compositions 2 and 3; (A) composition 2; (B) composition 3;
  • Figures 12A-12B illustrate scanning electron micrographs of the microstructure of embodiments of steel compositions 2 and 3 in the as-quenched condition; (A) composition 2; (B) composition 3;
  • Figures 13A-13B illustrate scanning electron micrographs of the microstructure of embodiments of steel compositions 2 and 3 in the quenched and tempered condition; (A) composition 2; (B) composition 3;
  • Figure 14 is a plot of hardness as a function of tempering temperature for embodiments of steel compositions 2 (solid squares) and 3 (open squares);
  • Figure 15 is a plot of hardness as a function of the average cooling rate between 800°C and 500°C for an embodiment of steel composition 2 steel and a reference Nb-V steel.
  • Embodiments of the present disclosure present compositions and methods of manufacture for low carbon steels microalloyed with boron.
  • boron/titanium (B/Ti) steels which exhibit controlled particulates of titanium nitride (TiN), and attendant improvements in toughness, are discussed in detail.
  • TiN titanium nitride
  • free boron may be substantially kept in solid solution, improving hardenability during austenite decomposition.
  • the size of TiN precipitates may be controlled by the cooling rate during casting.
  • size may comprise the diameter of the precipitates.
  • size may comprise the largest dimension of the precipitates.
  • fine precipitates of TiN having a mean size less than about 50 nm, may be produced. Due to the small size of these TiN precipitates, they are not detrimental to toughness. Additionally, these precipitates may inhibit excessive grain growth during processing operations such as reheating prior to hot rolling. By reducing austenite grain size, toughness may be improved, after accelerated cooling or quenching, due to the reduction in martensite/bainite packet size.
  • the mechanical properties and microstructure of the steel composition may be further influenced by heat treatments after hot rolling.
  • steel compositions may be cooled in air at rates less than about 1°C/sec after hot rolling and subjected to reheating into the austenitic range and quenching.
  • steel compositions may be cooled in air after hot rolling and subjected to reheating into the austenitic range and quenching and tempering.
  • steel compositions may be subject to accelerated cooling at rates between about 5 to 50°C/sec directly after hot rolling.
  • compositions processed in this manner especially in the case of compositions subjected to quenching and tempering.
  • samples subjected to quenching and tempering at about 500°C may exhibit yield and tensile strengths of about 118 and 127 ksi, respectively, with impact energies measured in the range of about 143-173 J at about -60°C.
  • samples subjected to accelerated cooling may exhibit good impact energies, especially for cooling rates of about 10-20°C/sec.
  • impact energies greater than about 220 J are observed for temperatures of -20°C and higher.
  • Figure 1 illustrates one embodiment of a method 100 of producing boron-titanium (B/Ti) steels.
  • the compositions may be produced in the form of pipes.
  • the method 100 of Figure 1 includes steel casting operations in blocks 110, 112, and 114, collectively referred to as steel casting operations 102, steel forming operations in blocks 116, 120, 122 and 124, collectively referred to as steel forming operations 104, and steel heat treatment operations in blocks 126 and 128, collectively referred to as heat treatment operations 106. It may be appreciated that, in some embodiments one or more of the heat treatment operations can be omitted partially or totally, as necessary.
  • the B/Ti steel is cast from the molten state during steel casting operations 102.
  • the steel casting operations 102 may comprise continuous casting operations.
  • the steel casting operations 102 can include iron melting/purification 110, ladle treatments 112, and continuous casting 114, as are known in the art.
  • the steel may comprise elements in the concentration ranges listed below in Table 1, where the concentrations are provided in weight percent (wt. %) on the basis of the total weight of the steel composition, unless otherwise noted.
  • the cast steel may comprise a boron-titanium steel alloy including not only carbon (C), boron (B), and titanium (Ti) but one or more of manganese (Mn), silicon (Si), nickel (Ni), chromium (Cr) molybdenum (Mo), vanadium (V), and niobium (Nb).
  • Impurities of sulfur (S), phosphorous (P), copper (Cu), and nitrogen (N) may also be present, however, the concentration of these impurities in one embodiment is preferably reduced to an amount as low as possible.
  • C is an element whose addition inexpensively raises the strength of the steel. If the C content is less than about 0.04%, it may be, in some embodiments, difficult to obtain the strength desired in the composition. On the other hand, in other embodiments, if the steel has a C content greater than about 0.12 wt. %, toughness and weldability may be adversely impacted. Therefore, in an embodiment, the C content may range between about 0.04 to 0.12 wt. %. In other embodiments, the C content may range between about 0.04 to 0.08 wt. %. This lower C range may enable compositions to be fabricated, optionally, without tempering (i.e. in the as-quenched condition), while still achieving good toughness.
  • B is an element whose addition is effective in increasing the hardenability of the steel.
  • B may improve hardenability by inhibiting the formation of ferrite.
  • the B content is less than about 0.0005 wt. %, in some embodiments, it may be difficult to obtain the desired hardenability of the steel.
  • the concentration of B in the composition may range between about 0.0005 to 0.003 wt. %. In other embodiments, the concentration of B in the composition may range between about 0.0005 to 0.002 wt. %. At least a portion of the B in the composition may be in its free, elemental state in solid solution.
  • Si is an element whose addition has a deoxidizing effect during the steel making process and also raises the strength of the steel. If the Si content is too low, in some embodiments, the steel may be susceptible to oxidation, with a high level of micro-inclusions. On the other hand, though, if the Si content of the steel is too high, in some embodiments both toughness and formability of the steel may decrease. Therefore, in certain embodiments of the composition, the concentration of Si may range between about 0.05 to 0.3 wt. %.
  • Mn and Cr are elements which may be employed in combination with B, Mo, and Ni to increase hardenability.
  • these alloying additions may assist in inhibiting the formation of ferrite and pearlite from austenite during cooling. They may further enable depression of the bainitic start temperature, improving microstructural refinement.
  • Mn may additionally provide solid solution hardening.
  • the concentration of Mn may range between about 0.6 to 1.6 wt. %.
  • Cr may be omitted from the composition. In other embodiments, the concentration of Cr may range up to about 0.5 wt. %.
  • Mo is an element used to increase the hardenability of the steel composition. Alloying additions of Mo may also reduce the segregation of phosphorous to grain boundaries, improving resistance to inter-granular fracture. Mo may further enhance the hardenability effects of B. In certain embodiments, Mo may be omitted from the composition. In other embodiments, the concentration of Mo may range up to about 0.5 wt. %.
  • Ni is an alloying addition which may increase hardenability and improve toughness. In certain embodiments, Ni may be omitted from the composition. In other embodiments, the concentration ofNi may range up to about 0.5 wt. %.
  • Ti is an element whose addition is effective in increasing the effectiveness of B in the steel, by fixing nitrogen impurities as TiN and inhibiting the formation of boron nitrides. If the Ti content is too low it may be difficult, in some embodiments, to obtain the desired effect of boron on hardenability. In an embodiment, if the Ti content is higher than about 0.03 wt. %, coarse TiN and TiC may be formed, adversely affecting hot ductility and toughness. Accordingly, in certain embodiments, the concentration of Ti may range between about 0.01 to 0.03 wt. %.
  • the concentration of Ti may be specified on the basis of the concentration of N, maintaining a ratio of Ti to N greater than about 3.4 (for concentrations in weight percent).
  • substantially all of the N present within the composition may be in the form of TiN.
  • greater than about 90%, greater than about 92%, greater than about 94%, greater than about 96%, greater than about 98%, and greater than about 99% of the N content of the composition may be present in the form of TiN.
  • the TiN may adopt forms including, but not limited to, particles.
  • Nb is an alloying addition which may be used to refine the austenitic grain size of the composition. Nb may further enhance the effects of boron on hardenability and provide precipitation hardening. In certain embodiments, Nb may be omitted from the composition. In other embodiments, the concentration of Nb may range up to about 0.05 wt. %.
  • V is an alloying addition that may be employed to provide precipitation hardening. In certain embodiments, V may be omitted from the composition. In other embodiments, the concentration of V may range up to about 0.15 wt. %.
  • O is an impurity which may be present in the steel composition, for example, in the form of oxides. As the oxygen content increases, impact properties may be impaired. Accordingly, a lower oxygen content is preferred.
  • the upper limit of the oxygen content may be about 0.0050 wt. %. In another embodiment, the upper limit of oxygen content is below about 0.0015 wt. %.
  • Cu is not needed in embodiments of the steel composition, but may be present. In some embodiments, depending on the manufacturing process, the presence of Cu may be unavoidable. Thereafter, in an embodiment, the maximum Cu content may be about 0.10 wt. % or less.
  • S, P, Ca, N, and the like are impurities and their concentration is preferably kept as low as possible.
  • the concentration of each of S, P, Ca, and N may be independently provided as: S not greater than about 0.005 wt. %, P not greater than about 0.015 wt. %, Ca not greater than about 0.003 wt. %, and N not greater than about 0.008 wt. %.
  • the concentration of each of S, P, Ca, and N may be independently provided as: S not greater than about 0.003 wt. %, P not greater than about 0.015 wt. %, Ca not greater than about 0.002, and N not greater than about 0.006 wt. %.
  • the liquid steel may be continuously cast in steel casting operation 114.
  • the liquid steel may be cast into a rod, although, it may be understood that other shapes may be cast.
  • the cooling rate of the cast rod may be selected so as to provide control over the size of TiN precipitates that form during solidification.
  • the cooling rate during casting may be maintained at a selected rate.
  • the cooling rate may be selected such that the size of the TiN precipitates is less than about 50 nm.
  • the cooling rate from casting may be maintained at a rate greater than about 5°C/min at about the center of the rod.
  • the cooling rate from casting may be maintained at a rate greater than about 10°C/min at about the center of the rod. In other embodiments, the cooling rate from casting may be maintained at a rate greater than about 20°C/min at about the center of the rod. In additional embodiments, the cooling rate from casting may be maintained at a rate greater than about 30°C/min at about the center of the rod.
  • the rod thus fabricated may be subsequently formed into a tubular bar or pipe in steel forming operations 104, and more particularly may be formed into a seamless pipe.
  • a solid, substantially cylindrical rod of steel may be subjected to a first reheating operation (block 116) into the austenitic range, up to a temperature of about 1200°C to 1300°C, preferably about 1250°C.
  • the rod may be further pierced, in certain preferred embodiments, utilizing the Mannessmann process at temperatures between about 1100 to 1200°C, and subsequently hot rolled at temperatures ranging between about 900 to 1100°C.
  • the seamless hot rolled tube of steel may possess an approximately uniform wall thickness, both circumferentially around the tube and longitudinally along the tube axis.
  • tubes formed in this manner may possess an outer diameter ranging between about 60 to 273 mm and wall thickness ranging between about 6 to 25 mm.
  • a solid bar possessing an outer diameter of about 290 mm may be hot rolled in this manner into a tube possessing an outer diameter of about 244.5 mm and a wall thickness of about 16 mm.
  • the cross-sectional area reduction experienced by the tube may provide a refined microstructure.
  • a refined microstructure advantageously allows obtaining desired mechanical properties within the fabricated tube.
  • the seamless hot rolled tube of steel so manufactured may then be cooled to room temperature.
  • the austenitic grain size of the steel, after hot rolling and prior to transformation may range between about 10 to 50 ⁇ m. In other embodiments, the austenitic grain size of the steel, after hot rolling and prior to transformation, may range between about 20 to 50 ⁇ m.
  • this degree of austenitic refinement may allow selected compositions to achieve a good balance of strength and toughness after accelerated cooling from the finish rolling temperature without the need for subsequent heat treatment, such as quenching or quenching and tempering.
  • this degree of austenitic refinement may allow compositions having elevated carbon concentrations to achieve a good balance of strength and toughness when subjected to heat treatments such as quenching and tempering.
  • Embodiments of the composition may be cooled from hot rolling by air cooling or accelerated cooling in block 124.
  • cooling rates less than about 1°C/sec may be achieved in tubes with wall thickness greater than about 8 mm.
  • Subsequent heat treatments may also be employed to improve the strength and toughness of the steel composition.
  • cooling may be performed directly from hot rolling, without an intermediate cooling step, to room temperature (block 124).
  • Several devices can be used to achieve cooling rates greater than that corresponding to natural air cooling, including, but not limited to, forced air flow, water sprays, and air-water mixture sprays.
  • the flow of the coolant may be directed to the outer tube wall, or to the inner and outer tube walls in order to improve microstructure homogeneity.
  • cooling rates ranging between about 5 to 50°C/sec may be achieved through accelerated cooling.
  • cooling rates ranging between about 10 to 50°C/sec may be employed.
  • cooling rates ranging between about 10 to 20°C/sec may be employed.
  • these cooling rates may be employed with tubes of wall thickness between about 8 mm and 25 mm.
  • Forming of the hot rolled tube may be completed through a plurality of finishing steps.
  • finishing steps may include cutting the tube to length, such as lengths of approximately 8 m to 15 m, cropping the ends of the tube, straightening the tube, and non-destructive testing (e.g., electromagnetic testing, ultrasound testing).
  • non-destructive testing e.g., electromagnetic testing, ultrasound testing.
  • One or more heat treatment operations 106 may optionally be performed upon the tube after the forming operations 104.
  • quenching may be performed in block 126.
  • the composition may be reheated a second time into the austenitic range, prior to quenching, to temperatures greater than about Ae3 (e.g., about 870-950°C). Soak times at maximum temperature may range between about 5 to 30 minutes. Quenching may be further performed with water sprays to cool the composition from about the maximum temperature to about room temperature.
  • tempering may be further performed upon quenched compositions in block 128.
  • Tempering may be performed by heating to temperatures ranging between about 400 to 700°C, holding at the tempering temperature for a selected duration, and air cooling from the tempering temperature to about room temperature. The compositions may be held at the tempering temperature for between about 10 to 60 minutes.
  • compositions 1, 2, and 3 The manufacture, microstructure, and mechanical properties of embodiments of three steel compositions of the present disclosure, referred to as compositions 1, 2, and 3, are discussed in the examples below. The performance benefits achieved from such compositions are further discussed. It may be understood that these examples are discussed for illustrative purposes and should not be construed to limit the scope of the disclosed embodiments.
  • compositions 1, 2, and 3 The concentrations of alloying elements present in compositions 1, 2, and 3 are illustrated below in Table 2.
  • composition 1 was designed to produce a fine bainitic structure after accelerated cooling from the austenitic range.
  • composition 2 having the highest carbon content, was designed for use with air cooling from hot rolling, followed by quenching and tempering, as discussed below. Due to its higher carbon content, Cr and Nb are substantially absent in composition 2, as compared with the other compositions. Further, composition 3, having the lowest carbon content, was designed to obtain high strength and good toughness in the as-quenched condition, without tempering.
  • compositions 1, 2 and 3 were melted in an approximately 20 kg vacuum induction furnace and electro-slag re-melted to decrease sulfur content. Subsequently, compositions 1, 2, and 3 were cast into slabs having a thickness of approximately 140 mm and hot rolled to a final thickness of about 16 mm. During hot rolling, reheating and finishing temperatures of about 1200-1250°C and 950-1000°C, respectively, were employed. All the hot-rolled plates were subsequently air cooled to about room temperature.
  • the hot rolled composition 1 was subjected to one of the following post-rolling processing sequences:
  • Quenching was performed in water, from a temperature of about 900-950°C to about room temperature, using moderate agitation. Where tempering operations were also performed, the composition was heated to between about 300°C to 450°C with a soaking time of about one hour.
  • Accelerated cooling was performed by cooling the composition in a mixture of air and water from the reheating temperature of about 900-950°C, to about room temperature, at cooling rates ranging between about 5 to 45°C/sec.
  • the reheating temperature prior to accelerated cooling was selected to have an austenitic microstructure representative of that industrially achieved just at the exit of the hot rolling mill.
  • the complete heat treatment was performed using a Gleeble 3500 thermo-mechanical simulator.
  • Hot rolled compositions 2 and 3 were subject to one of two post-rolling processing sequences:
  • Quenching was performed by heating the composition to a temperature of about 925°C (composition 2) or about 930°C (composition 3), with a soak time of about 10 min. Cooling was performed in water from the quenching temperature to about room temperature using moderate agitation. When tempering was performed, the composition was heated to temperatures ranging between about 400 to 700°C, with a soak time at maximum temperature of about 30 minutes.
  • compositions 1, 2, and 3 were reheated at about 5°C/sec to about 920°C, 925°C, and 930°C, respectively, with a soak time of about 10 min at maximum temperature.
  • the austenization temperatures were chosen to be approximately 20-30°C above the Ac3 temperature corresponding to the respective compositions. Cooling rates ranging between about 0.5 to 50°C/sec were examined in composition I and cooling rates ranging between about 0.2 to 50°C/sec were examined in compositions 2 and 3.
  • the resulting microstructures were further characterized using optical and scanning electron microscopy.
  • the mechanical properties of the compositions so fabricated were further evaluated by mechanical tests including one or more of tensile testing, hardness testing, and Charpy testing. In each case, tensile samples and full size Charpy samples were taken in the transversal direction. Tensile testing was performed in accordance with ASTM E8, "Standard Test Methods for Tension Testing of Metallic Materials", the entirety of which is incorporated herein by reference, and the reported results are averaged over two samples.
  • Example 1 Continuous Cooling Transformation (CCT) behavior and microstructural evaluation of Composition 1 for cooling rates between about 0.5 and 50°C / s.
  • the CCT diagram derived from dilatometric measurements of composition 1 is shown in Figure 2 . Illustrated in Figure 2 are traces of temperature as a function of cooling rate for transformations of about 5%, 20%, 50%, 80%, and 95%. Due to the reheating condition, about 920°C over about 10 min, the austenitic grain size prior to transformation was estimated to be about 10-20 ⁇ m, based upon the sample cooled at about 50°C/sec.
  • the microstructure observed under these conditions was again mainly bainitic, as illustrated in the micrographs corresponding to cooling rates of about 10°C/sec and 20°C/sec in Figure 3 . Notably, however, the bainitic structure was finer and substantially without the blocky regions of retained austenite.
  • Hardness measurements for composition 1 after cooling at different rates are also illustrated in CCT diagram of Figure 2 . It may be observed that the hardness ranges between about 262 Hv for cooling rates of about 2°C/sec to greater than about 340 Hv for cooling rates of about 50°C.
  • Example 2 Toughness Evaluation of Composition 1 - Accelerated Cooling Condition
  • composition 1 in the as-quenched condition The tensile and impact properties of composition 1 in the as-quenched condition are illustrated in Tables 4 and 5.
  • Table 4 Tensile properties of composition 1 after quenching Composition Condition YS (ksi) (ksi) UTS YS/UT S El (%) 1 As-Quenched 121 156 0.78 16
  • Table 5 Impact Energy and Hardness of composition 1 after quenching Composition Condition Hardness Charpy (10 x 10 mm) (Hv 1 kg) T (°C) CVN (J) Shear Area (%) 1 As-Quenched 363 25 150 79 0 82 43 -20 42 24
  • the as-quenched composition exhibited improvements in strength and impact energy over that of as-rolled samples (YS ⁇ 69 ksi, UTS ⁇ 99 ksi, CVN-6-8 J at 25°C to -20°C). This improvement may be ascribed to a general refinement of the microstructure and the substantial disappearance of large, blocky austenitic regions.
  • composition 1 in the quenched and tempered condition samples were quenched as discussed above and tempered at temperatures ranging between about 350°C to 440°C for about 1 hr.
  • the measured hardness values are illustrated in Figure 5 .
  • the hardness in the as-quenched condition is about 362 Hv, falling modestly with tempering at about 300 to 400°C to within about 350 to 335 Hv.
  • Samples tempered at about 440°C further exhibited a significant decrease in hardness, falling to about 280 +/- 20 Hv.
  • composition 1 Composition Condition (°C) Temper Yield Strength (ksi) UTS (ksi) YS/UTS El 1 As-Quenched N/A 121 156 0.78 16 1 As-Quenched and Tempered 410 129 138 0.94 14 1 Quenched and Tempered 440 129 141 0.91 16
  • Table 7 Impact Energy of Composition 1 Composition Condition Temper (°C) Impact Energy T(°C) CVN (J) Shear Area (%) 1 As-Quenched N/A 24 150 79 0 82 43 -20 42 24 1 Quenched and Tempered 410 24 215 100 0 171 74 -20 136 59 1 Quenched and Tempered 440 24 170 84 0 144 76 -20 113 49
  • the yield and tensile strengths measured in the quenched and tempered condition were about 129 ksi and about 138-141 ksi, respectively.
  • the yield strength measured in the as-quenched material was less, about 121 ksi, while the tensile strength was greater, about 156 ksi.
  • the impact energies of samples in the quenched and tempered condition were found to be greater than those of samples measured at comparable temperatures in the as-quenched condition.
  • samples tempered at 410 and 440°C exhibited impact energies of about 215 and 170 J, respectively, while the impact energy of the as-quenched material was about 150 J.
  • the difference in impact energies was even greater, with samples tempered at about 410 and 440°C exhibited impact energies of about 136 and 113 J, respectively, while the impact energy of the as-quenched material was about 42 J.
  • the microstructure of composition 1 is bainite and martensite, with a fine dispersion of carbides, which improves the yield strength of the quenched and tempered material over that of the as-quenched material alone.
  • impact energy and hardness values were also observed. Most notably, in the range of about 10 to 20°C/s, impact energy values greater than about 220 J at about -20°C were achieved with more than about 80% ductile area. Furthermore, hardness values ranged between about 300-320 Hv.
  • the CCT diagrams derived from dilatometric measurements of compositions 2 and 3 are shown in Figures 7 and 9 , respectively, for cooling rates of about 0.2, 0.5, 5, 10, 30, and 50°C/sec.
  • the transformation start temperatures shown in these figures were determined as the first deviation from linear behavior of both dilatometric curves.
  • the austenitic grain sizes of compositions 2 and 3 were estimated to be between about 20 to 30 ⁇ m from measurements on samples cooled at about 50°C/sec.
  • Figures 8 and 10 further illustrate optical micrographs of compositions 2 and 3 cooled at rates of about 0.2, 0.5, 1, 10, 30, and 50°C/sec.
  • compositions 2 and 3 From the measured CCT diagrams and the observed microstructures, the transformation behavior of compositions 2 and 3 may be identified.
  • Bainite is the main transformation product when cooling between about 5°C/sec to 30°C/sec. At lower cooling rates, polygonal ferrite is the predominant constituent. Martensite appears in composition 2 at cooling rates of about 10°C/sec and in composition 3 at about 30°C/sec and becomes the dominant phase when cooling at about 50°C/sec in both compositions.
  • composition 2 A number of notable differences between the two compositions may also be observed.
  • pearlite is found in large quantities, in addition to bainite.
  • composition 3 a more complex microstructure was observed, with a higher portion of bainite and some retained austenite, in addition to pearlite. Without being bound by theory, this difference may be ascribed to the lower carbon content of composition 3, which reduces the total fraction of pearlite, as well as the alloying additions of Cr and Nb, which encourage bainite formation.
  • the scale of the bainite differs between compositions 2 and 3.
  • the bainitic structure of composition 3 is generally finer than that of composition 2. Without being bound by theory, this observation is believed to be a consequence of the Cr and Nb alloying additions.
  • the tendency to form martensite is stronger in composition 2.
  • the lowest cooling rate at which martensite was observed in composition 2 was about 10°C/sec, while the lowest cooling rate at which martensite was observed in composition 3 was about 30°C/sec.
  • the concentration of martensite in composition 2 is similar or higher to that of bainite.
  • compositions 2 and 3 values as a function of cooling rate are also illustrated in Figures 11A and 11B .
  • Calculations performed with Creusot-Loire modeling (see Ph. Maynier, B. Jungmann, and J. Dollet, "Creusot-Loire system for the prediction of the mechanical properties of low alloy steel products", Hardenability concepts with applications to steels, Ed. D.V. Doane and J. S. Kirkaldy, The Metallurgical Society of AIME (1978), p. 518 ) are presented in the same graphs for comparison (dashed line). It is notable that, despite its lower carbon content, composition 3 presents a slightly higher hardness level than that of composition 2 for cooling rates below about 30°C/sec.
  • this increment in hardness may be ascribed to the microstructural refinement of composition 3 already mentioned. It is also possible that, at the lower cooling rates in composition 3, some Nb dissolved during the austenization stage re-precipitates as fine carbides, which may increase hardness.
  • composition 3 The tensile and impact properties measured for composition 3 in the as-quenched condition are presented in Tables 8 and 9 below. Hardness properties of composition 2 are also presented in Table 9. Corresponding SEM micrographs for the compositions are illustrated in Figures 12A and 12B . These results illustrate the effect of carbon on the microstructure and mechanical properties of the compositions.
  • Table 8 Tensile properties of as-auenched composition 3 Composition Condition YS (ksi) UTS (ksi) YS/UTS El (%) 3 As-Quenched 121 148 0.82 14.5 Table 9 - Impact Energy and Hardness of as-quenched compositions 2 and 3 Composition Condition Hardness Charpy (10 x 10 mm) (Hv 1Kg) T(°C) CVN (J) Shear Area (%) 2 As-Quenched 350 - - - 3 As-Quenched 339 20 163 100 0 161 100 -20 162 97 -40 96 47 -60 67 27
  • the microstructure of as-quenched composition 2 was primarily martensite, with some regions of bainite ( Figure 12A ). Further, the hardness of composition 2 was relatively high, about 350 Hv. Based upon experience with other systems, poor toughness was expected for this system and no tensile or impact tests were performed.
  • the as-quenched microstructure of composition 3 was predominantly bainitic, with small regions of martensite ( Figure 12B ).
  • tensile and impact tests were carried out and a beneficial combination of properties was obtained.
  • the yield strength was measured to be approximately 121 ksi, with a low yield strength to tensile strength ratio, about 0.82.
  • the ductile to brittle transition temperature measured as that corresponding to about 50% shear area, was found to be about -40°C.
  • the impact energy was measured to be substantially constant at about 160J between about -20°C to 20°C.
  • compositions 2 and 3 were found to yield beneficial properties in case of composition 3.
  • similar testing and evaluation were performed on samples of compositions 2 and 3 in the quenched and tempered condition.
  • Figures 13A and 13B present scanning electron micrographs of the microstructure of compositions 2 and 3 in the quenched and tempered condition.
  • the microstructure was composed mainly of slightly tempered bainite.
  • Hardness results from small samples of compositions 2 and 3 tempered between about 400 to 700°C are illustrated in Figure 14 . It may be observed that the response of both compositions exhibit similar evolution in hardness with increasing tempering temperature. As expected, owing to its greater carbon content, composition 2 was found to exhibit greater hardness than composition 3 in the as-quenched condition and at low tempering temperatures. Conversely, however, for tempering temperatures greater than about 550°C, the hardness of composition 3 was found to be greater than that of composition 2.
  • compositions 2 and 3 The mechanical properties after tempering at about 500°C for about 30 min are found to exhibit a good combination of strength and toughness in both compositions 2 and 3.
  • the yield strengths of compositions 2 and 3 are about 118 ksi and the ultimate tensile strengths are about 126-127 ksi.
  • compositions 2 and 3 further exhibit ductile to brittle transition temperatures below about -60°C, with nearly 100% of shear area over the temperature range examined.
  • compositions 2 and 3 exhibit nearly identical tensile and impact energy properties. Without being bound by theory, in comparing the chemistries of the two compositions, it appears that the reduction in carbon content in composition 3 is approximately offset by the alloying additions of Cr and Nb.
  • compositions 2 and 3 which were air-cooled after hot rolling, then reheated and quenched, exhibited good toughness when carbon content was maintained below about 0.07% (composition 3). Furthermore, good combinations of strength and toughness were obtained when quenching and tempering at temperatures of about 500°C. In this case, excellent mechanical properties were obtained for both compositions 2 and 3, with yield strengths of about 118 ksi and impact energies of about 175-179 J at about -40°C. Additionally, almost fully ductile fracture surfaces were observed for the range of testing temperatures studied, with the ductile to brittle transition temperature well below about -60°C for both materials.
  • low carbon steels having alloying additions of boron and titanium having alloying additions of boron and titanium are presented. Free nitrogen impurities are substantially consumed by reaction with titanium, forming TiN precipitates. Casting parameters are further selected so as to inhibit these precipitates from coarsening. For example, by employing cooling rates greater than about 10 to 40°C/min during casting, fine precipitates of TiN having a mean diameter less than about 50 nm may be achieved. Substantial removal of free nitrogen impurities further allows free boron to remain in solid solution, improving hardenability during austenite decomposition. These compositions may be cooled from hot rolling in air and quenched, quenching and tempering, or subjected to accelerated cooling directly after hot rolling at rates between about 5 to 50°C, yielding an excellent balance of strength and toughness.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Mechanical Engineering (AREA)
  • Thermal Sciences (AREA)
  • Physics & Mathematics (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Materials Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Manufacturing & Machinery (AREA)
  • Heat Treatment Of Steel (AREA)
  • Heat Treatment Of Articles (AREA)
  • Heat Treatment Of Sheet Steel (AREA)
EP10166261.7A 2009-06-17 2010-06-17 Bainitstähle mit Bor Active EP2287346B1 (de)

Applications Claiming Priority (1)

Application Number Priority Date Filing Date Title
US12/486,610 US20100319814A1 (en) 2009-06-17 2009-06-17 Bainitic steels with boron

Publications (2)

Publication Number Publication Date
EP2287346A1 true EP2287346A1 (de) 2011-02-23
EP2287346B1 EP2287346B1 (de) 2019-12-18

Family

ID=43086438

Family Applications (1)

Application Number Title Priority Date Filing Date
EP10166261.7A Active EP2287346B1 (de) 2009-06-17 2010-06-17 Bainitstähle mit Bor

Country Status (6)

Country Link
US (1) US20100319814A1 (de)
EP (1) EP2287346B1 (de)
JP (1) JP5787492B2 (de)
AR (1) AR077129A1 (de)
BR (1) BRPI1004267B1 (de)
MX (1) MX2010006761A (de)

Cited By (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN103451550A (zh) * 2012-06-01 2013-12-18 北京奇峰聚能科技有限公司 储能飞轮铸件用合金钢及储能飞轮铸件铸造方法
CN111876696A (zh) * 2020-07-23 2020-11-03 江阴兴澄特种钢铁有限公司 一种服役温度可达-60℃以下的x100管件用钢板及其制造方法
WO2024008920A1 (en) * 2022-07-08 2024-01-11 Tenaris Connections B.V. Steel composition for expandable tubular articles, expandable tubular article having this steel composition, manufacturing method thereof and use thereof

Families Citing this family (31)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2004097059A1 (es) 2003-04-25 2004-11-11 Tubos De Acero De Mexico, S.A. Tubo de acero sin costura para ser utilizado como canalizador y proceso de obtencíon del mismo
MXPA05008339A (es) * 2005-08-04 2007-02-05 Tenaris Connections Ag Acero de alta resistencia para tubos de acero soldables y sin costura.
US7744708B2 (en) * 2006-03-14 2010-06-29 Tenaris Connections Limited Methods of producing high-strength metal tubular bars possessing improved cold formability
WO2008000300A1 (en) * 2006-06-29 2008-01-03 Tenaris Connections Ag Seamless precision steel tubes with improved isotropic toughness at low temperature for hydraulic cylinders and process for obtaining the same
MX2007004600A (es) * 2007-04-17 2008-12-01 Tubos De Acero De Mexico S A Un tubo sin costura para la aplicación como secciones verticales de work-over.
US7862667B2 (en) 2007-07-06 2011-01-04 Tenaris Connections Limited Steels for sour service environments
EP2238272B1 (de) * 2007-11-19 2019-03-06 Tenaris Connections B.V. Hochfester bainitischer stahl für octg-anwendungen
CA2686301C (en) * 2008-11-25 2017-02-28 Maverick Tube, Llc Compact strip or thin slab processing of boron/titanium steels
EP2325435B2 (de) 2009-11-24 2020-09-30 Tenaris Connections B.V. Verschraubung für [ultrahoch] abgedichteten internen und externen Druck
US9163296B2 (en) 2011-01-25 2015-10-20 Tenaris Coiled Tubes, Llc Coiled tube with varying mechanical properties for superior performance and methods to produce the same by a continuous heat treatment
IT1403689B1 (it) 2011-02-07 2013-10-31 Dalmine Spa Tubi in acciaio ad alta resistenza con eccellente durezza a bassa temperatura e resistenza alla corrosione sotto tensioni da solfuri.
IT1403688B1 (it) 2011-02-07 2013-10-31 Dalmine Spa Tubi in acciaio con pareti spesse con eccellente durezza a bassa temperatura e resistenza alla corrosione sotto tensione da solfuri.
US8636856B2 (en) 2011-02-18 2014-01-28 Siderca S.A.I.C. High strength steel having good toughness
US8414715B2 (en) 2011-02-18 2013-04-09 Siderca S.A.I.C. Method of making ultra high strength steel having good toughness
US9403242B2 (en) * 2011-03-24 2016-08-02 Nippon Steel & Sumitomo Metal Corporation Steel for welding
US9340847B2 (en) 2012-04-10 2016-05-17 Tenaris Connections Limited Methods of manufacturing steel tubes for drilling rods with improved mechanical properties, and rods made by the same
CA2897451C (en) 2013-01-11 2019-10-01 Tenaris Connections Limited Galling resistant drill pipe tool joint and corresponding drill pipe
US9187811B2 (en) 2013-03-11 2015-11-17 Tenaris Connections Limited Low-carbon chromium steel having reduced vanadium and high corrosion resistance, and methods of manufacturing
US9803256B2 (en) 2013-03-14 2017-10-31 Tenaris Coiled Tubes, Llc High performance material for coiled tubing applications and the method of producing the same
EP2789700A1 (de) 2013-04-08 2014-10-15 DALMINE S.p.A. Dickwandige vergütete und nahtlose Stahlrohre und entsprechendes Verfahren zur Herstellung der Stahlrohre
EP2789701A1 (de) 2013-04-08 2014-10-15 DALMINE S.p.A. Hochfeste mittelwandige vergütete und nahtlose Stahlrohre und entsprechendes Verfahren zur Herstellung der Stahlrohre
KR102368928B1 (ko) 2013-06-25 2022-03-04 테나리스 커넥션즈 비.브이. 고크롬 내열철강
US9850553B2 (en) 2014-07-22 2017-12-26 Roll Forming Corporation System and method for producing a hardened and tempered structural member
BR102016001063B1 (pt) * 2016-01-18 2021-06-08 Amsted Maxion Fundição E Equipamentos Ferroviários S/A liga de aço para componentes ferroviários, e processo de obtenção de uma liga de aço para componentes ferroviários
US11124852B2 (en) 2016-08-12 2021-09-21 Tenaris Coiled Tubes, Llc Method and system for manufacturing coiled tubing
KR101777974B1 (ko) * 2016-08-23 2017-09-12 현대제철 주식회사 철근 및 그 제조 방법
US10434554B2 (en) 2017-01-17 2019-10-08 Forum Us, Inc. Method of manufacturing a coiled tubing string
US11701539B2 (en) * 2019-02-22 2023-07-18 Jaquish Biomedical Corporation Variable resistance exercise devices
CN110106445B (zh) * 2019-06-05 2021-04-16 上海大学 一种用于海洋平台铸造节点高强度高低温韧性用钢及其制备方法
EP4028563A4 (de) 2019-09-19 2022-07-27 Nucor Corporation Extrem wetterfester stahl für heissprägeanwendungen
CN118234546A (zh) 2021-10-06 2024-06-21 贾奎什生物医学公司 提供感觉反馈的健身杆

Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4812182A (en) * 1987-07-31 1989-03-14 Hongsheng Fang Air-cooling low-carbon bainitic steel
EP1008660A1 (de) * 1998-12-09 2000-06-14 Sumitomo Metal Industries Limited Niedrig legierter Stahl für ölindustrielle röhrenförmige Gegenstände
EP1914324A1 (de) * 2005-07-25 2008-04-23 Sumitomo Metal Industries, Ltd. Verfahren zur herstellung eines nahtlosen stahlrohrs
US20090010794A1 (en) * 2007-07-06 2009-01-08 Gustavo Lopez Turconi Steels for sour service environments
WO2009065432A1 (en) * 2007-11-19 2009-05-28 Tenaris Connections Ag High strength bainitic steel for octg applications

Family Cites Families (24)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3413166A (en) * 1965-10-15 1968-11-26 Atomic Energy Commission Usa Fine grained steel and process for preparation thereof
US3655465A (en) * 1969-03-10 1972-04-11 Int Nickel Co Heat treatment for alloys particularly steels to be used in sour well service
US3915697A (en) * 1975-01-31 1975-10-28 Centro Speriment Metallurg Bainitic steel resistant to hydrogen embrittlement
JPH07331381A (ja) * 1994-06-06 1995-12-19 Nippon Steel Corp 高強度高靭性継目無鋼管およびその製造法
GB2297094B (en) * 1995-01-20 1998-09-23 British Steel Plc Improvements in and relating to Carbide-Free Bainitic Steels
CA2231985C (en) * 1997-03-26 2004-05-25 Sumitomo Metal Industries, Ltd. Welded high-strength steel structures and methods of manufacturing the same
US5993570A (en) * 1997-06-20 1999-11-30 American Cast Iron Pipe Company Linepipe and structural steel produced by high speed continuous casting
US6866725B2 (en) * 2000-02-28 2005-03-15 Nippon Steel Corporation Steel pipe excellent in formability and method of producing the same
WO2002063058A1 (fr) * 2001-02-07 2002-08-15 Nkk Corporation Feuille d"acier mince et procede de fabrication de ladite feuille
US6669789B1 (en) * 2001-08-31 2003-12-30 Nucor Corporation Method for producing titanium-bearing microalloyed high-strength low-alloy steel
US6669285B1 (en) * 2002-07-02 2003-12-30 Eric Park Headrest mounted video display
WO2004097059A1 (es) * 2003-04-25 2004-11-11 Tubos De Acero De Mexico, S.A. Tubo de acero sin costura para ser utilizado como canalizador y proceso de obtencíon del mismo
AR047467A1 (es) * 2004-01-30 2006-01-18 Sumitomo Metal Ind Tubo de acero sin costura para pozos petroliferos y procedimiento para fabricarlo
US7566416B2 (en) * 2004-10-29 2009-07-28 Sumitomo Metal Industries, Ltd. Steel pipe for an airbag inflator and a process for its manufacture
JP4792778B2 (ja) * 2005-03-29 2011-10-12 住友金属工業株式会社 ラインパイプ用厚肉継目無鋼管の製造方法
MXPA05008339A (es) * 2005-08-04 2007-02-05 Tenaris Connections Ag Acero de alta resistencia para tubos de acero soldables y sin costura.
US7744708B2 (en) * 2006-03-14 2010-06-29 Tenaris Connections Limited Methods of producing high-strength metal tubular bars possessing improved cold formability
JP4751224B2 (ja) * 2006-03-28 2011-08-17 新日本製鐵株式会社 靭性と溶接性に優れた機械構造用高強度シームレス鋼管およびその製造方法
US20080226396A1 (en) * 2007-03-15 2008-09-18 Tubos De Acero De Mexico S.A. Seamless steel tube for use as a steel catenary riser in the touch down zone
CN101514433A (zh) * 2007-03-16 2009-08-26 株式会社神户制钢所 低温冲击特性优异的汽车用高强度电阻焊钢管及其制造方法
MX2007004600A (es) * 2007-04-17 2008-12-01 Tubos De Acero De Mexico S A Un tubo sin costura para la aplicación como secciones verticales de work-over.
JP5020690B2 (ja) * 2007-04-18 2012-09-05 新日本製鐵株式会社 機械構造用高強度鋼管及びその製造方法
US8414715B2 (en) * 2011-02-18 2013-04-09 Siderca S.A.I.C. Method of making ultra high strength steel having good toughness
US8636856B2 (en) * 2011-02-18 2014-01-28 Siderca S.A.I.C. High strength steel having good toughness

Patent Citations (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4812182A (en) * 1987-07-31 1989-03-14 Hongsheng Fang Air-cooling low-carbon bainitic steel
EP1008660A1 (de) * 1998-12-09 2000-06-14 Sumitomo Metal Industries Limited Niedrig legierter Stahl für ölindustrielle röhrenförmige Gegenstände
EP1914324A1 (de) * 2005-07-25 2008-04-23 Sumitomo Metal Industries, Ltd. Verfahren zur herstellung eines nahtlosen stahlrohrs
US20090010794A1 (en) * 2007-07-06 2009-01-08 Gustavo Lopez Turconi Steels for sour service environments
WO2009065432A1 (en) * 2007-11-19 2009-05-28 Tenaris Connections Ag High strength bainitic steel for octg applications

Non-Patent Citations (3)

* Cited by examiner, † Cited by third party
Title
"Corrosion 2001", part Paper 01077 16 March 2001, NACE INTERNATIONAL, Houston, Texas, article G.L.TURCONI ET AL.: "Improvement of resistance to SSC initiation and propagation of high strength OCTG through microstructure and precipitation control", pages: 1 - 15, XP009141583 *
G. GOMEZ ET AL: "Air cooled bainitic steels for strong, seamless pipes - Part 1 -alloy design, kinetics and microstructure", MATERIALS SCIENCE AND TECHNOLOGY, vol. 25, no. 12, 1 December 2009 (2009-12-01), pages 1501 - 1507, XP002611498 *
HONG-SHENG FANG ET AL.: "The Developing Prospect of Air-cooled Bainitic Steels", INTERNATIONAL JOURNAL OF ISSI, vol. 2, no. 2, 1 February 2005 (2005-02-01), pages 9 - 18, XP002611497 *

Cited By (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN103451550A (zh) * 2012-06-01 2013-12-18 北京奇峰聚能科技有限公司 储能飞轮铸件用合金钢及储能飞轮铸件铸造方法
CN103451550B (zh) * 2012-06-01 2015-07-15 北京奇峰聚能科技有限公司 储能飞轮铸件用合金钢及储能飞轮铸件铸造方法
CN111876696A (zh) * 2020-07-23 2020-11-03 江阴兴澄特种钢铁有限公司 一种服役温度可达-60℃以下的x100管件用钢板及其制造方法
CN111876696B (zh) * 2020-07-23 2021-08-24 江阴兴澄特种钢铁有限公司 一种服役温度可达-60℃以下的x100管件用钢板及其制造方法
WO2024008920A1 (en) * 2022-07-08 2024-01-11 Tenaris Connections B.V. Steel composition for expandable tubular articles, expandable tubular article having this steel composition, manufacturing method thereof and use thereof
NL2032426B1 (en) * 2022-07-08 2024-01-23 Tenaris Connections Bv Steel composition for expandable tubular products, expandable tubular article having this steel composition, manufacturing method thereof and use thereof

Also Published As

Publication number Publication date
BRPI1004267A2 (pt) 2012-03-20
US20100319814A1 (en) 2010-12-23
JP2011006790A (ja) 2011-01-13
BRPI1004267B1 (pt) 2020-12-22
MX2010006761A (es) 2010-12-22
JP5787492B2 (ja) 2015-09-30
AR077129A1 (es) 2011-08-03
EP2287346B1 (de) 2019-12-18

Similar Documents

Publication Publication Date Title
EP2287346B1 (de) Bainitstähle mit Bor
DK2495341T3 (en) High strength steel and good hardness
KR101222724B1 (ko) 연성이 우수한 고강도 강 시트의 제조 방법 및 그 제조방법에 의해 제조된 시트
US8414715B2 (en) Method of making ultra high strength steel having good toughness
KR20180099876A (ko) 고강도 강판 및 그 제조 방법
KR20080084748A (ko) 저온 충격 특성이 우수한 자동차용 고강도 전봉 강관과 그제조 방법
EP2559783A1 (de) Hochfeste heissgewalzte stahlplatte mit hervorragenden streckbarkeits- und ermüdungseigenschaften sowie herstellungsverfahren dafür
KR101908819B1 (ko) 저온에서의 파괴 개시 및 전파 저항성이 우수한 고강도 강재 및 그 제조방법
JP2010168624A (ja) 高周波焼入れ用圧延鋼材およびその製造方法
KR101908818B1 (ko) 저온에서의 파괴 개시 및 전파 저항성이 우수한 고강도 강재 및 그 제조방법
JP2020500262A (ja) 低温用中マンガン鋼材及びその製造方法
JP2010144226A (ja) 高周波焼入れ用圧延鋼材およびその製造方法
JP4547944B2 (ja) 高強度高靭性厚鋼板の製造方法
JP5459064B2 (ja) 高周波焼入れ用圧延鋼材およびその製造方法
EP3269837B1 (de) Mikro-legierter stahl und verfahren zur herstellung dieses stahls
JP5459063B2 (ja) 高周波焼入れ用圧延鋼材およびその製造方法
JP2023531248A (ja) 鋼組成物から高強度鋼管を製造する方法およびその鋼管から作られる構成部品
JP5459062B2 (ja) 高周波焼入れ用圧延鋼材およびその製造方法
EP3722449B1 (de) Warmgewalztes stahlblech für gewickelte rohre und verfahren zu dessen herstellung
JP5459065B2 (ja) 高周波焼入れ用圧延鋼材およびその製造方法
JP5194571B2 (ja) 引張強さ５７０Ｎ／ｍｍ２級以上の溶接割れ感受性に優れた高張力鋼の製造方法
JP2005120397A (ja) 絞り特性に優れた高強度鍛造部品
KR100398375B1 (ko) 고주파소입용비조질강선재의제조방법

Legal Events

Date Code Title Description
PUAI Public reference made under article 153(3) epc to a published international application that has entered the european phase

Free format text: ORIGINAL CODE: 0009012

AK Designated contracting states

Kind code of ref document: A1

Designated state(s): AL AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HR HU IE IS IT LI LT LU LV MC MK MT NL NO PL PT RO SE SI SK SM TR

AX Request for extension of the european patent

Extension state: BA ME RS

17P Request for examination filed

Effective date: 20110601

17Q First examination report despatched

Effective date: 20160229

RAP1 Party data changed (applicant data changed or rights of an application transferred)

Owner name: TENARIS CONNECTIONS B.V.

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: EXAMINATION IS IN PROGRESS

GRAP Despatch of communication of intention to grant a patent

Free format text: ORIGINAL CODE: EPIDOSNIGR1

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: GRANT OF PATENT IS INTENDED

RIN1 Information on inventor provided before grant (corrected)

Inventor name: GONZALO ROBERTO, GOMEZ

Inventor name: TERESA ESTELA, PEREZ

INTG Intention to grant announced

Effective date: 20190711

GRAS Grant fee paid

Free format text: ORIGINAL CODE: EPIDOSNIGR3

GRAA (expected) grant

Free format text: ORIGINAL CODE: 0009210

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: THE PATENT HAS BEEN GRANTED

AK Designated contracting states

Kind code of ref document: B1

Designated state(s): AL AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HR HU IE IS IT LI LT LU LV MC MK MT NL NO PL PT RO SE SI SK SM TR

REG Reference to a national code

Ref country code: GB

Ref legal event code: FG4D

REG Reference to a national code

Ref country code: CH

Ref legal event code: EP

REG Reference to a national code

Ref country code: DE

Ref legal event code: R096

Ref document number: 602010062400

Country of ref document: DE

REG Reference to a national code

Ref country code: IE

Ref legal event code: FG4D

REG Reference to a national code

Ref country code: AT

Ref legal event code: REF

Ref document number: 1214661

Country of ref document: AT

Kind code of ref document: T

Effective date: 20200115

REG Reference to a national code

Ref country code: NL

Ref legal event code: MP

Effective date: 20191218

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: LV

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: SE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: BG

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200318

Ref country code: NO

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200318

Ref country code: LT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: GR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200319

Ref country code: FI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

REG Reference to a national code

Ref country code: LT

Ref legal event code: MG4D

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: HR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: AL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: EE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: PT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200513

Ref country code: CZ

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: RO

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: NL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: SM

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: SK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: IS

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200418

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 602010062400

Country of ref document: DE

REG Reference to a national code

Ref country code: AT

Ref legal event code: UEP

Ref document number: 1214661

Country of ref document: AT

Kind code of ref document: T

Effective date: 20191218

PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: NO OPPOSITION FILED WITHIN TIME LIMIT

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: DK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: ES

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

26N No opposition filed

Effective date: 20200921

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: SI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: IT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: MC

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

REG Reference to a national code

Ref country code: CH

Ref legal event code: PL

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: PL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

GBPC Gb: european patent ceased through non-payment of renewal fee

Effective date: 20200617

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: LU

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20200617

REG Reference to a national code

Ref country code: BE

Ref legal event code: MM

Effective date: 20200630

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: LI

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20200630

Ref country code: IE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20200617

Ref country code: CH

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20200630

Ref country code: GB

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20200617

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: BE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20200630

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: TR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: MT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

Ref country code: CY

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: MK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20191218

P01 Opt-out of the competence of the unified patent court (upc) registered

Effective date: 20230526

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: FR

Payment date: 20230523

Year of fee payment: 14

Ref country code: DE

Payment date: 20230523

Year of fee payment: 14

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: AT

Payment date: 20230525

Year of fee payment: 14