EP3495523A1 - Hochfeste stahlplatte und herstellungsverfahren dafür - Google Patents

Hochfeste stahlplatte und herstellungsverfahren dafür Download PDF

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Publication number
EP3495523A1
EP3495523A1 EP17836781.9A EP17836781A EP3495523A1 EP 3495523 A1 EP3495523 A1 EP 3495523A1 EP 17836781 A EP17836781 A EP 17836781A EP 3495523 A1 EP3495523 A1 EP 3495523A1
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Prior art keywords
mass
temperature
less
cooling
amount
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EP17836781.9A
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English (en)
French (fr)
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EP3495523A4 (de
EP3495523B1 (de
Inventor
Elijah KAKIUCHI
Toshio Murakami
Shigeo Otani
Yuichi Futamura
Tadao Murata
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Kobe Steel Ltd
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Kobe Steel Ltd
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • 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
    • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
    • 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
    • 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
    • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—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

Definitions

  • the present disclosure relates to a high-strength sheet that can be used in various applications including automobile parts.
  • steel sheets used for automobile parts and the like are required to achieve both improvement in strength and improvement in impact resistance properties.
  • Patent Document 1 discloses a high-strength steel sheet in which an attempt is made to improve impact resistance properties by heating a slab to 1,210°C or higher and controlling the hot-rolling conditions to form fine TiN particles having a size of 0.5 ⁇ m or less, thereby suppressing the formation of AlN particles having a particle size of 1 ⁇ m or more that act as a starting point of low temperature fracture.
  • Patent Document 2 discloses a high-strength sheet in which an attempt is made to improve collision resistance properties by forming a network structure in which 50% or more of a ferrite grain size is in contact with a hard phase while defining the C amount to more than 0.45% and 0.77% or less, the Mn amount to 0.1% or more and 0.5% or less and the Si amount to 0.5% or less, and defining each addition amount of Cr, Al, N and O.
  • Patent Document 3 discloses a high-strength sheet in which an attempt is made to improve collision resistance properties by adding 3.5 to 10% of Mn, thereby adjusting the amount of retained austenite to 10% or more and an average interval of retained austenite to 1.5 ⁇ m or less.
  • Patent Document 4 discloses a high-strength sheet that has a tensile strength of 980 to 1,180 MPa and also exhibits satisfactory deep drawability.
  • steel sheets used for automobile parts are required to have sufficient strength and impact resistance properties while being made thinner.
  • steel sheets having higher tensile strength and excellent impact properties are required.
  • steel sheets are required to have not only high tensile strength and impact properties, but also excellent strength-ductility balance, high yield ratio, stretch formability and excellent hole expansion ratio.
  • the followings are required for each of the tensile strength, the strength-ductility balance, the yield ratio, the deep drawing properties and the hole expansion ratio.
  • the tensile strength is required to be 980 MPa or higher.
  • high yield strength (YS), in addition to high tensile strength (TS).
  • TS high tensile strength
  • a joint strength of the spot welded portion is also required as basic performances of the steel sheet for automobiles.
  • a cross tensile strength of the spot welded portion is required to be 6 kN or higher.
  • the product (TS ⁇ EL) of TS and total elongation (EL) is required to be 20,000 MPa% or higher.
  • the hole expansion ratio ⁇ showing hole expansion properties is 20% or more and the maximum forming height (forming height) showing stretch formability is 16 mm or more.
  • Patent Documents 1 to 4 it is difficult for the high-strength sheets disclosed in Patent Documents 1 to 4 to satisfy all of these requirements, and there has been required a high-strength steel sheet that can satisfy all of these requirements.
  • the embodiment of the present invention has been made to respond to these requirements, and it is an object thereof to provide a high-strength sheet in which all of tensile strength (TS), yield ratio (YR), the product (TS ⁇ EL) of (TS) and total elongation (EL), hole expansion ratio ( ⁇ ), thickness reduction ratio (RA) of the fracture portion during a tensile test, maximum forming height and cross tensile strength (SW cross tension) of the spot welded portion are at a high level, and a manufacturing method thereof.
  • TS tensile strength
  • YiR yield ratio
  • TS ⁇ EL product
  • ⁇ total elongation
  • ⁇ hole expansion ratio
  • RA thickness reduction ratio
  • SW cross tension maximum forming height and cross tensile strength
  • Aspect 1 of the present invention provides a high-strength sheet including:
  • Aspect 2 of the present invention provides the high-strength sheet according to aspect 1, in which the C amount is 0.30% by mass or less.
  • Aspect 3 of the present invention provides the high-strength sheet according to aspect 1 or 2, in which the Al amount is less than 0.10% by mass.
  • Aspect 4 of the present invention provides a method for manufacturing a high-strength sheet, which includes:
  • Aspect 5 of the present invention provides the manufacturing method according to aspect 4, wherein cooling to the cooling stopping temperature includes:
  • Aspect 6 of the present invention provides the manufacturing method according to aspect 4 or 5, in which the tempering parameter is 11,000 to 14,000 and the holding time is 1 to 150 seconds.
  • TS tensile strength
  • YiR yield ratio
  • TS ⁇ EL product of (TS)
  • EL total elongation
  • ⁇ hole expansion ratio
  • RA thickness reduction ratio
  • Fig. 1 is a diagram explaining a method for manufacturing a high-strength sheet according to the embodiment of the present invention, especially a heat treatment.
  • TS tensile strength
  • YiR yield ratio
  • TS ⁇ EL product of (TS)
  • EL total elongation
  • ⁇ thickness reduction ratio
  • SW cross tension maximum forming height and cross tensile strength
  • MA cross tension
  • MA has an average size of 1.0 ⁇ m or less
  • a half-width of the concentration distribution of Mn in the carbon-concentrated region that is equal to the amount of retained austenite is 0.3% by mass or more
  • the high-strength sheet according to the embodiment of the present invention has the Mn-concentrated region formed by holding in a two-phase coexistence region intermediate between an Ac 1 point and an Ac 3 point, more specifically at a temperature between the Ac 1 point and 0.2 ⁇ Ac 1 point + 0.8 ⁇ Ac 3 point for a predetermined time, followed by holding at a temperature of the Ac 3 point or higher for a predetermined time, in the austenitizing step of a heat treatment during manufacturing.
  • the carbon-concentrated region corresponding to retained austenite (the same amount as that of retained austenite) is formed. In this carbon-concentrated region, both the Mn-concentrated region and the Mn-nonconcentrated region are formed.
  • the region containing a large amount of Mn and the region that does not contain a large amount of Mn exist in some carbon-enriched regions (retained austenite). Therefore, when the distribution of the Mn concentration is measured with respect to the entire carbon-concentrated region (i.e., corresponding to the entire retained austenite), there is a certain degree or more of variation in Mn concentration. Specifically, the half-width of the concentration distribution of Mn becomes 0.3% by mass or more.
  • Variation in amount of Mn contained in retained austenite means that it is possible to be provided with retained austenite with various degrees of stability. Therefore, residual austenite with low stability that causes strain induced transformation at a relatively small amount of strain and residual austenite with high stability that causes strain induced transformation at a large amount of strain coexist, thus making it possible to cause strain induced transformation in various strain regions. As a result, the n value can be increased in a wide strain region to enhance strain dispersibility, thus enabling realization of high stretch formability.
  • the high-strength sheet according to the embodiment of the present invention and a manufacturing method thereof will be described in detail below.
  • Ferrite generally has excellent workability but has a problem such as low strength. As a result, a large amount of ferrite leads to a decrease in yield ratio. Therefore, a ferrite fraction was set at 5% or less (5 volume % or less).
  • the ferrite fraction is preferably 3% or less, and more preferably 1% or less.
  • the ferrite fraction can be determined by observing with an optical microscope and measuring the white region by the point counting method. By such a method, it is possible to determine the ferrite fraction by an area ratio (area %). The value obtained by the area ratio may be directly used as the value of the volume ratio (volume %).
  • the total fraction of tempered martensite and tempered bainite is preferably 70% or more.
  • the retained austenite causes the TRIP phenomenon of being transformed into martensite due to strain induced transformation during working such as press working, thus making it possible to obtain large elongation. Martensite thus formed has high hardness. Therefore, excellent strength-ductility balance can be obtained.
  • the amount of retained austenite is set at 10% or more (10 volume % or more), it is possible to realize TS ⁇ EL of 20,000 MPa% or more and excellent strength-ductility balance.
  • the amount of retained austenite is preferably 15% or more.
  • MA is abbreviation of a martensite-austenite constituent and is a composite (complex structure) of martensite and austenite.
  • ferrite including tempered martensite and untempered martensite in X-ray diffraction
  • austenite by X-ray diffraction
  • Co-K ⁇ ray can be used as an X-ray source.
  • MA is a hard phase and the vicinity of matrix/hard phase interface acts as a void forming site during deformation.
  • the average size of MA is preferably 0.8 ⁇ m or less.
  • carbon-concentrated region that is equal to the amount of retained austenite means the region that corresponds to retained austenite.
  • the half-width of the concentration distribution of Mn is 0.3% by mass or more, preferably 0.5% by mass or more, more preferably 0.6% by mass or more, and still more preferably 0.75% by mass or more.
  • retained austenite carbon-enriched region
  • Retained austenite with low stability causes strain induced transformation at a small amount of strain, thus turning to martensite.
  • Retained austenite with high stability does not cause strain induced transformation at a small amount of strain, and causes strain induced transformation, thus turning martensite only after a large amount of strain is applied.
  • retained austenite having a wide range of stability exists, retained austenite continuously causes strain induced transformation from the time when a small amount of strain is applied immediately after the start of forming to the time when forming progresses and a large amount of strain is applied.
  • the n value can be increased over a wide range of strain to enhance the strain dispersibility, thus enabling realization of high stretch formability.
  • X-ray small angle scattering means that the size distribution of fine particles (e.g., cementite particles dispersed in a steel sheet) contained in the steel sheet can be obtained by irradiating the steel sheet with X-rays and measuring scattering of X-rays transmitted through the steel sheet.
  • fine particles e.g., cementite particles dispersed in a steel sheet
  • X-ray small angle scattering it is possible to analyze the size and the fraction of cementite particles using the q value and the scattering intensity.
  • the q value is an index of the size of particles (e.g., cementite particles) in the steel sheet.
  • the "q value of 1 nm -1 " corresponds to cementite particles having a particle size of about 1 nm.
  • the scattering intensity is an index of the volume fraction of particles (e.g., cementite particles) in the steel sheet. The larger the scattering intensity, the larger the volume fraction of cementite becomes.
  • the scattering intensity at a certain q value semi-quantitatively indicates the volume fraction of cementite particles of the size corresponding to the q value.
  • the scattering intensity at the q value of 1 nm -1 semi-quantitatively indicates the volume fraction of fine cementite particles having a size of about 1 nm.
  • large scattering intensity at the q value of 1 nm -1 indicates large volume fraction of fine cementite particles having a size of about 1 nm.
  • the scattering intensity at the q value of 1 nm -1 is 1.0 cm -1 or less
  • the volume fraction of fine cementite particles having a size of about 1 nm existing in the steel sheet is a predetermined value (the value corresponding to the scattering intensity of 1.0 cm -1 ) or less.
  • the steel sheet in which "the scattering intensity at the q value of 1 nm -1 is 1.0 cm -1 or less" is excellent in collision resistance properties since the volume fraction of cementite having a size of about 1 nm is suppressed to a low value.
  • the steel sheet according to the embodiment of the present invention by suppressing the volume fraction of fine cementite to a low value, more specifically, by setting the scattering intensity at the q value of 1 nm -1 at 1 cm -1 or less, fine carbide formed in laths of tempered martensite is reduced to enhance the deformability in martensite. Thus, fracture of the steel sheet upon collision is suppressed to improve collision resistance properties of the steel sheet.
  • X-ray small angle scattering was measured using a Nano-viewer, Mo tube manufactured by Rigaku Corporation. A 3 mm ⁇ disk-shaped sample was cut out from the steel sheet and samples having a thickness of 20 ⁇ m were cut out from the vicinity of the thickness of 1/4 and then used. Data at the q value of 0.1 to 10 nm -1 were collected. Among them, absolute intensity was determined for the q value of 1 nm -1 .
  • steel structures other than the above-mentioned ferrite, tempered martensite, tempered bainite retained austenite and cementite are not specifically defined.
  • pearlite, untempered bainite, untempered martensite and the like may exist, in addition to the steel structures such as ferrite.
  • the steel structure such as ferrite satisfies the above-mentioned structure conditions, the effects of the present invention are exhibited even if pearlite or the like exists in the steel.
  • composition of the high-strength sheet according to the embodiment of the present invention will be described below. Main elements C, Si, Al, Mn, P and S will be described. Note that all percentages as unit with respect to the composition are by mass.
  • Carbon (C) is an element indispensable for ensuring properties such as high strength-ductility balance (TS ⁇ EL balance) by increasing the amount of desired structure, especially retained ⁇ . In order to effectively exhibit such effect, there is a need to add C in the amount of 0.15% or more. However, the amount of more than 0.35% is not suitable for welding. The amount is preferably 0.18% or more, and more preferably 0.20% or more. The amount is preferably 0.30% or less. If the C amount is 0.25% or less, welding can be easily performed.
  • Si and Al each have the effect of suppressing the precipitation of cementite, thus remaining retained austenite. In order to effectively exhibit such effect, there is a need to add Si and Al in the total amount of 0.5% or more. If the total amount of Si and Al exceeds 3.0%, the deformability of the steel is degraded, thus degrading TS ⁇ EL and forming height.
  • the total amount is preferably 0.7% or more, and more preferably 1.0% or more.
  • the total amount is preferably 2.5% or less.
  • Al may be added in the amount enough to function as an deoxidizing element, i.e., less than 0.10% by mass.
  • Al may be added in a larger amount of 0.7% by mass or more.
  • Mn suppresses the formation of ferrite.
  • Mn is an element indispensable for improving the stretch formability by forming the Mn-concentrated region and forming retained austenites with different stabilities.
  • the amount is preferably 1.5% or more, and more preferably 2.0% or more.
  • the amount is preferably 3.5% or less.
  • the content of P is set at 0.05% or less (including 0%).
  • the content is 0.03% or less (including 0%).
  • S inevitably exists as an impurity element. If more than 0.01% of S exists, sulfide-based inclusions such as MnS are formed and act as a starting point of cracking, thus degrading ⁇ . Therefore, the content of S is set at 0.01% or less (including 0%). The content is preferably 0.005% or less (including 0%).
  • the balance is composed of iron and inevitable impurities. It is permitted to mix, as inevitable impurities, trace elements (e.g., As, Sb, Sn, etc.) incorporated according to the conditions of raw materials, materials, manufacturing facilities and the like. There are elements whose content is preferably as small as possible, like P and S, that are therefore inevitable impurities in which the composition range is separately defined as mentioned above. Therefore, "inevitable impurities" constituting the balance as used herein means the concept excluding elements whose composition range is separately defined.
  • trace elements e.g., As, Sb, Sn, etc.
  • any other element may be further included.
  • the high-strength sheet has TS of 980 MPa or higher.
  • TS is 1,180 MPa or higher. If TS is lower than 980 MPa, excellent fracture properties can be more surely obtained, but it is not preferable since withstand load upon collision decreases.
  • the high-strength sheet has an yield ratio of 0.75 or more. This makes it possible to realize a high yield strength combined with the above-mentioned high tensile strength and to use the final product obtained by working such as deep drawing under high stress.
  • the high-strength sheet has a yield ratio of 0.80 or more.
  • TS ⁇ EL is 20,000 MPa% or more. By having TS ⁇ EL of 20,000 MPa% or more, it is possible to obtain high-level strength-ductility balance that has both high strength and high ductility. Preferably, TS ⁇ EL is 23,000 MPa% or more.
  • the maximum forming height is an index used for evaluation of the stretch formability.
  • the maximum forming height is taken as a punch stroke at the occurrence of fracture where the load rapidly decreases in the load-stroke diagram.
  • a lubricating polyethylene sheet is interposed between the punch and the steel sheet, and stretch forming is performed at a blank holding force of 1,000 kgf, and then the maximum forming height is determined by measuring the height at the time of fracture (punch stroke).
  • the high-strength steel sheet according to the embodiment of the present invention has the maximum forming height of 20 mm or more, and preferably 21 mm or more.
  • a hole expansion ratio ⁇ is determined in accordance with Japan Iron and Steel Federation Standard JFS T1001.
  • JFS T1001 Japan Iron and Steel Federation Standard JFS T1001.
  • ⁇ % d ⁇ d 0 / d 0 ⁇ 100
  • the high-strength sheet according to the embodiment of the present invention has a hole expansion ratio ⁇ of 20% or more, and preferably 30% or more. This makes it possible to obtain excellent workability such as press formability.
  • a tensile test was performed at a deformation rate of 10 mm/min and the test piece was fractured. Then, the fracture surface was observed and the value (t 1 /t 0 ) obtained by dividing a thickness t 1 in a thickness direction of the fracture surface by an original thickness t 0 was taken as a thickness reduction ratio.
  • the thickness reduction rate in this test is 50% or more, preferably 52% or more, and more preferably 55% or more. This makes it possible to obtain a steel sheet having excellent impact resistance properties since the steel sheet is hardly fractured even if it deforms greatly upon collision.
  • Cross tensile strength of spot welding was evaluated in accordance with JIS Z 3137. Two 1.4 mm-thick steel sheets laid one upon another were used as the conditions of spot welding. Using a dome radius type electrode, spot welding was performed under a welding pressure of 4 kN by increasing a current by 0.5 kA in a range from 6 kA to 12 kA, and the current value (minimum current value) at which dust was generated during welding was examined. A cross joint spot-welded at a current that is 0.5 kA lower than the minimum current value was used as a test piece for measurement of a cross tensile strength. Samples having a cross tensile strength of 6 kN or higher were rated "Good". The cross tensile strength is preferably 8 kN or higher, and more preferably 10 kN or higher.
  • cross tensile strength is 6 kN or higher, it is possible to obtain parts having high bonding strength during welding when automobile parts and the like are manufactured from the steel sheet.
  • the inventors of the present application have found that the above-mentioned desired steel structure is attained by subjecting a rolled material with predetermined composition to a heat treatment (multi-step austempering treatment) mentioned later, thus obtaining a high-strength steel sheet having the above-mentioned desired properties.
  • Fig. 1 is a diagram explaining a method for manufacturing a high-strength sheet according to the embodiment of the present invention, especially a heat treatment.
  • the rolled material to be subjected to the heat treatment is usually produced by cold-rolling after subjecting to hot-rolling.
  • the process is not limited thereto, and the rolled material may be produced by any one of hot-rolling and cold-rolling.
  • the conditions of hot-rolling and cold-rolling are not particularly limited.
  • a rolled material is held in a two-phase coexistence region intermediate between an Ac 1 point and an Ac 3 point, more specifically at a temperature T 1 (Ac 1 ⁇ T 1 ⁇ 0.2 ⁇ Ac 1 point + 0.8 ⁇ Ac 3 ) between the Ac 1 point and 0.2 ⁇ Ac 1 point + 0.8 ⁇ Ac 3 point for 5 seconds or more and, as shown in [3] and [4] of Fig. 1 , the rolled material is held at a heating temperature T 2 to a temperature T 2 (Ac 3 ⁇ T 2 ) of the Ac 3 point or higher for 5 to 600 second, thus austenitizing the rolled material.
  • the rolled material is heated to the temperature T 1 , followed by holding for 5 seconds or more.
  • the holding time is preferably 900 seconds or less.
  • the holding temperature T 1 may be held at a constant temperature between the Ac 1 point and 0.2 ⁇ Ac 1 point + 0.8 ⁇ Ac 3 point, as shown in [2] of Fig. 1 , and the holding temperature may be varied between the Ac 1 point and 0.2 ⁇ Ac 1 point + 0.8 ⁇ Ac 3 point, for example, slow heating is performed at a temperature between the Ac 1 point and 0.2 ⁇ Ac 1 point + 0.8 ⁇ Ac 3 point.
  • the amount of Mn-concentrated austenite decreases to reduce the variation in the concentration of Mn in retained austenite (carbon-concentrated region), thus failing to obtain sufficient stretch formability.
  • the Mn concentration of austenite decreases to reduce the variation in the concentration of Mn in retained austenite (carbon-concentrated region), thus failing to obtain sufficient stretch formability.
  • the longer the holding time at the temperature T1, the better, and the holding time is preferably 900 seconds or less from the viewpoint of the productivity.
  • the temperature T 1 is preferably between 0.9 ⁇ Ac 1 point + 0.1 ⁇ Ac 3 point and 0.3 ⁇ Ac 1 point + 0.7 ⁇ Ac 3 point, and the holding time at the temperature T 1 is preferably 10 seconds or more and 800 seconds or less.
  • the temperature T 1 is more preferably between 0.8 ⁇ Ac 1 point + 0.2 x Ac 3 point and 0.4 ⁇ Ac 1 point + 0.6 ⁇ Ac 3 point, and the holding time at the temperature T 1 is more preferably 30 seconds or more and 600 seconds or less.
  • the heating rate to the temperature T 1 shown as [1] in Fig. 1 is preferably 5 to 20°C/sec.
  • a material is heated to a temperature T 2 (Ac 3 ⁇ T 2 ) of an Ac 3 point or higher, followed by holding at the temperature T 2 , thus austenitizing the rolled material.
  • the holding time at the temperature T 2 is 5 to 600 seconds.
  • the portion which was ferrite turns to austenite when heated to the temperature T 1 .
  • Mn is not concentrated. Therefore, the Mn-nonconcentrated region exists in austenite, together with the Mn-concentrated region, and it becomes possible to increase the variation in the concentration of Mn in retained austenite (carbon-concentrated region) in the high-strength sheet after the heat treatment, thus enabling realization of the high stretch formability.
  • the ferrite fraction of the obtained high-strength sheet exceeds 5%, leading to a decrease in YR.
  • the temperature T 2 is preferably the Ac 3 point + 50°C or lower.
  • the holding time at the temperature T 2 is more than 600 seconds, the Mn concentration in the Mn-concentrated region decreases to reduce the variation in the Mn concentration in retained austenite, thus degrading the stretch formability.
  • the temperature T 2 is preferably the Ac 3 point + 10°C or higher, and the holding time at the temperature T 2 is preferably 10 to 450 seconds.
  • the temperature T 2 is more preferably the Ac 3 point + 20°C or higher, and the holding time at the temperature T 2 is more preferably 20 to 300 seconds.
  • Heating from the temperature T 1 to the temperature T 2 shown in [3] of Fig. 1 is preferably performed at a heating rate of 0.1°C/sec or more and less than 10°C/sec.
  • the Ac 1 point and the Ac 3 point may be determined by the measurement, or may be calculated by a generally known calculation formula using the composition.
  • cooling is performed to a cooling stopping temperature T 3 between 100°C or higher and lower than 300°C at an average cooling rate of 10°C/sec or more.
  • the cooling stopping temperature is controlled to a temperature in a range of 100°C or higher and lower than 300°C, the final amount of retained austenite is controlled by adjusting the amount of austenite remaining without being transformed into martensite.
  • Cooling is performed at an average cooling rate of 10°C/sec or more between at least 650°C and 300°C. This is because, by setting the average cooling rate at 10°C/sec or more, the formation of ferrite during cooling is suppressed to form a structure composed mainly of fine martensite.
  • Preferred example of such cooling includes cooling (slow cooling) to a rapid cooling starting temperature T 4 of 650°C or higher at relatively low average cooling rate of 0.1°C/sec or more and 10°C/sec or less, as shown in [5] of Fig. 1 , followed by cooling (rapid cooling) from the rapid cooling starting temperature T 4 to a cooling stopping temperature T 3 of 300°C or lower at an average cooling rate of 10°C/sec or more and less than 200°C/sec, as shown in [6] of Fig. 1 .
  • rapid cooling starting temperature T 4 By setting the rapid cooling starting temperature T 4 at 650°C or higher, it is possible to suppress the formation of ferrite during cooling (slow cooling).
  • cooling rate is less than 10°C/sec, ferrite is formed to decrease YR. MA becomes coarse, thus decreasing the hole spreading ratio.
  • the cooling rate is preferably 15°C/°C or higher, and more preferably 20°C/sec or more.
  • the cooling stopping temperature T 3 is preferably 120°C or higher and 280°C or lower, and more preferably 140°C or higher and 260°C or lower.
  • holding may be performed at the cooling stopping temperature T 3 .
  • the holding time is preferably 1 to 150 seconds. Even if the holding time is more than 150 seconds, the productivity of the steel sheet is degraded though properties of the obtained steel sheet are not significantly improved. Therefore, the holding time is preferably set at 150 seconds or less.
  • heating is performed from the above cooling stopping temperature to a reheating temperature in a range of 300°C to 500°C at a reheating rate of 30°C/sec or more. Rapid heating enables a decrease in retention time in a temperature range where precipitation and growth of carbide are promoted, thus making it possible to suppress the formation of fine carbide.
  • the reheating rate is preferably 60°C/sec or more, and more preferably 70°C/sec.
  • Such rapid heating can be achieved by a method such as high-frequency heating or electric heating.
  • a tempering parameter P represented by the following equation (1) is set at 10,000 or more and 14,500 or less and the holding time is set at 1 to 150 seconds.
  • the tempering parameter P When the tempering parameter P is less than 10,000, carbon diffusion from martensite to austenite does not sufficiently occur and austenite becomes unstable, thus failing to ensure the amount of retained austenite, leading to insufficient TS ⁇ EL balance. If the tempering parameter P is more than 14,500, the formation of carbide cannot be prevented even by a short-time treatment, thus failing to ensure the amount of retained austenite, leading to degradation of TS ⁇ EL balance. Even if the tempering parameter is appropriate, carbide is formed in martensitic laths if the heating rate is too low and heating time is too long, so that crack propagation easily occurs during collision deformation, thus degrading collision resistance properties. The amount of carbide in martensite laths can be determined from the scattering intensity of X-ray small angle scattering.
  • the reheating temperature T 5 is lower than 300°C, diffusion of carbon becomes insufficient, thus failing to obtain sufficient amount of retained austenite, leading to degradation of TS ⁇ EL. If the reheating temperature T 5 is higher than 500°C, retained austenite is decomposed into cementite and ferrite, thus failing to ensure properties because of insufficient retained austenite.
  • the holding time is less than 1 second, carbon diffusion may be insufficient, similarly. Therefore, it is preferred to hold at a reheating temperature T 5 for 1 second or more. If the holding time is more than 150 seconds, carbon may precipitate as cementite, similarly. Therefore, the holding time is preferably 150 seconds or less.
  • the reheating temperature T 5 is preferably 320 to 480°C, and more preferably 340 to 460°C.
  • the tempering parameter P is preferably 10,500 to 14,500, and the holding time at this time is preferably 1 to 150 seconds.
  • the tempering parameter P is more preferably 11,000 to 14,000, and the holding time at this time is preferably 1 to 100 seconds, and more preferably 1 to 60 seconds.
  • cooling may be performed to the temperature of 200°C or lower, for example, room temperature.
  • the average cooling rate to 200°C or lower is preferably 10°C/sec.
  • the high-strength sheet according to the embodiment of the present invention can be obtained by the above-mentioned heat treatment.
  • a steel sheet having a thickness of 2.5 mm was produced by multistage rolling after heating to 1,200°C. At this time, the end temperature of hot-rolling was set at 880°C. After cooling to 600°C at 30°C/sec, cooling was stopped and the steel sheet was inserted into a furnace heated to 600°C, held for 30 minutes and then furnace-cooled to obtain a hot-rolled steel sheet.
  • the hot-rolled steel sheet was subjected to pickling to remove the scale on the surface, and then cold-rolled to reduce the thickness to 1.4 mm.
  • This cold rolled sheet was subjected to a heat treatment to obtain samples.
  • the heat treatment conditions are shown in Table 2.
  • the number in parentheses, for example, [2] in Table 2 corresponds to the process of the same number in parentheses in Fig. 1 .
  • sample No. 1 is sample in which austenitizing was performed only at a temperature of the Ac 3 point or higher corresponding to a temperature T 2 without performing in two stages of austenitizing at a temperature T 1 and austenitizing at the temperature T 2 .
  • Sample No. 9 is sample in which cooling was performed to a reheating temperature, followed by holding the same temperature instead of cooling to a cooling stopping temperature between 100°C or higher and lower than 300°C (samples in which the steps corresponding to [7] to [8] in Fig. 1 were skipped).
  • Samples 15 and 31 to 36 are samples in which the heating temperature T 2 and the rapid cooling starting temperature T 4 were set at the same temperature. After the austenitizing, cooling was performed to a cooling stopping temperature T 3 in a single stage.
  • Reheating corresponding to [8] was performed by an electric heating method.
  • the ferrite fraction, the total fraction of tempered martensite and tempered bainite (mentioned as "tempered M/B” in Table 3), the amount of retained (amount of retained ⁇ ), the half-width of the concentration distribution in the carbon-concentrated region, and the scattering intensity at the q value of 1 nm -1 in X-ray small angle scattering were determined by the above-mentioned methods.
  • a two-dimensional microfocused X-ray diffraction apparatus (RINT-RAPID II) manufactured by Rigaku Corporation was used. The results are shown in Table 3.
  • the steel structure (balance structure) other than the steel structure shown in Table 3 is untempered martensite in samples excluding sample No. 9, or untempered bainite in sample No. 9.
  • No. Steel No. Steel structure Ferrite Tempered M/B Amount of retained ⁇ Average size of MA Half-width of concentration distribution of Mn in carbon-concentrated region Scattering intensity at q value of 1 nm -1 % % % ⁇ m % by mass cm -1 1 a 0 72 17.3 0.67 *0.23 *2.42 2 a 0 72 17.7 0.90 *0.17 0.74 3 a 0 71 17.1 0.67 *0.22 *2.34 4 a *30.7 *40 16.2 0.80 *0.24 0.72 5 a *30.5 *36 16.8 0.87 *0.24 0.74 6 a *13.7 61 18.0 0.70 0.71 0.68 7 a 0 *53 17.5 *2.07 0.86 0.71 8 a 0 72 17.1 0.60 *0.16 0.71 9 a 0
  • Samples Nos. 13, 15, 18, 21 and 28 to 36 are Examples that satisfy all requirements (composition, manufacturing conditions and steel structure) defined in the embodiment of the present invention. All of these samples achieve the tensile strength (TS) of 980 MPa or higher, the yield ratio (YR) of 0.75 or more, TS ⁇ EL of 20,000 MPa% or higher, the hole expansion ratio ( ⁇ ) of 20% or more, the maximum forming height of 16 mm or more, the SW cross tension of 6 kN or higher, and the R5 tensile thickness reduction ratio (RA) of 50% or more.
  • TS tensile strength
  • YiR yield ratio
  • TS ⁇ EL 20,000 MPa% or higher
  • ⁇ hole expansion ratio
  • RA tensile thickness reduction ratio
  • sample No. 1 since holding was performed only at a temperature of an Ac 3 point or higher corresponding to a temperature T 2 without performing austenitizing in two stages of austenitizing at a temperature T 1 and austenitizing at the temperature T 2 , the sample exhibits small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height. Since the holding time [7] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.
  • Sample No. 2 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height because of low holding temperature T 1 .
  • Sample No. 3 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height because of high holding temperature T 1 . Since the holding time [7] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.
  • samples Nos. 4 and 5 after heating to the heating temperature T 1 and holding, the temperature that is the same as T 1 was selected as the heating temperature T 2 , thus failing to austenitize at sufficiently high temperature. Therefore, samples exhibited a large amount of ferrite, low total fraction of tempered martensite and tempered bainite, and small half-width of the concentration distribution of Mn in the carbon-concentrated region. As a result, samples exhibited low tensile strength, low yield ratio and low maximum forming height.
  • Sample No. 6 exhibited large amount of ferrite because of low heating temperature T 2 , leading to low yield ratio.
  • Sample No. 7 exhibited low total fraction of tempered martensite and tempered bainite and large average size of MA because of high cooling stopping temperature T 3 , thus decreasing the hole expansion ratio.
  • Sample No. 8 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region because of short holding time at the heating temperature T 1 , thus decreasing the maximum forming height.
  • Sample No. 10 exhibited small amount of retained austenite because of low cooling stopping temperature T 3 , thus decreasing low value of TS ⁇ EL and maximum forming height.
  • Sample No. 16 exhibited large amount of ferrite and low total fraction of tempered martensite and tempered bainite because of low rapid cooling starting temperature T 4 , thus decreasing the tensile strength and the yield ratio. Since the holding time [9] was as long as 300 seconds, carbide (cementite) was precipitated. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.
  • Sample No. 19 exhibited the parameter increased to 14,604 because of high reheating temperature T 5 , thus decreasing the amount of retained austenite. As a result, the value of TS ⁇ EL and maximum forming height decreased. Because of large scattering intensity of X-ray small angle scattering, it can be said that cementite having a size of about 1 nm has a large volume fraction. As a result, collision resistance properties (thickness reduction ratio) were degraded.
  • Sample No. 20 exhibited the parameter decreased to 9,280 because of low reheating temperature T 5 , thus decreasing the amount of retained austenite. As a result, the value of TS ⁇ EL and maximum forming height decreased.
  • Sample No. 22 exhibited small C amount, insufficient amount of retained austenite, thus decreasing the value of TS ⁇ EL and maximum forming height.
  • Sample No. 23 exhibited large Mn amount and insufficient amount of retained austenite, thus decreasing the value of TS ⁇ EL and maximum forming height.
  • Sample No. 24 exhibited small Mn amount, large amount of ferrite, and insufficient total amount of tempered martensite and tempered bainite. As a result, the tensile strength and the yield ratio decreased.
  • Sample No. 25 exhibited small amount of Si + Al, insufficient total amount of tempered martensite and tempered bainite, small amount of retained austenite, and large average size of MA. As a result, the value of TS ⁇ EL, the hole expansion ratio and the maximum forming height decreased.
  • Sample No. 26 exhibited large C amount, thus decreasing the SW cross tensile strength.
  • Sample No. 27 exhibited large amount of Si + Al, thus decreasing the value of TS ⁇ EL and the maximum forming height.
  • Sample No. 28 exhibited small half-width of the concentration distribution of Mn in the carbon-concentrated region and low maximum forming height because of high holding temperature T 1 .
  • the manufacturing method according to the embodiment of the present invention enables the production of the steel sheet that satisfies the composition and the steel structure defined in the embodiment of the present invention.

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WO2020127557A1 (de) * 2018-12-19 2020-06-25 Voestalpine Stahl Gmbh Verfahren zur herstellung von thermo-mechanisch hergestellten warmbanderzeugnissen
WO2020127558A1 (de) * 2018-12-19 2020-06-25 Voestalpine Stahl Gmbh Verfahren zur herstellung von konventionell warmgewalzten warmbanderzeugnissen
WO2020127555A1 (de) * 2018-12-19 2020-06-25 Voestalpine Stahl Gmbh Verfahren zur herstellung von thermo-mechanisch hergestellten profilierten warmbanderzeugnissen
EP4180545A4 (de) * 2020-08-27 2023-12-06 Nippon Steel Corporation Warmgewalztes stahlblech
EP4180546A4 (de) * 2020-08-27 2023-12-06 Nippon Steel Corporation Warmgewalztes stahlblech
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WO2020127561A1 (de) * 2018-12-19 2020-06-25 Voestalpine Stahl Gmbh Verfahren zur herstellung von konventionell warmgewalzten, profilierten warmbanderzeugnissen
WO2020127557A1 (de) * 2018-12-19 2020-06-25 Voestalpine Stahl Gmbh Verfahren zur herstellung von thermo-mechanisch hergestellten warmbanderzeugnissen
WO2020127558A1 (de) * 2018-12-19 2020-06-25 Voestalpine Stahl Gmbh Verfahren zur herstellung von konventionell warmgewalzten warmbanderzeugnissen
WO2020127555A1 (de) * 2018-12-19 2020-06-25 Voestalpine Stahl Gmbh Verfahren zur herstellung von thermo-mechanisch hergestellten profilierten warmbanderzeugnissen
EP4180545A4 (de) * 2020-08-27 2023-12-06 Nippon Steel Corporation Warmgewalztes stahlblech
EP4180546A4 (de) * 2020-08-27 2023-12-06 Nippon Steel Corporation Warmgewalztes stahlblech
EP4206343A4 (de) * 2020-08-27 2023-12-13 Nippon Steel Corporation Warmgewalztes stahlblech
EP4206344A4 (de) * 2020-08-27 2023-12-13 Nippon Steel Corporation Warmgewalztes stahlblech

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