EP3408418A1 - A hot rolled precipitation strengthened and grain refined high strength dual phase steel sheet possessing 600 mpa minimum tensile strength and a process thereof - Google Patents

A hot rolled precipitation strengthened and grain refined high strength dual phase steel sheet possessing 600 mpa minimum tensile strength and a process thereof

Info

Publication number
EP3408418A1
EP3408418A1 EP17740800.2A EP17740800A EP3408418A1 EP 3408418 A1 EP3408418 A1 EP 3408418A1 EP 17740800 A EP17740800 A EP 17740800A EP 3408418 A1 EP3408418 A1 EP 3408418A1
Authority
EP
European Patent Office
Prior art keywords
dual phase
steel sheet
phase steel
daim
max
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
EP17740800.2A
Other languages
German (de)
French (fr)
Other versions
EP3408418B1 (en
Inventor
Appa Rao Chintha
Saurabh KUNDU
Prashant Pathak
Sushil Kumar GIRI
Soumendu MONIA
Subhankar Das BAKSHI
G Senthil KUMAR
Vinay V. MAHASHABDE
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.)
Tata Steel Ltd
Original Assignee
Tata Steel 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 Tata Steel Ltd filed Critical Tata Steel Ltd
Publication of EP3408418A1 publication Critical patent/EP3408418A1/en
Application granted granted Critical
Publication of EP3408418B1 publication Critical patent/EP3408418B1/en
Active legal-status Critical Current
Anticipated expiration legal-status Critical

Links

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
    • C21D8/02—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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • 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
    • C21D1/185—Hardening; Quenching with or without subsequent tempering from an intercritical temperature
    • 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
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • 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
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • 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
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • 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
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips 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
    • 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
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/34—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tyres; for rims
    • 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
    • C22C38/02—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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • 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
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • 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
    • C21D2211/005—Ferrite
    • 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
    • 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
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
    • C21D8/0426—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
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0447—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment
    • C21D8/0463—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment following hot rolling

Definitions

  • the present disclosure relates to a process for producing hot rolled high strength dual phase steel.
  • the disclosure further relates to hot rolled high strength dual phase steel with > 600 MPa tensile strength and 25% total elongation.
  • European patent EP1398392A1 and US patent US8337643 dlsdose a method of producing a hot rolled dual phase (ferrite + martensite) steel of minimum tensile strength of 590 MPa. Though the proposed steels achieved the strength level, it contains high amount of Si (minimum 0.5 wt.% in European patent and 0.2 wt.% in US patent). Presence of Si will lead to surface scales, generally called as tiger marks.
  • European Patent EP2053139B1 discloses a method in which a hot rolled steel sheet is subjected to heat treatment after forming so as to achieve a tensile strength varying in the range of 440 to 640 MPa.
  • the heat treatment after forming which is an essential part of the disclosure, is likely to add to the processing cost and hence is not suitable for mass production.
  • European Patent EP2578714A1 discloses a method of producing hot-rolled steel sheets with a minimum tensile strength of 590 MPa with excellent bake hardenabtlity and stretch-flangeability. According to the proposed method the steel must contain 1.7 to 2.5 wt% of Mn, When added in such large amounts, Mn tends to segregate in the central portion in the thickness direction, which not only induces cracking during press forming but also leads to Inconsistency in achieving the desired stretch-flangeability.
  • the automotive wheel is composed of a disk and a rim. While the disc Is press formed, the rim is flared and then roll formed after flash butt welding. Therefore, the material needed to form the disk needs to have good deep drawability, stretch formability and stretchability, whereas the material needed to form ttie rim needs to have good formability after welding.
  • Hot rolled DP steels with a tensile strength of 600 MPa have become a very popular choice for wheel disc applications owing to their superior strength and formability and at the same time good stretchability (high n value) and spot weldability.
  • HR-DP 600 steel it is difficult to produce the HR-DP 600 steel in any mill because many process parameters, e.g. the finish rolling temperature, cooling rate etc. are needed to be optimized and fine tuned keeping in mind the mill configuration e.g. the length of the run out table, water volume available etc. in order to obtain the desired microstructural features which in turn will decide the final mechanical properties.
  • All the existing patents and literature have considerable amount of Si to increase ferrite strength in order to fatigue life of the steel.
  • Another object of the disclosure is to propose process of producing hot rolled precipitation strengthened high strength dual phase steel sheet, with lower percentage of Si.
  • Another object of the disclosure to propose hot rolled precipitation strengthened high strength dual phase steel sheet the tensile strength more than 600 MPa with lower percentage of Si.
  • Still another object of the disclosure is to propose hot rolled precipitation strengthened high strength dual phase steel sheet, with lower percentage of Si.
  • the disclosure provides a process for producing dual phase steel sheet.
  • the process comprises steps of making a liquid steel having chemical composition in wt% of C: 0.03 - 0.12 , Mn: 0.8 - 1.5, S: ⁇ 0.1, Cr: 0.3 - 0.7 ,S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max; continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840 ⁇ 30 deg.
  • FRT finish rolling temperature
  • FIG. 1 illustrates various steps for a process for making a high strength dual phase steel in accordance with an embodiment of the disclosure.
  • FIG. 2 illustrates a schematic diagram of cooling profile to obtain the high strength dual phase steel in accordance with an embodiment of the disclosure.
  • FIG. 3 illustrates a tensile stress - strain plot of strip 1 accordance with an embodiment of the disclosure.
  • FIG. 4 illustrates an optical micrograph of Strip 1 (Nital etched) in accordance with an embodiment of the disclosure.
  • FIG. 5 illustrates an optical image of Lepera etched sample white phase: Martensite (a); Dark phase: Ferrite (a! In accordance with an embodiment of the disclosure.
  • FIG. 6 illustrates an optical image of Le pera etched sample: Fine grains as small as 2 ⁇ can be noticed in accordance with an embodiment of the disclosure.
  • FIG. 7 illustrates scanning electron microscopy images of strip 1 in accordance with an embodiment of the disclosure.
  • FIG 8(a) illustrates bright field TEM micrograph of one of the precipitates in the ferrite matrix
  • Various embodiments of the disclosure provide a process for producing dual phase steel sheet comprising steps of making a liquid steel having chemical composition in wt% of C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si: ⁇ 0.1, Cr: 0.3 - 0.7 ,S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max; continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840 ⁇ 30 deg.
  • FRT finish rolling temperature
  • a dual phase steel sheet comprising a chemical composition In wt% C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si O.i, Cr: 0.3 - 0.7, S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max.
  • FIG.1 Shown in FIG.1 is a process (100) for producing dual phase steel sheet
  • Step (104) a liquid steel is made. Following is the composition of the liquid steel (in wt.%) C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si: ⁇ 0.1, Cr: 0.3 - 0.7 ,S- 0.008 max, P- 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V-0.06 max.
  • the liquid steel in wt.% C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si: ⁇ 0.1, Cr: 0.3 - 0.7 ,S- 0.008 max, P- 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V-0.06 max.
  • each alloying element and the limitations imposed on each element are essential for achieving the target microstructure and properties.
  • Carbon is one of the most effective and economical strengthening elements. Carbon combines with Nb or V to form carbides or carbonitrides which bring about precipitation strengthening. This requires a minimum of 0.03%C In the steel. However, In order to have good weld-ability, the carbon content has to be restricted to less than 0.12%.
  • Mn 0.8-1.5%: Manganese not only Imparts solid solution strengthening to the ferrite but It also lowers the austenite to ferrite transformation temperature thereby refining the ferrite grain size. However, the Mn level cannot be increased to beyond 1.5% as at such high levels It enhances centerllne segregation during continuous casting.
  • Si ⁇ 0.1 wt % Silicon like Mn is a very efficient solid solution strengthening element
  • Si leads surface scale problems In hot rolling and hence it should be restricted to less than 0.1% in order to prevent the formation of surface scales.
  • Nb 0.035% maximum: Niobium is the most potent microalloying element for grain refinement even when it is added in very small amounts. When in solid solution it lowers the austenite to ferrite transformation temperature which not only refines the ferrite grain size but also promotes the formation of lower transformation products like balnlte. However, to ensure the effectiveness of Nb, it should not be allowed to precipitate before the transformation temperature Is reached. To ensure that the entire Nb content remains in solution before rolling commences and it is alone added, the maximum Nb content is restricted to 0.035%.
  • V 0.06% maximum: Microalloying by Vanadium also leads to precipitation strengthening as well as grain refinement The solubility of Vanadium in austenite is more than that of other microalloying elements and so it is more likely to remain in solution prior to transformation. During phase transformation, vanadium precipitates as carbides and/or nitrides, depending on the relative carbon and nitrogen contents, at grain boundaries resulting !n precipitation strengthening as well as grain refinemenL In order to achieve the desired strengthening, it Is required to add either Nb or V. Both can also be added. If V alone is added, it is required up to 0.06 wt.%.
  • P 0.025% maximum: Phosphorus content should be restricted to 0.025% maximum as higher phosphorus levels can lead to reduction in toughness and weldability due to segregation of P into grain boundaries.
  • N ⁇ 0.007 Too high N content raises the dissolution temperature of Nb(C, N) and hence reduces the effectiveness of Nb. Reducing nitrogen levels also positively affects ageing stability and toughness in the heat-affected zone of the weld seam, as well as resistance to inter- crystalline stress-corrosion cracking. Thus N levels should be preferably kept below 0.007.
  • Al 0.01 to 0.1 Al is used to remove undesirable oxygen from molten steel and hence steel contains some amount of Al, may be upto 0.05 wt.%. Excess (high) Al in steel making is a major problem as it decreases hot deformation of cast slab besides nozzle clogging during casting. Therefore, Al needs to be restricted to 0.1 wt.%.
  • step (108) the liquid steel is continuously casted into a slab.
  • the liquid steel of the specified composition is first continuously casted either In a conventional continuous caster or a thin slab caster.
  • a thin slab caster the temperature of the cast slab is not allowed to drop to a temperature below 950 °C This is because if the thin slab temperature falls below 950 °C, Nb precipitation occurs. Then it becomes difficult to completely dissolve the precipitates in the subsequent reheating process rendering them ineffective for precipitation strengthening.
  • the slab After casting the slab with the specified composition, the slab is reheated to a temperature of 1100 to 1200 °C for a duration of 20 minutes to 2 hours.
  • the reheating temperature should be above 1100 °C, to ensure complete dissolution of any precipitates of Nb or/and V that may have formed in the preceding processing steps.
  • a reheating temperature greater than 1200 °C is also not desirable because it leads to grain coarsening of austenite and/or excessive scale loss.
  • Step (112) the slab is hot rolled into a hot rolled sheet at finish rolling temperature (F T) 840 ⁇ 30°C.
  • the hot rolling constitute of a roughing step above the recrystaliization temperature and a finishing step below the recrystaliization temperature, when rolling is done in a conventional hot strip mill.
  • the deformation schedule is designed in such a manner that the cast structure is destroyed in the Initial stands and finishing is done below the recrystaliization temperature. More specifically the finish rolling in either set up should be done at a temperature, TF KT given by 840 +/- 30 °C.
  • Laminar cooling on the Run-Out-Table (ROT): At step (116) the hot rolled sheet is cooled on a Run Out Table at cooling rate 40 - 70°C/s. The said cooling rate is maintained to achieve intermediate temperature O ) 720 ⁇ Tm ⁇ 650.
  • the cooling rate should be higher than 40 * C/s to prevent formation of pearlite. Any pearlite, or degenerate pearlite if formed leads to deterioration of both, tensile strength as well as stretch flangeability. High cooling rate also results in lowering the ferrite start temperature which leads to refinement of the ferrite grain size. It also prevents the growth of the ferrite. By increasing the cooling rate and controlling rolling schedule, the desired grain size of 2-6 pm can be achieved. The cooling rate may not be more than 70°C s because then the desired amount of ferrite will not form. This fast cooling is continued up to an intermediate temperature.
  • the intermediate temperature (T IN T) should be 650 ⁇ Tim- ⁇ 720° C.
  • the strip is allowed to naturally cool while being transferred on RoT.
  • the duration of air cooling is critical and is 5 to 7 seconds. If the strip is allowed to cool for less than 5 seconds, then sufficient amount of ferrite will not be formed. On the other hand, if the strip is allowed to air cool more than 7 seconds then it will results in insufficient amount of marten site.
  • austenite transforms to ferrite.
  • entire austenite will not transform to ferrite as time is not sufficient for complete transformation.
  • remaining austenite at the end of natural cooling will be enriched with carbon because ferrite cannot accommodate average carbon content in the steel.
  • the strip Is further cooled rapidly after naturally being cooled at step (120). This ensures the transformation of remaining carbon enriched austenite to marbensite.
  • the cooling rate during this period Is 40 - 70 °C/s to achieve coiling temperature below 400° C.
  • the coifing temperature can be as low as 100 deg.C
  • the microstructure obtained comprises maitensite particle phase in the ferrite matrix.
  • the microstructure is uniform or in other words maitensite phase is distributed uniformly throughout the ferrite matrix. Furthermore, bainite or degenerate pearlite/ pearlite and grain boundary cemenb ' te is avoided and high strength dual phase steel sheet achieves good work hardening rate, low yield point and continuous yielding.
  • the contribution of each of the microstructural components is described below: a) Ferrite:
  • the hot rolled steel sheet according to the present disclosure has 75-90 % ferrite (by vol.).
  • the ferrite Is strengthened by solid solution strengthening contributions from Mn. Using suitable processing conditions, the grain size Is restricted to 2 - 5 pm.
  • Maitensite The amount of maitensite in the microstructure is 10-25% (by vol.). The strengthening from maitensite comes from its structure, carbon content and higher dislocation density.
  • Bainite The amount of martensite in the microstructure is less than 5% (by vol.).
  • the high strength dual phase steel sheet has got improved fatigue life due to the presence of fine precipitate in the ferrite matrix coupled with martensite as second phase.
  • the yield stress of the high strength dual phase steel sheet obtained is 350 - 500 MPa.
  • the tensile strength obtained is min. 600 MPa.
  • the min. uniform elongation is 16% and 22% minimum total elongation.
  • strain hardening exponent ("n") of the high strength dual phase steel sheet is 0.15 - 0.16.
  • Yield to Tensile strength (ratio) of the dual phase steel is 0.6 - 0.8 and the hole expansion ratio in punched condition is about 40%.
  • a slab of the composition (given in Table 1) according to the process 100 (Strip 1) was continuously cast in a CSP mill. Slab was hot rolled. The ROT cooling was done in accordance with the present disclosure and the cooling profile is given in FIG. 2.
  • the mechanical properties steels sheet are listed in Table 2, 3 & 4.
  • the microstructures of the steels are shown in FIGS. 4, 5, 6 & 7. It is dear from the mechanical properties and the miaostructures achieved, that the target properties can be achieved when the chemistry and ROT cooling parameters do conform to the requirements of the disclosure.
  • FIGS. 4, 5, 6 & 7 consist of ferrite and martensite.
  • Tensile test samples with 50 mm gauge length were prepared In accordance to ASTM E8 standard. Typical tensile test plot is given in FIG. 3. It is evident from the figure and table that newly developed steel has minimum 600 MPa tensile strength, 16 % uniform elongation and minimum 22 % total elongation, the strip has high strain hardening co-efficient 0.15, yield ratio (Yield strength to Tensile strength) between 0.6 & 0.8. The steel has dispersion of fine precipitates in ferrite matrix.
  • the identities of these precipitates are confirmed using Energy Dispersive Spectroscopy (EDS) and Selective Area Diffraction (SAD) techniques in TEM.
  • EDS Energy Dispersive Spectroscopy
  • SAD Selective Area Diffraction
  • the precipitates are majorly Nb(C,N) as described in Figure 8 a-f.
  • the steel also has very fine average grain size below 3 ⁇ ,

Landscapes

  • 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 Sheet Steel (AREA)
  • Heat Treatment Of Steel (AREA)
  • Heat Treatment Of Strip Materials And Filament Materials (AREA)

Abstract

The invention relates to a process for producing dual phase steel sheet comprises steps of making a liquid steel having chemical composition in wt% of C: 0.03 - 0.12, Mn: 0.8 1.5, Si:<0.1, O: 0.3 0.7,9 0.008 max, P - 0.025 max, Al- 0,01 to 0.1, N- - 0.007 max Nb: 0.005 0,035, and V- 0.06 max, continuous casting the liquid steel in a slab, hot rolling the dab into a hot rolled sheet at finish rolling temperature (FRT) 840 ±30 deg. C, cooling the hot rolled sheet on Run Out Table at cooling rate,10 - 70oC/s achieving intermediate temperature (TINT) 720 ≤ TINT ≤ 650; natural cooling the hot rolled sheet for duration 5 7 - seconds and rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to martensite, at cooling rate of 40 -70 deg. C/s to achieve coiling temperature below 400 - deg. C.

Description

TITLE ; A HOT ROLLED PRECIPITATION STRENGTHENED AND GRAIN REFINED HIGH STRENGTH DUAL PHASE STEEL SHEET POSSESSING 600 MPa MINIMUM TENSILE STRENGTH AND A PROCESS THEREOF
FIELD OF THE DISCLOSURE
The present disclosure relates to a process for producing hot rolled high strength dual phase steel. The disclosure further relates to hot rolled high strength dual phase steel with > 600 MPa tensile strength and 25% total elongation.
BACKGROUND OF THE DISCLOSURE
Motor vehicle fuel consumption and resultant emission is one of the major contributors to air pollution. Light-weight environmental friendly vehicle design is required to address the problems of environmental pollution. Successful light-weight motor vehicles require utilization of advanced high strength high strength steel (AHSS) sheets. However, because of its poor formability, the AHSS sheet cannot be applied easily to a wide variety of motor vehicle components. Hence, the ductility and formabllity required for AHSS sheet becomes increasingly demanding. Therefore addressing the present scenario has necessitated development of a hot rolled steel sheet with high tensile strength coupled with excellent uniform elongation, working hardening rate and total elongation for automotive component such as wheel web applications.
Hence, In order to replace the existing grades of steel used for automotive structural and wheel web applications, it Is necessary to develop hot-rolled steel sheets which not only possess a minimum tensile strength of 600 MPa but also have good formabllity and good surface quality.
European patent EP1398392A1 and US patent US8337643 dlsdose a method of producing a hot rolled dual phase (ferrite + martensite) steel of minimum tensile strength of 590 MPa. Though the proposed steels achieved the strength level, it contains high amount of Si (minimum 0.5 wt.% in European patent and 0.2 wt.% in US patent). Presence of Si will lead to surface scales, generally called as tiger marks.
European Patent EP2053139B1 discloses a method in which a hot rolled steel sheet is subjected to heat treatment after forming so as to achieve a tensile strength varying in the range of 440 to 640 MPa. However, the heat treatment after forming, which is an essential part of the disclosure, is likely to add to the processing cost and hence is not suitable for mass production.
European Patent EP2578714A1 discloses a method of producing hot-rolled steel sheets with a minimum tensile strength of 590 MPa with excellent bake hardenabtlity and stretch-flangeability. According to the proposed method the steel must contain 1.7 to 2.5 wt% of Mn, When added in such large amounts, Mn tends to segregate in the central portion in the thickness direction, which not only induces cracking during press forming but also leads to Inconsistency in achieving the desired stretch-flangeability.
It is also important to understand automotive wheel to develop the steel. The automotive wheel is composed of a disk and a rim. While the disc Is press formed, the rim is flared and then roll formed after flash butt welding. Therefore, the material needed to form the disk needs to have good deep drawability, stretch formability and stretchability, whereas the material needed to form ttie rim needs to have good formability after welding. After the wheel discs and rims are formed by their respective processes, they are assembled by means of spot welding or arc welding. Hence the materials for both rim and disc use need to have good spot weldability. From the point of view of application, the most important functional requirement for auto- wheels is durability, which can be Increased by increasing the fatigue strength of the wheel material.
The various studies conducted in the recent past show that precipitation hardened steels and dual phase (DP) steels are both suitable for wheel disc application. From the fatigue strength considerations, the upper limit of the tensile strength of steels for wheel use is ~ 600 MPa (or 85 ksi) p". Irie, K. Tsunoyama, M. Shinozakl and T. Kato: SAE Paper No. 880695, 1988], This is because when the tensile strength is increased beyond 600 MPa, the consequent increased notch sensitivity results In lowering of the fatigue strength. Hot rolled DP steels with a tensile strength of 600 MPa (or H -DP 600) have become a very popular choice for wheel disc applications owing to their superior strength and formability and at the same time good stretchability (high n value) and spot weldability. However, it is difficult to produce the HR-DP 600 steel in any mill because many process parameters, e.g. the finish rolling temperature, cooling rate etc. are needed to be optimized and fine tuned keeping in mind the mill configuration e.g. the length of the run out table, water volume available etc. in order to obtain the desired microstructural features which in turn will decide the final mechanical properties. All the existing patents and literature have considerable amount of Si to increase ferrite strength in order to fatigue life of the steel.
OBJECTS OF THE DISCLOSURE
In view of the foregoing limitations inherent in the prior-art, it is an object of the disclosure to propose process for producing hot rolled precipitation strengthened high strength dual phase steel sheet the tensile strength more than 600 Pa with lower percentage of Si.
Another object of the disclosure is to propose process of producing hot rolled precipitation strengthened high strength dual phase steel sheet, with lower percentage of Si.
Another object of the disclosure to propose hot rolled precipitation strengthened high strength dual phase steel sheet the tensile strength more than 600 MPa with lower percentage of Si.
Still another object of the disclosure is to propose hot rolled precipitation strengthened high strength dual phase steel sheet, with lower percentage of Si.
SUMMARY OF THE DISCLOSURE
The disclosure provides a process for producing dual phase steel sheet. The process comprises steps of making a liquid steel having chemical composition in wt% of C: 0.03 - 0.12 , Mn: 0.8 - 1.5, S:<0.1, Cr: 0.3 - 0.7 ,S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max; continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840 ± 30 deg. C; cooling the hot rolled sheet on Run Out Table at cooling rate 40 - 70°C/s achieving intermediate temperature (Tim-) 720 < TINT≤ 650; natural cooling the hot rolled sheet for duration 5 - 7 seconds; and rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to marfcensite, at cooling rate of 40 - 70 deg. Cjs to achieve coiling temperature below 400 deg. C.
BRIEF DESCRIPTION OF THE ACCOMPANYING DRAWINGS
FIG. 1 illustrates various steps for a process for making a high strength dual phase steel in accordance with an embodiment of the disclosure. FIG. 2 illustrates a schematic diagram of cooling profile to obtain the high strength dual phase steel in accordance with an embodiment of the disclosure.
FIG. 3 illustrates a tensile stress - strain plot of strip 1 accordance with an embodiment of the disclosure.
FIG. 4 illustrates an optical micrograph of Strip 1 (Nital etched) in accordance with an embodiment of the disclosure.
FIG. 5 illustrates an optical image of Lepera etched sample white phase: Martensite (a); Dark phase: Ferrite (a!) In accordance with an embodiment of the disclosure.
FIG. 6 illustrates an optical image of Le pera etched sample: Fine grains as small as 2 μτη can be noticed in accordance with an embodiment of the disclosure.
FIG. 7 illustrates scanning electron microscopy images of strip 1 in accordance with an embodiment of the disclosure.
FIG 8(a) illustrates bright field TEM micrograph of one of the precipitates in the ferrite matrix; 8 (b) Dark field image of 8 (a); 8 (c) selected area diffraction pattern from Nb(C,N) precipitate, 8 (d) Dark field image showing Nb(C,N) precipitates; 8 (e) EDS spectrum of the precipitate and 8 (f) Composition of the precipitate
DETAILED DESCRIPTION OF A PREFERRED EMBODIMENT OF THE DISCLOSURE
Various embodiments of the disclosure provide a process for producing dual phase steel sheet comprising steps of making a liquid steel having chemical composition in wt% of C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si:<0.1, Cr: 0.3 - 0.7 ,S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max; continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840 ± 30 deg. C; cooling the hot rolled sheet on Run Out Table at cooling rate 40 - 70°C/s achieving intermediate temperature (TINT) 720 < Tim-≤ 650; natural cooling the hot rolled sheet for duration 5 - 7 seconds; and rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to martensite, at cooling rate of 40 - 70 deg. C/s to achieve coiling temperature below 400 deg. C. Another embodiment of the disclosure provide A dual phase steel sheet, comprising a chemical composition In wt% C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si O.i, Cr: 0.3 - 0.7, S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max.
Shown in FIG.1 is a process (100) for producing dual phase steel sheet At Step (104) a liquid steel is made. Following is the composition of the liquid steel (in wt.%) C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si:<0.1, Cr: 0.3 - 0.7 ,S- 0.008 max, P- 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V-0.06 max.
The addition of each alloying element and the limitations imposed on each element are essential for achieving the target microstructure and properties.
C: 0.03-0.12%: Carbon is one of the most effective and economical strengthening elements. Carbon combines with Nb or V to form carbides or carbonitrides which bring about precipitation strengthening. This requires a minimum of 0.03%C In the steel. However, In order to have good weld-ability, the carbon content has to be restricted to less than 0.12%.
Mn: 0.8-1.5%: Manganese not only Imparts solid solution strengthening to the ferrite but It also lowers the austenite to ferrite transformation temperature thereby refining the ferrite grain size. However, the Mn level cannot be increased to beyond 1.5% as at such high levels It enhances centerllne segregation during continuous casting.
Si < 0.1 wt % : Silicon like Mn is a very efficient solid solution strengthening element However, Si leads surface scale problems In hot rolling and hence it should be restricted to less than 0.1% in order to prevent the formation of surface scales.
Nb: 0.035% maximum: Niobium is the most potent microalloying element for grain refinement even when it is added in very small amounts. When in solid solution it lowers the austenite to ferrite transformation temperature which not only refines the ferrite grain size but also promotes the formation of lower transformation products like balnlte. However, to ensure the effectiveness of Nb, it should not be allowed to precipitate before the transformation temperature Is reached. To ensure that the entire Nb content remains in solution before rolling commences and it is alone added, the maximum Nb content is restricted to 0.035%.
V: 0.06% maximum: Microalloying by Vanadium also leads to precipitation strengthening as well as grain refinement The solubility of Vanadium in austenite is more than that of other microalloying elements and so it is more likely to remain in solution prior to transformation. During phase transformation, vanadium precipitates as carbides and/or nitrides, depending on the relative carbon and nitrogen contents, at grain boundaries resulting !n precipitation strengthening as well as grain refinemenL In order to achieve the desired strengthening, it Is required to add either Nb or V. Both can also be added. If V alone is added, it is required up to 0.06 wt.%.
P: 0.025% maximum: Phosphorus content should be restricted to 0.025% maximum as higher phosphorus levels can lead to reduction in toughness and weldability due to segregation of P into grain boundaries.
S: 0.008% maximum: The Sulphur content has to be limited otherwise it results in a very high inclusion level that deteriorates formability.
N < 0.007: Too high N content raises the dissolution temperature of Nb(C, N) and hence reduces the effectiveness of Nb. Reducing nitrogen levels also positively affects ageing stability and toughness in the heat-affected zone of the weld seam, as well as resistance to inter- crystalline stress-corrosion cracking. Thus N levels should be preferably kept below 0.007.
Al 0.01 to 0.1: Al is used to remove undesirable oxygen from molten steel and hence steel contains some amount of Al, may be upto 0.05 wt.%. Excess (high) Al in steel making is a major problem as it decreases hot deformation of cast slab besides nozzle clogging during casting. Therefore, Al needs to be restricted to 0.1 wt.%.
At step (108) the liquid steel is continuously casted into a slab.
The liquid steel of the specified composition is first continuously casted either In a conventional continuous caster or a thin slab caster. When cast In a thin slab caster, the temperature of the cast slab is not allowed to drop to a temperature below 950 °C This is because if the thin slab temperature falls below 950 °C, Nb precipitation occurs. Then it becomes difficult to completely dissolve the precipitates in the subsequent reheating process rendering them ineffective for precipitation strengthening.
Reheating: After casting the slab with the specified composition, the slab is reheated to a temperature of 1100 to 1200 °C for a duration of 20 minutes to 2 hours. The reheating temperature should be above 1100 °C, to ensure complete dissolution of any precipitates of Nb or/and V that may have formed in the preceding processing steps. A reheating temperature greater than 1200 °C is also not desirable because it leads to grain coarsening of austenite and/or excessive scale loss.
At Step (112) the slab is hot rolled into a hot rolled sheet at finish rolling temperature (F T) 840 ± 30°C.
The hot rolling constitute of a roughing step above the recrystaliization temperature and a finishing step below the recrystaliization temperature, when rolling is done in a conventional hot strip mill. In case Continuous Strip Processing is used for producing this steel, where there is no separate roughing mill, the deformation schedule is designed in such a manner that the cast structure is destroyed in the Initial stands and finishing is done below the recrystaliization temperature. More specifically the finish rolling in either set up should be done at a temperature, TFKT given by 840 +/- 30 °C.
Laminar cooling on the Run-Out-Table (ROT): At step (116) the hot rolled sheet is cooled on a Run Out Table at cooling rate 40 - 70°C/s. The said cooling rate is maintained to achieve intermediate temperature O ) 720 < Tm≤ 650.
The cooling rate should be higher than 40* C/s to prevent formation of pearlite. Any pearlite, or degenerate pearlite if formed leads to deterioration of both, tensile strength as well as stretch flangeability. High cooling rate also results in lowering the ferrite start temperature which leads to refinement of the ferrite grain size. It also prevents the growth of the ferrite. By increasing the cooling rate and controlling rolling schedule, the desired grain size of 2-6 pm can be achieved. The cooling rate may not be more than 70°C s because then the desired amount of ferrite will not form. This fast cooling is continued up to an intermediate temperature. The intermediate temperature (TINT) should be 650 < Tim- < 720° C.
At step (120), the strip is allowed to naturally cool while being transferred on RoT. The duration of air cooling is critical and is 5 to 7 seconds. If the strip is allowed to cool for less than 5 seconds, then sufficient amount of ferrite will not be formed. On the other hand, if the strip is allowed to air cool more than 7 seconds then it will results in insufficient amount of marten site.
During this period, austenite transforms to ferrite. However, entire austenite will not transform to ferrite as time is not sufficient for complete transformation. As a result remaining austenite at the end of natural cooling will be enriched with carbon because ferrite cannot accommodate average carbon content in the steel.
At step (124) the strip Is further cooled rapidly after naturally being cooled at step (120). This ensures the transformation of remaining carbon enriched austenite to marbensite. The cooling rate during this period Is 40 - 70 °C/s to achieve coiling temperature below 400° C. The coifing temperature can be as low as 100 deg.C
The strengthening contributions from solid solution elements and microalloying elements are restricted. Also, the extent of possible grain refinement, by controlled rolling and cooling is limited to 2 pm due to which the high strength dual phase steel obtained.
The microstructure obtained comprises maitensite particle phase in the ferrite matrix. The microstructure is uniform or in other words maitensite phase is distributed uniformly throughout the ferrite matrix. Furthermore, bainite or degenerate pearlite/ pearlite and grain boundary cemenb'te is avoided and high strength dual phase steel sheet achieves good work hardening rate, low yield point and continuous yielding. The contribution of each of the microstructural components is described below: a) Ferrite: The hot rolled steel sheet according to the present disclosure has 75-90 % ferrite (by vol.). The ferrite Is strengthened by solid solution strengthening contributions from Mn. Using suitable processing conditions, the grain size Is restricted to 2 - 5 pm. This grain refinement of ferrite leads to strengthening of the ferrite, the amount of which is decided by the Hall-Petch relationship. Also it is precipitation strengthened by the formation of fine Nb,V(CN) precipitates. b) Maitensite: The amount of maitensite in the microstructure is 10-25% (by vol.). The strengthening from maitensite comes from its structure, carbon content and higher dislocation density.
c) Bainite: The amount of martensite in the microstructure is less than 5% (by vol.).
The high strength dual phase steel sheet has got improved fatigue life due to the presence of fine precipitate in the ferrite matrix coupled with martensite as second phase. The yield stress of the high strength dual phase steel sheet obtained is 350 - 500 MPa. The tensile strength obtained is min. 600 MPa. The min. uniform elongation is 16% and 22% minimum total elongation.
Further, strain hardening exponent ("n") of the high strength dual phase steel sheet is 0.15 - 0.16. Yield to Tensile strength (ratio) of the dual phase steel is 0.6 - 0.8 and the hole expansion ratio in punched condition is about 40%.
Experimental Analysis
For the purpose of example only, a slab of the composition (given in Table 1) according to the process 100 (Strip 1) was continuously cast in a CSP mill. Slab was hot rolled. The ROT cooling was done in accordance with the present disclosure and the cooling profile is given in FIG. 2. The mechanical properties steels sheet are listed in Table 2, 3 & 4. The microstructures of the steels are shown in FIGS. 4, 5, 6 & 7. It is dear from the mechanical properties and the miaostructures achieved, that the target properties can be achieved when the chemistry and ROT cooling parameters do conform to the requirements of the disclosure.
The optical (both Nital and Le pera etched) and SEM microstructures are presented in FIGS. 4, 5, 6 & 7 which consist of ferrite and martensite. Tensile test samples with 50 mm gauge length were prepared In accordance to ASTM E8 standard. Typical tensile test plot is given in FIG. 3. It is evident from the figure and table that newly developed steel has minimum 600 MPa tensile strength, 16 % uniform elongation and minimum 22 % total elongation, the strip has high strain hardening co-efficient 0.15, yield ratio (Yield strength to Tensile strength) between 0.6 & 0.8. The steel has dispersion of fine precipitates in ferrite matrix. The identities of these precipitates are confirmed using Energy Dispersive Spectroscopy (EDS) and Selective Area Diffraction (SAD) techniques in TEM. The precipitates are majorly Nb(C,N) as described in Figure 8 a-f. The steel also has very fine average grain size below 3 μηη,
Table 1: Composition (in wt.%) of the samples tested
Table 2: Tensile properties of the tested samples
Table 3: Hole expansion of the tested samples
Table 4: Quantification of microstructural constituents

Claims

WE CLAIM :
1. A process for producing dual phase steel sheet, comprising steps of :-
- making a liquid steel having chemical composition in wt% of
C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si:<0.1, Cr: 0.3 - 0.7, S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max;
- continuous casting the liquid steel into a slab;
- hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840 ± 30°C;
- cooling the hot rolled sheet on Run Out Table at cooling rate 40 - 70°cys achieving intermediate temperature fJ ) 720 < Tmr≤ 650;
- natural cooling the hot rolled sheet for duration 5 - 7 seconds; and
- rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to martensite, at cooling rate of 40 - 70 ° s to achieve coiling temperature below 400° C.
2. The process as claimed in claim 1, wherein the dual phase steel sheet being reheated to a temperature range 1100-1200°C for dissolution of precipitates.
3. The process as claimed in claim 1, wherein yield stress of the dual phase steel sheet is 350 - 500 MPa.
4. The process as claimed in daim 1, wherein the dual phase steel sheet have 600 MPa minimum tensile strength.
5. The process as daimed in daim 1, wherein the dual phase steel sheet have 16% minimum uniform elongation.
6. The process as claimed in daim 1, wherein the dual phase steel sheet have 22% minimum total elongation.
7. The process as daimed in daim 1, wherein strain hardening exponent fn") of the dual phase steel sheet is 0.15 - 0.16
8. The process as daimed in daim l, wherein yield to tensile strength (ratio) of the dual phase steel is 0.6 - 0.8.
9. The process as daimed in daim 1, wherein the dual phase steel sheet have hole expansion ratio in punched condition is about 40%.
10. The process as daimed in daim 1, wherein the dual phase steel sheet have 75 - 90% ferrite , 10 - 25% martenslte and < 5% bainite by volume.
11. The process as claimed in daim 1, wherein the grain size of high strength dual phase steel is 2-5 Mm.
12. A dual phase steel sheet, comprising: a chemical composition in wt% C: 0.03 - 0.12 , Mn: 0.8 - 1.5, Si:<0.1, Cr: 0.3 - 0.7, S- 0.008 max, P - 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005 - 0.035, and V- 0.06 max.
13. The dual phase steel sheet as daimed in daim 12, wherein yield stress of the dual phase steel sheet is 350 - 500 MPa.
14. The dual phase steel sheet as claimed In daim 12, wherein the dual phase steel sheet has 600 MPa min. tensile strength.
15. The dual phase steel sheet as daimed in daim 12, wherein the dual phase steel sheet has 16% min uniform elongation.
16. The dual phase steel sheet as daimed in daim 12, wherein the dual phase steel sheet has 22% min total elongation.
17. The dual phase steel sheet as daimed in daim 12, wherein strain hardening exponent ("n") of the dual phase steel sheet is 0.15 - 0.16.
18. The dual phase steel sheet as claimed in daim 12, wherein yield to tensile strength (ratio) is 0.6 - 0.8.
19. The dual phase steel sheet as claimed in daim 12, wherein hole expansion ratio in punched condition is about 40%.
20. The dual phase steel sheet as daimed in daim 12, wherein the dual phase steel sheet has 75 - 90% ferrite, 10 - 25% martensite and < 5% bainite by volume.
21. The dual phase steel sheet as daimed in daim 12, wherein the grain size of dual phase steel sheet is 2-5 μιτι.
EP17740800.2A 2017-02-10 2017-05-10 A hot rolled precipitation strengthened and grain refined high strength dual phase steel sheet possessing 600 mpa minimum tensile strength and a process thereof Active EP3408418B1 (en)

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
IN201731004831 2017-02-10
PCT/IN2017/050171 WO2018146695A1 (en) 2017-02-10 2017-05-10 A hot rolled precipitation strengthened and grain refined high strength dual phase steel sheet possessing 600 mpa minimum tensile strength and a process thereof

Publications (2)

Publication Number Publication Date
EP3408418A1 true EP3408418A1 (en) 2018-12-05
EP3408418B1 EP3408418B1 (en) 2023-05-10

Family

ID=59366465

Family Applications (1)

Application Number Title Priority Date Filing Date
EP17740800.2A Active EP3408418B1 (en) 2017-02-10 2017-05-10 A hot rolled precipitation strengthened and grain refined high strength dual phase steel sheet possessing 600 mpa minimum tensile strength and a process thereof

Country Status (6)

Country Link
US (1) US20200123630A1 (en)
EP (1) EP3408418B1 (en)
JP (1) JP7063810B2 (en)
KR (1) KR20190131408A (en)
ES (1) ES2951778T3 (en)
WO (1) WO2018146695A1 (en)

Families Citing this family (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN111020096B (en) * 2019-11-22 2021-05-28 辽宁科技大学 Single LF (low frequency) process low-nitrogen control method for dual-phase automobile steel DP590
CN113926892B (en) * 2020-06-29 2024-07-12 宝山钢铁股份有限公司 Stamping forming process and application of hot-rolled ultra-high strength dual-phase steel part with tensile strength of more than or equal to 980MPa
CN112430772A (en) * 2020-09-28 2021-03-02 甘肃酒钢集团宏兴钢铁股份有限公司 CSP flow-based medium-temperature coiling type hot rolling DP600 production method
EP4317515A1 (en) * 2021-04-02 2024-02-07 Baoshan Iron & Steel Co., Ltd. Dual-phase steel and hot-dip galvanized dual-phase steel having tensile strength greater than or equal to 980mpa and method for manufacturing same by means of rapid heat treatment
CN115029622B (en) * 2022-04-29 2023-05-23 武汉钢铁有限公司 High-surface-quality hot-rolled dual-phase steel and production process thereof
CN115491601A (en) * 2022-09-20 2022-12-20 武汉钢铁有限公司 Economical magnet yoke steel with yield strength of 350MPa grade produced by CSP production line and production method

Family Cites Families (16)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS57137426A (en) * 1981-02-20 1982-08-25 Kawasaki Steel Corp Production of low yield ratio, high tensile hot rolled steel plate by mixed structure
JP3235416B2 (en) * 1995-07-24 2001-12-04 住友金属工業株式会社 Manufacturing method of high strength hot rolled steel sheet with excellent workability and fatigue properties
JP4051999B2 (en) 2001-06-19 2008-02-27 Ｊｆｅスチール株式会社 High tensile hot-rolled steel sheet excellent in shape freezing property and durability fatigue property after forming, and method for producing the same
JP4502646B2 (en) 2004-01-21 2010-07-14 株式会社神戸製鋼所 High-strength hot-rolled steel sheet with excellent workability, fatigue characteristics and surface properties
US8337643B2 (en) 2004-11-24 2012-12-25 Nucor Corporation Hot rolled dual phase steel sheet
JP4661306B2 (en) * 2005-03-29 2011-03-30 Ｊｆｅスチール株式会社 Manufacturing method of ultra-high strength hot-rolled steel sheet
JP5040197B2 (en) 2006-07-10 2012-10-03 Ｊｆｅスチール株式会社 Hot-rolled thin steel sheet with excellent workability and excellent strength and toughness after heat treatment and method for producing the same
KR101142620B1 (en) * 2007-03-27 2012-05-03 신닛뽄세이테쯔 카부시키카이샤 High-strength hot rolled steel sheet being free from peeling and excelling in surface and burring properties and process for manufacturing the same
JP5348071B2 (en) 2010-05-31 2013-11-20 Ｊｆｅスチール株式会社 High strength hot rolled steel sheet and method for producing the same
KR20120097173A (en) * 2011-02-24 2012-09-03 현대제철 주식회사 High strength steel sheet and method of manufacturing the same
JP2012197516A (en) 2012-05-08 2012-10-18 Sumitomo Metal Ind Ltd Method for manufacturing hot-rolled steel sheet
BR112014031739B1 (en) * 2012-06-26 2019-05-28 Nippon Steel & Sumitomo Metal Corporation HIGH-RESISTANCE HOT LAMINATED STEEL SHEET AND SAME PRODUCTION METHOD
KR20150025952A (en) * 2013-08-30 2015-03-11 현대제철 주식회사 High strength plated hot-rolled steel sheet and method of manufacturing the same
JP6275510B2 (en) * 2014-02-27 2018-02-07 Ｊｆｅスチール株式会社 High strength hot rolled steel sheet and method for producing the same
CN104451402A (en) 2014-12-19 2015-03-25 山东钢铁股份有限公司 700MPa-grade hot-rolled dual-phase steel and manufacturing method thereof
DE102016121905A1 (en) * 2016-11-15 2018-05-17 Salzgitter Flachstahl Gmbh Method for producing dual-phase steel wheel discs with improved cold workability

Also Published As

Publication number Publication date
KR20190131408A (en) 2019-11-26
WO2018146695A1 (en) 2018-08-16
ES2951778T3 (en) 2023-10-24
US20200123630A1 (en) 2020-04-23
JP2020509151A (en) 2020-03-26
JP7063810B2 (en) 2022-05-09
EP3408418B1 (en) 2023-05-10

Similar Documents

Publication Publication Date Title
EP3408418B1 (en) A hot rolled precipitation strengthened and grain refined high strength dual phase steel sheet possessing 600 mpa minimum tensile strength and a process thereof
JP3889766B2 (en) High-strength hot-rolled steel sheet excellent in hole expansion workability and its manufacturing method
KR100942087B1 (en) High strength hot rolled steel sheet excellent in bore expanding workability and method for production thereof
JP5041084B2 (en) High-tensile hot-rolled steel sheet excellent in workability and manufacturing method thereof
KR100958019B1 (en) Dual phase steel sheet and method for manufacturing the same
JP2016534230A (en) High hardness hot rolled steel product and method for producing the same
WO2008007753A1 (en) Hot-rolled steel sheets excellent both in workability and in strength and toughness after heat treatment and process for production thereof
JP2009503267A (en) Method for producing high-strength steel sheet having excellent ductility and steel sheet produced thereby
KR20030076430A (en) High tensile hot rolled steel sheet excellent in elongation property and elongation flanging property, and method for producing the same
JP3889765B2 (en) High-strength hot-rolled steel sheet excellent in hole expansion workability and its manufacturing method
RU2393238C1 (en) Procedure for production of plate iron low-alloyed strip
CN112292472B (en) High-strength steel sheet having excellent collision resistance and method for producing same
CN112739834A (en) Hot-rolled steel sheet and method for producing same
KR20240000646A (en) Hot rolled steel sheet with high hole expansion ratio and manufacturing process thereof
KR100903546B1 (en) High tensile hot rolled steel sheet excellent in shape freezing property and endurance fatigue characteristics after forming
JP4867177B2 (en) High tensile hot rolled steel sheet excellent in bake hardenability and formability and method for producing the same
JP4848722B2 (en) Method for producing ultra-high-strength cold-rolled steel sheet with excellent workability
KR102451005B1 (en) High-strength steel sheet having excellent thermal stability and method for mnufacturing thereof
KR101630977B1 (en) High strength hot rolled steel sheet having excellent formability and method for manufacturing the same
KR101543836B1 (en) High strength hot rolled steel sheet having excellent impact resistance and formability and method for manufacturing the same
KR101505252B1 (en) Cold-rolled steel sheet for outcase of car having low yield ratio with excellent formability and method of manufacturing the same
KR101657835B1 (en) High strength hot-rolled steel sheet having excellent press formability and method for manufacturing the same
RU2803955C1 (en) Cold-rolled and annealed steel sheet and production method
KR101467026B1 (en) Steel sheet and method of manufacturing the same
KR20100047021A (en) Hot-rolled steel sheet having ultra-high strength, and method for producing the same

Legal Events

Date Code Title Description
STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: UNKNOWN

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

Free format text: STATUS: THE INTERNATIONAL PUBLICATION HAS BEEN MADE

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

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

Free format text: STATUS: REQUEST FOR EXAMINATION WAS MADE

17P Request for examination filed

Effective date: 20180829

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 RS SE SI SK SM TR

AX Request for extension of the european patent

Extension state: BA ME

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

Free format text: STATUS: REQUEST FOR EXAMINATION WAS MADE

DAV Request for validation of the european patent (deleted)
DAX Request for extension of the european patent (deleted)
STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: EXAMINATION IS IN PROGRESS

17Q First examination report despatched

Effective date: 20210525

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

INTG Intention to grant announced

Effective date: 20221205

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 RS 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: AT

Ref legal event code: REF

Ref document number: 1566724

Country of ref document: AT

Kind code of ref document: T

Effective date: 20230515

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: 602017068586

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: NL

Ref legal event code: FP

REG Reference to a national code

Ref country code: LT

Ref legal event code: MG9D

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

Ref country code: NL

Payment date: 20230712

Year of fee payment: 7

REG Reference to a national code

Ref country code: AT

Ref legal event code: MK05

Ref document number: 1566724

Country of ref document: AT

Kind code of ref document: T

Effective date: 20230510

REG Reference to a national code

Ref country code: ES

Ref legal event code: FG2A

Ref document number: 2951778

Country of ref document: ES

Kind code of ref document: T3

Effective date: 20231024

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

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: 20230510

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: 20230911

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: 20230810

Ref country code: AT

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: 20230510

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

Ref country code: GB

Payment date: 20230712

Year of fee payment: 7

Ref country code: ES

Payment date: 20230717

Year of fee payment: 7

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

Ref country code: RS

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: 20230510

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: 20230510

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: 20230510

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: 20230510

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: 20230910

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: 20230510

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: 20230811

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

Ref country code: FR

Payment date: 20230712

Year of fee payment: 7

Ref country code: DE

Payment date: 20230712

Year of fee payment: 7

Ref country code: BE

Payment date: 20230712

Year of fee payment: 7

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

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: 20230510

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: 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: 20230510

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: 20230510

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: 20230510

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: 20230510

Ref country code: LU

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

Effective date: 20230510

Ref country code: LI

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

Effective date: 20230531

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: 20230510

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: 20230510

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: 20230510

Ref country code: CH

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

Effective date: 20230531

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 602017068586

Country of ref document: DE

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

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: 20230510

REG Reference to a national code

Ref country code: IE

Ref legal event code: MM4A

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

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: 20230510

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: IE

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

Effective date: 20230510

26N No opposition filed

Effective date: 20240213

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

Ref country code: IE

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

Effective date: 20230510

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: 20230510

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: 20230510