WO2016157896A1 - Tôle d'acier laminée à chaud et son procédé de production - Google Patents

Tôle d'acier laminée à chaud et son procédé de production Download PDF

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WO2016157896A1
WO2016157896A1 PCT/JP2016/001834 JP2016001834W WO2016157896A1 WO 2016157896 A1 WO2016157896 A1 WO 2016157896A1 JP 2016001834 W JP2016001834 W JP 2016001834W WO 2016157896 A1 WO2016157896 A1 WO 2016157896A1
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steel sheet
hot
lath
phase
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PCT/JP2016/001834
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English (en)
Japanese (ja)
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田中　孝明
太郎 木津
俊介 豊田
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Ｊｆｅスチール株式会社
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Priority to CN201680020526.3A priority Critical patent/CN107429362B/zh
Priority to MX2017012493A priority patent/MX2017012493A/es
Priority to JP2016549181A priority patent/JP6075517B1/ja
Priority to US15/561,436 priority patent/US20180119240A1/en
Priority to EP16771783.4A priority patent/EP3279353B1/fr
Priority to KR1020177029834A priority patent/KR101989262B1/ko
Publication of WO2016157896A1 publication Critical patent/WO2016157896A1/fr

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    • C—CHEMISTRY; METALLURGY
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    • 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
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    • 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")
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • 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
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    • 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
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    • 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
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    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
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    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
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    • 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
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • 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
    • C21D2211/004—Dispersions; Precipitations
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
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    • 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
    • C21D8/0273—Final recrystallisation annealing

Definitions

  • the present invention has a high tensile strength (TS) of 780 MPa or more, excellent stretch flangeability and punching, suitable for structural steel materials such as automobile parts and other transportation machinery and construction steel materials. Further, the present invention relates to a hot-rolled steel sheet having excellent properties and also excellent in production stability and a method for producing the same.
  • TS tensile strength
  • Patent Document 1 discloses that a steel sheet structure is a ferrite single-phase structure having a low dislocation density and excellent workability, and further, fine carbides are dispersed and precipitated in ferrite to strengthen precipitation, thereby hot rolling.
  • a steel sheet having improved strength while maintaining stretch flangeability of the steel sheet is disclosed.
  • the normal burring process is performed using a steel sheet punched into a predetermined shape.
  • the punching clearance usually changes due to die temperature rise or die wear due to continuous pressing. May occur.
  • Patent Document 2 discloses that Fe-type carbides having a bainite phase exceeding 92% by volume and an average interval of bainite laths of 0.60 ⁇ m or less and precipitated in grains among all Fe-type carbides.
  • a high-strength hot-rolled steel sheet with improved mass production punchability by making the number ratio of 10% or more is disclosed.
  • JP 2012-26034 A Japanese Unexamined Patent Publication No. 2014-205888
  • the steel sheet described in Patent Document 1 achieves both high strength and excellent stretch flangeability.
  • the steel sheet structure is substantially a ferrite single phase, there is almost no inclusion that becomes a starting point of voids when the steel sheet is punched.
  • the punched end surface may be roughened.
  • the steel sheet described in Patent Document 2 has a steel sheet structure mainly composed of a predetermined bainite by controlling hot rolling conditions, thereby obtaining excellent punchability.
  • a bainite structure is characterized in that mechanical properties such as tensile strength tend to fluctuate with respect to fluctuations in the coiling temperature.
  • the present invention has been developed in view of the above situation, has a tensile strength (TS): high strength of 780 MPa or more, has excellent stretch flangeability and punchability, and further has manufacturing stability.
  • An object of the present invention is to provide a hot-rolled steel sheet that is excellent in combination with its advantageous production method.
  • the present inventors diligently studied a method for increasing the strength of a steel sheet while maintaining workability, particularly stretch flangeability, and having excellent punchability and suppressing variations in mechanical properties with respect to variations in manufacturing conditions. As described above, in order to improve the stretch flangeability of the steel sheet, it is effective to make the strength in the metal structure uniform. As such a method, a method of increasing strength by solid solution strengthening or precipitation strengthening as a ferrite single phase structure, or a method of increasing strength by strengthening structure as a bainite single phase structure can be considered.
  • a steel sheet having a bainite single-phase structure has excellent stretch flangeability.
  • a steel sheet having a bainite single-phase structure has a large number of Fe-based carbides in the bainite structure, which serves as a starting point for voids at the time of punching, and thus has excellent punchability.
  • the mechanical properties such as strength greatly vary depending on the transformation temperature in the bainite structure, there is a concern that the variation in mechanical properties with respect to variations in hot rolling conditions such as the coiling temperature will increase.
  • the present inventors considered to reduce the influence of fluctuations in hot rolling conditions by adding a tempering treatment to a structure mainly composed of bainite and martensite.
  • a tempering treatment to a bainite or martensite structure
  • variations in mechanical properties due to changes in hot rolling conditions are greatly reduced, but at the same time the steel sheet strength is greatly reduced.
  • the steel sheet since the form of the Fe-based carbide in the tempered bainite or tempered martensite phase varies depending on the annealing conditions, the steel sheet does not necessarily have excellent punchability depending on the annealing conditions.
  • the present inventors suppressed the reduction of the steel sheet strength as described above when adding a tempering treatment to the structure mainly composed of bainite and further martensite, and excellent stretch flangeability and punchability.
  • MC type carbides such as TiC are dispersed and precipitated inside the lath and within the lath boundary, which suppresses coarsening of the lath during annealing and further disappearance of the lath due to recovery, and maintains high steel sheet strength even after annealing. I found out that I can do it.
  • MC type carbide is TiC, NbC, VC, (Ti where the atomic ratio of M element (M element includes Ti, Nb, V, Mo, etc.) and C is approximately 1: 1. , Mo) C and other carbides.
  • M element does not have to be a single type, and may be a composite carbide containing a plurality of metal elements.
  • N-containing carbonitrides and composite carbonitrides may be used.
  • the present inventors have further studied earnestly, by appropriately controlling the thermal history when cooling from the highest heating temperature to room temperature in the annealing process, the remaining structure other than the tempered martensite phase and the tempered bainite phase, In particular, the inventors have found that the formation of martensite phase, coarse pearlite phase and retained austenite phase is suppressed, and that, in addition to high strength and excellent punchability, it can also have excellent stretch flangeability.
  • the present invention was completed after further studies based on the above findings.
  • the gist configuration of the present invention is as follows. 1. % By mass C: 0.03% or more and 0.20% or less, Si: 0.4% or less, Mn: 0.5% to 2.0%, P: 0.03% or less, S: 0.03% or less, Al: 0.1% or less, N: 0.01% or less and Ti: 0.03% or more and 0.15% or less, with the balance consisting of Fe and inevitable impurities,
  • the total area ratio of the tempered bainite phase and the tempered martensite phase is 70% or more, and the total area ratio of the coarse pearlite phase, the martensite phase and the retained austenite phase is 10% or less,
  • the tempered bainite phase and the tempered martensite phase have lath having an average width of 1.0 ⁇ m or less as a substructure, and among Fe-based carbides precipitated in the lath and at the lath boundary, the proportion of those having an aspect ratio of 5 or less 80% or more, and MC type carbide having an average particle size of 20 nm
  • V 0.01% to 0.3%
  • Mo 0.01% to 0.3%
  • the hot-rolled steel sheet according to 1 or 2 comprising:
  • composition further, by mass%, one or more of REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn and Cs
  • REM halogen-containing compound
  • the steel material having the composition according to any one of 1 to 4 above is heated to an austenite single phase region, subjected to hot rolling consisting of rough rolling and finish rolling, and after the finish rolling is finished, A hot rolling process for cooling and winding, After the hot rolling step, pickling the steel sheet, and then having a continuous annealing step for continuous annealing,
  • the finish rolling temperature is 850 ° C. or more and 1000 ° C. or less
  • the average cooling rate to 500 ° C. is 30 ° C./s or more
  • the winding temperature is 500 ° C. or less
  • the maximum heating temperature of the steel sheet is 700 ° C.
  • the time when the steel plate temperature when heating the steel plate to the maximum heating temperature is 600 ° C. or more and 700 ° C. or less is 20 seconds or more and 1000 seconds or less, The time when the steel sheet temperature is over 700 ° C. is set to 200 seconds or less, The average cooling rate up to 530 ° C when cooling the steel sheet from the maximum heating temperature is 8 ° C / s or more and 25 ° C / s or less, and after the cooling is stopped, the holding time in the temperature range of 470 ° C or more and 530 ° C or less is set.
  • tensile strength (TS) high strength of 780 MPa or more, excellent stretch flangeability, suitable for structural steel materials such as automobile machinery and other transportation machinery parts and construction steel materials.
  • FIG. 2 is a schematic diagram showing an example of a structure in which a tempered bainite phase and a tempered martensite phase have lath as a substructure, and Fe-based carbides are precipitated and lath MC-type carbides are dispersed and precipitated in the lath inside and lath boundaries.
  • C 0.03% to 0.20% C improves the strength of the steel and promotes the formation of bainite and martensite during hot rolling. Therefore, in the present invention, the C content needs to be 0.03% or more. On the other hand, when the C content exceeds 0.20%, the carbon equivalent becomes excessively large and the weldability of the steel sheet is deteriorated. Therefore, the C content is set to 0.03% or more and 0.20% or less. Preferably, it is 0.04% or more and 0.18% or less, more preferably more than 0.05% and 0.15% or less.
  • Si 0.4% or less Si is usually positively contained in a high-strength steel sheet as an effective element for improving the steel sheet strength without reducing ductility (elongation). However, if the Si content exceeds 0.4%, an oxide is formed on the surface of the steel sheet during heat treatment, which causes deterioration of plating adhesion. Therefore, the Si content is 0.4% or less. Preferably it is 0.3% or less, More preferably, it is 0.2% or less. Note that Si may be added to an impurity level, or may be O%.
  • Mn 0.5% or more and 2.0% or less
  • Mn is an element that contributes to increasing the strength of steel by solid solution.
  • Mn is an element that promotes the formation of bainite and martensite during hot rolling by improving hardenability. In order to obtain such an effect, the Mn content needs to be 0.5% or more.
  • the Mn content is 0.5% or more and 2.0% or less.
  • they are 0.8% or more and 1.8% or less, More preferably, they are 1.0% or more and 1.7% or less.
  • P 0.03% or less
  • P is a harmful element that segregates at grain boundaries to reduce elongation, induces cracking during processing, and further degrades impact resistance. Therefore, the P content is 0.03% or less. However, excessive P removal leads to an increase in refining time and cost, so the P content is preferably 0.002% or more.
  • S 0.03% or less S is present in steel as MnS or TiS, and promotes the generation of voids during the punching of hot-rolled steel sheets. Moreover, since S becomes a starting point of generation of voids during processing, it reduces stretch flangeability. Therefore, it is preferable to reduce the S content as much as possible, and the content is 0.03% or less. Preferably it is 0.01% or less. However, since excessive desulfurization leads to an increase in refining time and cost, the S content is preferably 0.0002% or more.
  • Al 0.1% or less
  • Al is an element that acts as a deoxidizer. In order to obtain such an effect, it is preferable to contain 0.01% or more of Al. However, if Al exceeds 0.1%, it remains as an Al oxide in the steel sheet, and the Al oxide tends to aggregate and become coarse and deteriorate stretch flangeability. Therefore, the Al content is 0.1% or less.
  • N 0.01% or less N is present as coarse TiN in the steel, and promotes the generation of coarse voids during the punching of hot-rolled steel sheets. Moreover, since N becomes a starting point of generation of coarse voids during processing, it reduces stretch flangeability. For this reason, it is preferable to reduce N content as much as possible, and set it as 0.01% or less. Preferably it is 0.006% or less. However, excessive N removal leads to an increase in refining time and cost, so the N content is preferably 0.0005% or more.
  • Ti 0.03% or more and 0.15% or less
  • Ti is an indispensable element for increasing the strength of a steel sheet by forming MC-type carbides and suppressing lath coarsening in the annealing process.
  • MC type carbide increases the strength of the steel sheet by precipitation strengthening.
  • the Ti content is set to 0.03% or more and 0.15% or less.
  • they are 0.04% or more and 0.14% or less, More preferably, they are 0.05% or more and 0.13% or less.
  • V 0.01% or more and 0.3% or less
  • Nb 0.01% or more and 0.1% or less
  • Mo 0.01 as necessary for the purpose of further increasing the strength. % Or more and 0.3% or less can be contained.
  • V 0.01% or more and 0.3% or less
  • V forms MC-type carbides and, like Ti, contributes to increasing the strength of the steel sheet by suppressing coarsening of the lath and precipitation strengthening in the annealing process.
  • the V content is preferably 0.01% or more and 0.3% or less.
  • V may form MC-type carbides alone, or may form composite carbides with Ti, Nb, and Mo. These carbide compositions have no influence on the effects of the invention.
  • Nb 0.01% or more and 0.1% or less Nb forms MC-type carbides and, like Ti, contributes to increasing the strength of the steel sheet by suppressing lath coarsening and precipitation strengthening in the annealing process.
  • Nb it is necessary to contain 0.01% or more of Nb.
  • the Nb content is preferably 0.01% or more and 0.1% or less.
  • Nb may form MC-type carbides alone, or may form composite carbides with Ti, V, and Mo. These carbide compositions have no influence on the effects of the invention.
  • Mo 0.01% or more and 0.3% or less Mo forms MC-type composite carbide by compound addition with Ti, and, like Ti, increases the strength of the steel sheet by suppressing lath coarsening and precipitation strengthening in the annealing process. Contribute. In order to obtain such an effect, it is necessary to contain 0.01% or more of Mo. On the other hand, if the Mo content exceeds 0.3%, the central segregation becomes prominent and causes punching deterioration. Therefore, the Mo content is preferably 0.01% or more and 0.3% or less. Mo may form a composite carbide with Nb or V, but these carbide compositions do not affect the effects of the invention.
  • the hot-rolled steel sheet of the present invention may contain B: 0.0002% to 0.010% as necessary for the purpose of improving the hardenability during hot rolling.
  • B 0.0002% or more and 0.010% or less
  • B is an element that segregates at austenite grain boundaries and suppresses the formation and growth of ferrite to improve hardenability and promote the formation of bainite and martensite.
  • the B content is preferably 0.0002% or more.
  • the B content exceeds 0.010%, a hard iron boride is formed, which causes stretch flangeability deterioration. Therefore, when it contains B, it is preferable to make the content into 0.0002% or more and 0.010% or less. Further, it is more preferably 0.0002% or more and 0.0050% or less, and further preferably 0.0004% or more and 0.0030% or less.
  • the hot-rolled steel sheet of the present invention further includes REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn, and the above composition.
  • One or more of Cs may be contained in a total of 1.0% or less.
  • Components other than the above are Fe and inevitable impurities.
  • Total area ratio of tempered bainite phase and tempered martensite phase 70% or more
  • the hot-rolled steel sheet of the present invention has a structure mainly composed of tempered bainite and tempered martensite, which has both high strength and excellent punchability.
  • tempered bainite phase and the tempered martensite phase are less than 70%, a hot rolled steel sheet having desired high strength and punchability cannot be obtained.
  • the reason why the tempered bainite phase and the tempered martensite fraction are not individually defined is that the tempered bainite and tempered martensite after annealing have a structure that cannot be distinguished.
  • the total area ratio of the tempered bainite phase and the tempered martensite phase is 70% or more. Preferably it is 75% or more, more preferably 80% or more. Further, the total area ratio of the tempered bainite phase and the tempered martensite phase may be 100%.
  • the hot-rolled steel sheet of the present invention has a structure mainly composed of tempered bainite and tempered martensite.
  • the remaining structure other than martensite include Fe carbide, coarse pearlite, fine pearlite, pseudo pearlite, bainite, martensite, and retained austenite.
  • these structures especially when coarse pearlite, martensite, and retained austenite are present in the metal structure, stretch flangeability is significantly deteriorated. Therefore, the sum of the area ratios of the coarse pearlite phase, martensite phase and residual austenite phase is set to 10% or less.
  • coarse pearlite has a lamella spacing of 0.2 ⁇ m or more
  • fine pearlite has a lamella spacing of less than 0.2 ⁇ m
  • pseudo pearlite does not clearly observe a pearlite lamella.
  • the lamella spacing can be obtained by observing the structure with a scanning electron microscope.
  • examples of the remaining structure other than the tempered bainite phase, the tempered martensite phase, the coarse pearlite phase, the martensite phase, and the retained austenite phase include a ferrite phase, a pseudo pearlite phase, and a fine pearlite phase. Such a remaining structure is acceptable if the total area ratio is 30% or less.
  • FIG. 1 is a schematic diagram showing an example of a structure in which a tempered bainite phase and a tempered martensite phase have lath as a substructure, and Fe-based carbides precipitate and MC-type carbides disperse and precipitate in the lath inside and lath boundaries. The figure is shown.
  • the average width of the laths that the tempered bainite phase and the tempered martensite phase have as a substructure is 1.0 ⁇ m or less.
  • it is 0.8 micrometer or less, More preferably, it is 0.6 micrometer or less.
  • the lower limit is not particularly limited, but is usually about 0.1 ⁇ m.
  • Fe-based carbides deposited in the lath and lath boundary are the origin of voids during punching This contributes to improved punchability.
  • an Fe-based carbide having an aspect ratio of 5 or less has a large effect, and an excellent punchability can be exhibited by setting the ratio to 80% or more. Therefore, the proportion of Fe-based carbides precipitated in the lath and at the lath boundary is set to 80% or more with an aspect ratio of 5 or less. Preferably it is 85% or more.
  • the upper limit is not particularly limited and may be 100%.
  • the Fe-based carbide is ⁇ carbide (cementite) or ⁇ carbide.
  • the alloy element may be dissolved in the carbide.
  • the aspect ratio is the ratio of the length of the major axis to the minor axis of the Fe-based carbide precipitated in the lath and at the lath boundary.
  • MC type carbide dispersed and precipitated in the lath and lath boundary 20 nm or less
  • MC type carbide finely dispersed and precipitated in the lath and lath boundary has a pinning effect during annealing of the steel sheet.
  • the average particle diameter of MC type carbide exceeds 20 nm, the number of MC type carbide particles contributing to pinning is insufficient, the pinning effect is insufficient, and the steel sheet strength is reduced.
  • the average particle diameter of MC type carbides dispersed and precipitated in the lath and lath boundaries of the tempered bainite phase and the tempered martensite phase is set to 20 nm or less. Preferably it is 15 nm or less.
  • the lower limit is not particularly limited, but is usually about 1 nm.
  • the proportion of MC type carbide having a particle size exceeding 50 nm is preferably 10% or less.
  • the average dislocation density is in the following range.
  • Average dislocation density 1.0 ⁇ 10 14 m ⁇ 2 or more and 5.0 ⁇ 10 15 m ⁇ 2 or less
  • the steel sheet having a bainite and martensite structure is tempered to vary the variation in hot rolling conditions. Is reduced.
  • the average dislocation density of the steel sheet after annealing exceeds 5.0 ⁇ 10 15 m ⁇ 2 , the tempering of the steel sheet is insufficient, and the influence of fluctuations in hot rolling conditions cannot be sufficiently mitigated.
  • the average dislocation density is usually 1.0 ⁇ 10 14 m ⁇ 2 or more.
  • the average dislocation density is 1.0 ⁇ 10 14 m ⁇ 2 or more and 5.0 ⁇ 10 15 m ⁇ 2 or less.
  • it is 1.0 ⁇ 10 14 m ⁇ 2 or more and 2.0 ⁇ 10 15 m ⁇ 2 or less.
  • the method for producing a hot-rolled steel sheet of the present invention was obtained after heating the steel material having the above-described composition to the austenite single-phase region, subjecting it to hot rolling consisting of rough rolling and finish rolling, and finishing the finish rolling.
  • a hot rolling process for cooling and winding the steel sheet After the hot rolling step, pickling the steel sheet, and then having a continuous annealing step for continuous annealing,
  • the finish rolling temperature is 850 ° C. or more and 1000 ° C. or less
  • the average cooling rate to 500 ° C. is 30 ° C./s or more
  • the winding temperature is 500 ° C.
  • the maximum heating temperature of the steel sheet is 700 ° C. or more (A 3 points + A 1 point) / 2 or less
  • the time when the steel plate temperature when heating the steel plate to the maximum heating temperature is 600 ° C. or more and 700 ° C. or less is 20 seconds or more and 1000 seconds or less
  • the time when the steel sheet temperature is over 700 ° C. is set to 200 seconds or less
  • the melting method of a steel raw material is not specifically limited, Well-known melting methods, such as a converter and an electric furnace, are employable. Moreover, after melting, it is preferable to use a continuous casting method to form a slab (steel material) from the viewpoint of productivity and the like, but as a slab by a known casting method such as an ingot-bundling rolling method or a thin slab continuous casting method. Also good.
  • the steel material obtained as described above is subjected to hot rolling consisting of rough rolling and finish rolling, and the steel material is heated in the austenite single phase region prior to rough rolling. If the steel material before rough rolling is not heated in the austenite single-phase region, remelting of Ti carbide, etc. present in the steel material will not proceed, and MC type carbide will precipitate finely during annealing after hot rolling. I will not. Therefore, prior to rough rolling, the steel material is heated to an austenite single phase region, preferably 1150 ° C. or higher.
  • the upper limit of the heating temperature is not particularly specified, but when the heating temperature becomes higher than necessary, the yield decreases due to oxidation of the slab surface, so the heating temperature is usually 1350 ° C. or lower.
  • the steel material is not heated or directly heated after being heated for a short time. You may roll.
  • Finish rolling temperature 850 ° C. or more and 1000 ° C. or less
  • the temperature on the finish rolling delivery side needs to be 850 ° C. or higher, preferably 880 ° C. or higher.
  • the finish rolling temperature exceeds 1000 ° C., the surface properties of the steel sheet deteriorate.
  • the upper limit of finish rolling temperature shall be 1000 degrees C or less. Preferably it is 970 degrees C or less.
  • Each temperature such as the coiling temperature including the above finish rolling temperature is the temperature of the steel sheet surface.
  • cooling rate to 500 ° C 30 ° C / s or more
  • the cooling rate to 500 ° C. after finishing rolling needs to be 30 ° C./s or more.
  • it is 50 ° C./s or more.
  • the upper limit of the cooling rate is not particularly limited, but is usually about 300 ° C./s.
  • Winding temperature 500 ° C. or less Optimization of the winding temperature is important for controlling the steel sheet structure after hot rolling.
  • the coiling temperature exceeds 500 ° C., the lath width of the bainite becomes large, so that the lath width of the tempered bainite after annealing cannot be set to a predetermined value.
  • the lower limit of the coiling temperature is not particularly limited, but if the coiling temperature is excessively lowered, the cooling cost is only unnecessarily expensive.
  • the winding temperature is preferably 0 ° C. or higher. More preferably, it is 200 ° C. or higher.
  • the hot rolled steel sheet is pickled and subjected to a continuous annealing step for continuous annealing.
  • Maximum heating temperature of steel sheet 700 ° C or higher (A 3 points + A 1 point) / 2 or lower Optimization of the maximum heating temperature of steel sheets in the continuous annealing process is sufficient for the effects of manufacturing conditions fluctuations during hot rolling due to annealing It is important to reduce and achieve the desired high strength. If the maximum heating temperature of the steel sheet is less than 700 ° C, it is difficult to control the dislocation density in bainite and martensite within an appropriate range. For this reason, the influence of fluctuations in manufacturing conditions during hot rolling is sufficiently reduced.
  • the heating temperature of the steel sheet is less than 700 ° C, the aspect ratio of Fe-based carbides inside and between the laths tends to be large, and the proportion of Fe-based carbides with an aspect ratio of 5 or less should be in the desired range Is difficult.
  • the maximum heating temperature of the steel sheet exceeds (A 3 points + A 1 point) / 2, the MC type carbides become prominently coarsened, so that the lath coarsening in bainite and martensite can be sufficiently suppressed. Disappear. Further, by promoting austenitization, the bainite and martensite fractions are lowered, and the desired tempered bainite and tempered martensite fractions cannot be obtained.
  • the maximum heating temperature of the steel sheet in the continuous annealing process is 700 ° C. or more (A 3 points + A 1 point) / 2 or less.
  • the temperature is preferably 700 ° C. or higher and ⁇ (A 3 points + A 1 point) / 2 ⁇ ⁇ 10 ° C. or lower.
  • the points A 1 and A 3 can be calculated by the following formula.
  • a 1 point 751-26.6 x [% C] + 17.6 x [% Si]-11.6 x [% Mn] + 22.5 x [% Mo] + 233 ⁇ [% Nb] ⁇ 39.7 ⁇ [% V] ⁇ 57 ⁇ [% Ti] ⁇ 895 ⁇ [% B] ⁇ 169 ⁇ [% Al]
  • a 3 points 937-476.5 x [% C] + 56 x [% Si]-19.7 x [% Mn] + 38.1 x [% Mo] + 124.8 ⁇ [% V] + 136.3 ⁇ [% Ti] ⁇ 19 ⁇ [% Nb] + 3315 ⁇ [% B]
  • [% X] means the content of X element in steel (mass%).
  • Time when the steel plate temperature is 600 ° C or higher and 700 ° C or lower when the steel plate is heated to the maximum heating temperature 20 seconds or more and 1000 seconds or less It is desirable to appropriately control the heating heat history when heating the steel plate to the maximum heating temperature This is important for imparting high strength and excellent punchability to the steel sheet.
  • the pinning effect by MC type carbide is used to suppress the coarsening of the lath. In order to exhibit this pinning effect, it is necessary to sufficiently disperse MC type carbides in bainite and martensite before the lath starts coarsening. According to the study by the present inventors, precipitation of MC type carbide begins to occur remarkably at 600 ° C. or higher.
  • the coarsening and disappearance of lath occurs remarkably when the temperature exceeds 700 ° C. Therefore, the coarsening and disappearance of the lath can be suppressed by maintaining the steel plate temperature in the temperature range of 600 ° C. or more and 700 ° C. or less for a certain period of time and sufficiently depositing the MC type carbide.
  • the holding time in this temperature range is insufficient, since the coarsening of the lath is started before the MC type carbide is sufficiently precipitated, the pinning effect is not sufficiently exhibited and the lath becomes coarse. .
  • the holding time in the temperature range of 600 ° C. or more and 700 ° C. or less exceeds 1000 seconds, the Fe-based carbides precipitated inside and between the laths re-dissolve, and the prior austenite grain boundaries and packet grain boundaries In other words, there is no Fe-based carbide in the lath and between the laths, which moves to the block grain boundaries and effectively contributes to the improvement of punchability. Therefore, in order to obtain a steel sheet having excellent punchability, the holding time in the temperature range where the steel sheet temperature is 600 ° C. or higher and 700 ° C. or lower needs to be 1000 seconds or shorter. Preferably it is 800 seconds or less, more preferably 500 seconds or less.
  • the steel plate temperature here is the temperature of the steel plate surface.
  • Time when the steel plate temperature exceeds 700 ° C .: 200 seconds or less In the temperature range where the steel plate temperature exceeds 700 ° C., lath coarsening occurs remarkably. As described above, in the present invention, the movement of the lath boundary is suppressed and the coarsening of the lath is suppressed by the pinning effect by the MC type carbide finely dispersed and precipitated. And thereby, the steel plate strength is maintained. However, if the holding time in this temperature range becomes excessively long, the coarsening of the lath cannot be suppressed. For this reason, from the viewpoint of preventing the coarsening of the lath, the holding time in the temperature range where the steel plate temperature exceeds 700 ° C. is set to 200 seconds or less.
  • Preferably it is 180 seconds or less, more preferably 150 seconds or less.
  • the time when the steel plate temperature is higher than 700 ° C. is less than 10 seconds, the ductility of the steel plate is somewhat inferior, and therefore it is preferable to set it to 10 seconds or more.
  • Average cooling rate up to 530 ° C when cooling the steel plate from the maximum heating temperature 8 ° C / s or more and 25 ° C / s or less.
  • Controlling is important for obtaining excellent stretch flangeability.
  • the average cooling rate up to 530 ° C. is lower than 8 ° C./s, pearlite transformation cannot be suppressed during cooling, and coarse pearlite is produced in a predetermined amount or more. For this reason, stretch flangeability falls.
  • the average cooling rate is excessively high, it is difficult to maintain for a predetermined time in a temperature range of 470 ° C. to 530 ° C., which will be described later. Therefore, the average cooling rate up to 530 ° C. when cooling the steel plate from the maximum heating temperature is 25 ° C./s or less.
  • Holding time in the temperature range of 470 ° C or higher and 530 ° C or lower 10 seconds or longer
  • the steel plate is controlled and cooled as described above, and then held in the temperature range of 470 ° C or higher and 530 ° C or lower for a certain period of time. It is important to obtain excellent stretch flangeability.
  • the holding temperature after the cooling is stopped exceeds 530 ° C., coarse pearlite is generated, so that stretch flangeability is deteriorated.
  • the holding temperature after stopping cooling is lower than 470 ° C., the transformation from austenite to bainite is delayed.
  • C concentrates in the untransformed austenite region and stabilizes the austenite, so the transformation is not completed. Then, in the subsequent cooling, untransformed austenite is transformed into martensite or remains in the steel sheet structure as retained austenite, so that stretch flangeability is lowered. Further, when the steel plate is held for 10 seconds or more in the temperature range of 470 ° C. or more and 530 ° C. or less, the transformation of most austenite to bainite is completed, and then the martensite fraction generated when cooled is determined to be a predetermined value. Can be reduced to a range. Accordingly, after the controlled cooling is stopped, the holding time in the temperature range from 470 ° C. to 530 ° C.
  • the subsequent cooling conditions are not particularly limited, and may be cooled to room temperature by any cooling method.
  • the total holding time in the temperature range of 600 ° C to 700 ° C is 1000 seconds or less.
  • the steel plate may be immersed in a zinc pot to be a hot dip galvanized steel plate, or may be further heat-treated to obtain an alloyed hot dip galvanized steel plate. .
  • aluminum or aluminum alloy can be plated for hot dipping.
  • temper rolling may be applied to the steel sheet continuously in the annealing line or using another line according to a conventional method.
  • the hot-rolled steel sheet produced as described above may be separately subjected to electrogalvanizing treatment or hot dip galvanizing.
  • the hot-rolled steel sheet of the present invention is suitable as a steel sheet for automobile undercarriage, and is suitable for press forming performed at normal room temperature, and has excellent heat treatment characteristics.
  • the hot-rolled steel sheet produced as described above is also suitable as a warm-formed material steel sheet that is immediately press-formed after heating the steel sheet before pressing from 400 ° C to 700 ° C.
  • the molten steel having the composition shown in Table 1 was melted and continuously cast by a generally known method to obtain a slab (steel material) having a thickness of 300 mm. These slabs are heated to the temperatures shown in Table 2, roughly rolled, and finish rolling is completed at the temperatures shown in Table 2. After finishing rolling, the slabs are cooled at the average cooling rate shown in Table 2, and the windings shown in Table 2 are used. The coil was wound at the coiling temperature to obtain a hot rolled steel sheet having a thickness of 3.2 mm. Furthermore, these hot-rolled steel sheets were pickled by a generally known technique and annealed under the conditions shown in Table 2 in a continuous annealing line. Moreover, about some steel plates, the hot dip galvanization process and also the alloying process were performed in the continuous annealing line, and it was set as the hot dip galvanized steel plate and the galvannealed steel plate.
  • Test specimens were collected from the hot-rolled steel sheets thus obtained, and subjected to structure observation, measurement of average dislocation density, tensile test, hole expansion test, punching test, and production stability evaluation. The evaluation results are shown in Table 3.
  • the test method was as follows.
  • the area ratio of the sum of the martensite phase and the retained austenite phase was determined without particularly distinguishing between the martensite phase and the retained austenite phase.
  • the thin film produced from the hot-rolled steel sheet is observed with a transmission electron microscope, and the lath width of tempered bainite and tempered martensite is measured.
  • the average particle size of the MC type carbides precipitated inside the lath and at the lath boundary was determined.
  • transmission electron micrographs of 120 mm x 80 mm in size taken at 10 fields of view at a magnification of 30000 times were measured for three or more consecutive laths.
  • Draw 5 straight lines at an interval of 10mm perpendicular to the long axis measure the length of each line segment that intersects the lath boundary, and calculate the average length of the obtained line segments as the average lath width. It was.
  • the proportion of Fe-based carbides deposited inside the lath and at the lath boundary with an aspect ratio of 5 or less is a minimum of 100 in the total of 5 fields deposited inside the lath and at the lath boundary using photographs taken at a magnification of 165000 times.
  • the ratio of those having an aspect ratio of 5 or less was determined.
  • the average particle size of MC type carbide was measured using a photograph taken at a magnification of 300,000 times, and the diameter of MC type carbides such as TiC of at least 100 total in 5 fields was measured.
  • the diameter d def was obtained.
  • the lower limit of the measured particle diameter is 2 nm.
  • the measurement chart uses an X-ray diffractometer to measure the diffraction intensity of the (110), (211), and (220) planes of 1 / 4-thick ⁇ -iron using CoK ⁇ rays. Then, the half width of the peak value of the reflection intensity of each crystal plane is obtained, and the local strain ⁇ ′ applied to the steel sheet is determined by the following equations (1) and (2).
  • Hole expansion ratio ⁇ (%) ⁇ (d ⁇ d 0 ) / d 0 ⁇ ⁇ 100
  • TS tensile strength
  • ⁇ hole expansion rate
  • JIS No. 5 tensile test piece JIS Z 2001
  • C direction perpendicular direction
  • TS perpendicular direction
  • ⁇ perpendicular direction

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Abstract

Cette invention présente une composition spécifiée en composants et elle est façonnée en une structure dans laquelle : la surface totale d'une phase bainitique revenue et d'une phase martensitique revenue est supérieure ou égale à 70 % ; la surface totale d'une phase perlitique grossière, d'une phase martensitique et d'une phase austénitique résiduelle est inférieure ou égale à 10 % ; la phase bainitique revenue et la phase martensitique revenue ont une latte de largeur moyenne inférieure ou égale à 1,0 μm sous forme de structure auxiliaire ; la proportion de matériau ayant un rapport d'aspect de 5 ou moins dans un carbure de fer précipité à l'intérieur de la latte et au niveau des limites de la latte est supérieure ou égale à 80 % ; et un carbure de type MC ayant un diamètre moyen de particule de 20 nm ou moins est dispersé et précipité à l'intérieur de la latte et au niveau des limites de la latte. En outre, la densité de dislocation moyenne est de 1,0 × 1014m-2 à 5,0 × 1015m-2 inclus.
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