EP2314729A1 - Feuilles d'acier biphase à haute résistance ayant d'excellentes propriétés de déformation dynamique et son procédé de préparation - Google Patents

Feuilles d'acier biphase à haute résistance ayant d'excellentes propriétés de déformation dynamique et son procédé de préparation Download PDF

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EP2314729A1
EP2314729A1 EP10181225A EP10181225A EP2314729A1 EP 2314729 A1 EP2314729 A1 EP 2314729A1 EP 10181225 A EP10181225 A EP 10181225A EP 10181225 A EP10181225 A EP 10181225A EP 2314729 A1 EP2314729 A1 EP 2314729A1
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Prior art keywords
deformation
strength
strain
steel sheet
phase
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German (de)
English (en)
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EP2314729B2 (fr
EP2314729B1 (fr
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Akihiro Uenishi
Manabu Takahashi
Yukihisa Kuriyama
Yasuharu Sakuma
Osamu Kawano
Junichi Wakita
Hidesato Mabuchi
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Nippon Steel Corp
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Nippon Steel Corp
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Priority claimed from JP19029797A external-priority patent/JP3530347B2/ja
Priority claimed from JP22300897A external-priority patent/JP3936440B2/ja
Priority claimed from JP25893897A external-priority patent/JP3839928B2/ja
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    • 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
    • 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
    • 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
    • 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
    • C22C38/04—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
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
    • C23C2/0224—Two or more thermal pretreatments
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/024—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by cleaning or etching
    • 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

Definitions

  • the present invention relates to dual-phase type high-strength steel sheets, for automobiles use, which have excellent dynamic deformation properties and exhibit excellent impact absorption properties, and are intended to be used as structural members and reinforcing materials primarily for automobiles, as well as to a method of producing them.
  • high-strength steels have been increasing for the purpose of achieving lighter weight vehicle bodies in consideration of fuel consumption restrictions on automobiles and even more applications for high-strength steel are expected as domestic and foreign restrictions, relating to estimated impact absorption properties in automobile accidents, become rapidly more broad and strict.
  • front side members can allow impact energy to be absorbed through collapse of the member, thus lessening the impact experienced by passengers.
  • the present inventors have reported on the high strain rate properties and impact energy absorption properties of high-strength thin steel sheets in CAMP-ISIJ Vol.9 (1966), pp.1112-1115 , wherein they explain that the dynamic strength at a high strain rate of 10 3 (s -1 ) increases dramatically compared to the static strength at a low strain rate speed of 10 -3 (s -1 ), that absorption energy during crashes is increased by greater steel material strengths, that the strain rate dependency of materials depends on the structure of the steel, and that TRIP type steel (Transformation induced plasticity type steel) and dual-phase (hereunder, "DP”) type steel exhibit both excellent press formability and high impact absorption properties. Also, the present inventors have already filed Japanese Patent Applications No. 8-98000 and No.
  • the present invention has been proposed as a means of overcoming the problems described above, and provides dual-phase type high-strength steel sheets for automobiles use, which have excellent impact absorption properties and excellent dynamic deformation properties, as well as a method of producing them.
  • the invention further provides dual-phase type high-strength steel sheets, for automobiles, with excellent dynamic deformation properties, which are high-strength, steel sheets used for automotive parts, such as front side members, and which are selected based on exact properties and standards for impact energy absorption during collisions and can reliably provide guaranteed safety, as well as a method of producing them.
  • the invention still further provides dual-phase type high-strength steel sheets for automobiles with excellent dynamic deformation properties, which exhibit all the properties suitable for press forming of members, including excellent shape fixability, excellent stretchability and excellent flangeability, as well as a method of producing them.
  • Impact absorbing members such as front side members of automobiles are produced by bending and press forming of steel sheets. Because impacts during automobile collisions are absorbed by such members which have undergone press forming, they must have high impact absorption properties even after having undergone the pre-deformation corresponding to the press forming. At the current time, however, no attempt has been made to obtain high-strength steel sheets with excellent impact absorption properties as actual members, with consideration of both the increase in the deformation stress by press forming and the increase in deformation stress due to a higher strain rate, as was mentioned above.
  • steel sheets with a dual-phase (DP) structure are ideal as high-strength steel sheets with excellent impact absorption properties for actual members which are press formed as described above. It was demonstrated that such steel sheets with a dual-phase microstructure, which is a composite microstructure wherein the dominating phase is a ferrite phase responsible for the increase in deformation resistance by an increased strain rate, and the second phase includes a hard martensite phase, have excellent dynamic deformation properties.
  • the microstructure of the final steel sheets is a composite structure wherein the dominating phase is ferrite and another low temperature product phase includes a hard martensite phase at a volume fraction of 3 ⁇ 50% after deformation at 5% equivalent strain of the steel sheet.
  • the volume fraction of the martensite phase since high-strength steel sheets and even steel sheets with high dynamic deformation properties cannot be obtained if the martensite phase is less than 3%, the volume fraction of the martensite phase must be at least 3%. Also, if the martensite phase exceeds 50%, this results in a smaller volume fraction of the ferrite phase responsible for greater deformation resistance due to increased deformation speed, making it impossible to obtain steel sheets with excellent dynamic deformation properties compared to static deformation strength while also hindering press formability, and therefore it was found that the volume fraction of the martensite phase must be 3 ⁇ 50%.
  • the present inventors then pursued experimentation and research based on these findings and, as a result, found that although the degree of pre-deformation corresponding to press forming of impact absorbing members such as front side members sometimes reaches a maximum of over 20%, depending on the location, the majority are locations with 0% ⁇ 10% of equivalent strain, and that by understanding the effect of pre-deformation in this range, it is possible to estimate the behavior of the member as a whole after pre-deformation. Consequently, according to the invention, a deformation of from 0% to 10% of equivalent strain was selected as the amount of pre-deformation applied to members during press forming.
  • Fig. 1 is a graph showing the relationship between the absorption energy (Eab) of a press formed member during collision and the material strength (S), for the different steel types shown in Table 5, according to an example to be described later.
  • the material strength S is the tensile strength (TS) according to the common tensile test.
  • the member absorption energy (Eab) is the absorption energy in the lengthwise direction (direction of the arrow) along a press formed member such as shown in Fig. 2 , upon collision with a 400 kg mass weight at a speed of 15 m/sec, to a crushing degree of 100 mm.
  • hat-shaped section 1 consists of a 2.0 mm-thick steel sheet formed into a hat-shaped section 1 with a steel sheet 2 of the same thickness and the same type of steel, joined together by spot welding, the hat-shaped section 1 having a corner radius of 2 mm, and with spot welding points indicated by 3.
  • Fig. 1 From Fig. 1 it is seen that the member absorption energy (Eab) tends to increase with the strength of materials under normal tensile testing, though with considerable variation.
  • the materials in Fig. 1 were subjected to pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and then the static deformation strength ⁇ s when deformed in a strain rate range of 5 x 10 -4 - 5 x 10 -3 (s -1 ) and the dynamic deformation strength ⁇ d when deformed in a strain rate range of 5 x 10 2 - 5 x 10 3 (s -1 ) after the pre-deformation, were measured.
  • a classification was possible based on ( ⁇ d - ⁇ s).
  • the symbols plotted in Fig. 1 were as follows:
  • Fig. 3 shows the relationship between the work hardening coefficient of a steel sheet and the dynamic energy absorption which indicates the member impact absorption properties, for a class of materials with the same yield strength.
  • increased work hardening coefficients of the steel sheets result in improved member impact absorption properties (dynamic energy absorption)
  • the work hardening coefficient of a steel sheet can properly indicate the member impact absorption properties so long as the yield strength class is the same.
  • the yield strength x work hardening coefficient can be an indicator of the member impact absorption properties.
  • work hardening coefficient was expressed in terms of an n value of 5% ⁇ 10% strain in consideration of the strain undergone by members during press forming, from the viewpoint of improving the dynamic energy absorption, work hardening coefficients of under 5% strain or work hardening coefficients of even more than 10% strain may be preferred.
  • a schematic illustration of this test method is shown in Fig. 7 .
  • Fig. 5 , 4 is a worktop
  • 5 is a test piece
  • 6 is a spot welding section.
  • Fig. 6 7 is a hat-shaped test piece and 8 is a spot welding section.
  • Fig. 7 9 is a worktop
  • 10 is a test piece
  • 11 is a falling weight (150 kg)
  • 12 is a frame
  • 13 is a shock absorber.
  • the work hardening coefficient and yield strength of each steel sheet was determined in the following manner.
  • the steel sheet was shaped into a JIS-#5 test piece (gauge length: 50 mm, parallel width: 25 mm), subjected to tensile test at a strain rate of 0.001 (s -1 ) to determine the yield strength and work hardening coefficient (n value at 5% ⁇ 10% strain).
  • the steel sheet used had a sheet thickness of 1.2 mm and the steel sheet composition contained C at 0.02-0.25 wt%, either or both Mn and Cr at a total of 0.15-3.5 wt% and one or more of Si, Al and P at a total of 0.02-4.0 wt%, with the remainder Fe as the main component.
  • Fig. 8 is a graph showing the relationship between the average value ⁇ dyn of the deformation stress in the range of 3 ⁇ 10% of equivalent strain when deformed in a strain rate range of 5 x 10 2 - 5 x 10 3 (s -1 ) and the static material strength (TS), as an index of the impact energy absorption property upon collision according to the invention, where the static material strength (TS) is the tensile strength (TS: MPa) in the static tensile test as measured in a strain rate range of 5 x 10 -4 - 5 x 10 -3 (s -1 ).
  • impact absorbing members such as front side members have a hat-shaped cross-sectional shape
  • the present inventors have found that despite deformation proceeding up to a high maximum strain of over 40%, at least 70% of the total absorption energy is absorbed in a strain range of 10% or lower in a high-speed stress-strain diagram. Therefore, the dynamic deformation resistance with high-speed deformation at 10% or lower was used as the index of the high-speed collision energy absorption property.
  • the index used for the impact energy absorption property was the average stress: ⁇ dyn in the range of 3 ⁇ 10% of equivalent strain when deformed in a strain rate range of 5 x 10 2 -5 x 10 3 (s -1 ) high-speed tensile deformation.
  • the average stress: ⁇ dyn of 3 ⁇ 10% upon high-speed deformation generally increases with increasing static tensile strength ⁇ maximum stress (TS: MPa) in a static tensile test measured in a stress rate range of 5 x 10 -4 - 5 x 10 -3 (s -1 ) ⁇ of the steel material prior to pre-deformation or baking treatment. Consequently, increasing the static tensile strength (which is synonymous with the static material strength) of the steel material directly contributes to an improved impact energy absorption property of the member. However, increased strength of the steel results in poorer press formability into members, making it difficult to obtain members with the necessary shapes. Consequently, steels having a high ⁇ dyn with the same tensile strength TS are preferred.
  • steel sheets wherein the average value ⁇ dyn (MPa) of the deformation stress in the range of 3 ⁇ 10% of equivalent strain when deformed in a strain rate range of 5 x 10 2 - 5 x 10 3 (s -1 ), after pre-deformation of more than 0% and less than or equal to 10% of equivalent strain satisfies the inequality: aden ⁇ 0.766 x TS + 250 as expressed in terms of the tensile strength (TS: MPa) in the static tensile test as measured in a strain rate range of 5 x 10 -4 - 5 x 10 -3 (s -1 ) prior to pre-deformation, have higher impact energy absorption properties as actual members compared to other steels, and that the impact energy absorption property is improved without increasing the overall weight of the member, making it possible to provide high-strength steel sheets with high dynamic deformation resistance.
  • YS(0)/TS'(5) is no greater than 0.7, which amount is dependent on the initial microstructure, the amount of solid solution elements in the low temperature product phase other than the martensite phase and the main ferrite phase, and the deposited state of carbides, nitrides and carbonitrides.
  • YS(0) is the yield strength
  • TS'(5) is the tensile strength (TS') in the static tensile test with pre-deformation at 5% of equivalent strain or after further bake hardening treatment (BH treatment). It was also demonstrated that steel sheets with even more excellent dynamic deformation properties can be obtained when the yield strength: YS(0) x work hardening coefficient is at least 70.
  • the martensite is at a volume fraction of 3 ⁇ 50%, and preferably 3 ⁇ 30%.
  • the average grain size of the martensite is preferably no greater than 5 ⁇ m, and the average grain size of the ferrite is preferably no greater than 10 ⁇ m. That is, the martensite is hard, and contributes to a decrease in the yield ratio and an improvement in the work hardening coefficient, by producing a mobile dislocations primarily in adjacent ferrite grains; however, by satisfying the restrictions mentioned above it is possible to disperse fine martensite in the steel, so that the improvement in the properties spreads throughout the entire steel sheet.
  • this dispersion of fine martensite in the steel can help to avoid deterioration in the hole expansion ratio and tensile strength x total elongation, which is an adverse effect of the hard martensite. Also, because it is possible to reliably achieve work hardening coefficient ⁇ 0.130, tensile strength x total elongation ⁇ 18,000 and hole expansion ratio ⁇ 1.2, it is thereby possible to improve the impact absorption properties and press formability.
  • the yield ratio With a martensite volume fraction of less than 3%, the yield ratio becomes larger while the press formed member cannot exhibit an excellent work hardening property (work hardening coefficient ⁇ 0.130) after it has undergone collision deformation, and since the deformation resistance (load) stays at a low level, and the dynamic energy absorption is low preventing improvement in the impact absorption properties.
  • the volume fraction of the martensite is preferred to be no greater than 30%.
  • the ferrite is present at a volume fraction of preferably at least 50%, and more preferably at least 70%, and its average grain size (mean circle equivalent diameter) is preferably no greater than 10 ⁇ m, and more preferably no greater than 5 ⁇ m, with the martensite preferably adjacent to the ferrite. This aids the fine dispersion of the martensite in the ferrite matrix, while effectively extending the property-improving effect, beyond simply a local effect, to the entire steel sheet, favorably acting to prevent the adverse effects of the martensite.
  • the structure of the remainder present with the martensite and ferrite may be a mixed structure comprising a combination of one or more from among pearlite, bainite, retained ⁇ , etc., and although primarily bainite is preferred in cases which require hole expansion properties, since retained ⁇ undergoes work-induced transformation into martensite by press forming, experimental results have shown that including retained austenite prior to press forming has an effect even in preferred small amounts (5% or less).
  • the ratio of the martensite and ferrite particle sizes is no greater than 0.6, and the ratio of the hardnesses to be at least 1.5.
  • Dual-phase type high-strength steel sheets with excellent dynamic deformation properties which are used according to the invention are steel sheets containing the following chemical compositions, in terms of weight percentage: C at 0.02-0.25%, either or both Mn and Cr at a total of 0.15-3.5%, one or more from among Si, Al and P at a total of 0.02-4.0%, if necessary also one or more from among Ni, Cu and Mo at a total of no more than 3.5%, one or more from among Nb, Ti and V at no more than 0.30%, and either or both Ca and REM at 0.0005-0.01% for Ca and 0.005-0.05% for REM, with the remainder Fe as the primary component.
  • C is the element which most strongly affects the microstructure of the steel sheet, and if its content is too low it will become difficult to obtain martensite with the desired amount and strength. Addition in too great an amount leads to unwanted carbide precipitation, inhibited increase in deformation resistance at higher strain rates and overly high strength, as well as poor press formability and weldability; the content is therefore 0.02-0.25 wt%.
  • Mn, Cr: Mn and Cr have an effect of stabilizing austenite and guaranteeing sufficient martensite, and are also solid solution hardening elements; they must therefore be added in a minimum amount of 0.15 wt%, but if added in too much the aforementioned effect becomes saturated thus producing adverse effects such as preventing ferrite transformation, and thus they are added in the maximum amount of 3.5 wt%.
  • Si, Al, P are useful elements for producing martensite, and they promote production of ferrite and suppress precipitation of carbides, thus having the effect of guaranteeing sufficient martensite, as well as a solid solution hardening effect and a deoxidization effect.
  • P can also promote martensite formation and solid solution hardening, similar to Al and Si. From this standpoint, the minimum amount of Si + Al + P added must be at least 0.02 wt%. On the other hand, excessive addition will saturate this effect and result instead in brittleness, and therefore the maximum amount of addition is no more than 4.0 wt%.
  • Si scales can be avoided by adding Si at no greater than 0.1 wt%, and conversely by adding it at 1.0 wt% or greater Si scales can be produced over the entire surface so that they are not conspicuous.
  • the P content may be kept at no greater than 0.05%, and preferably no greater than 0.02%.
  • Ni, Cu, Mo are added when necessary, and are austenite-stabilizing elements similar to Mn, which increase the hardenability of the steel, and are effective for adjustment of the strength. From the standpoint of weldability and chemical treatment, they can be used when the amounts of C, Si, Al and Mn are restricted, but if the total amount of these elements added exceeds 3.5 wt% the dominant ferrite phase will tend to be hardened, thus inhibiting the increase in deformation resistance by a greater strain rate, as well as raising the cost of the steel sheet; the amount of these elements added is therefore 3.50 wt% or lower.
  • Nb, Ti, V These elements are added when necessary, and are effective for strengthening the steel sheet through formation of carbides, nitrides and carbonitrides. However, when added at greater than 0.3 wt% they are deposited in large amounts in the dominant ferrite phase or at the grain boundaries as carbides, nitrides and carbonitrides, becoming a source of the mobile dislocation during high speed deformation, and inhibiting the increase in deformation resistance by greater strain rates. In addition, the deformation resistance of the dominant phase becomes higher than necessary, thus wasting the C and leading to higher costs; the maximum amount to be added is therefore 0.3 wt%.
  • B is an element which is effective for strengthening since it improves the hardenability of the steel by suppressing production of ferrite, but if it is added at greater than 0.01 wt% its effect will be saturated, and therefore B is added at a maximum of 0.01 wt%.
  • Ca is added to at least 0.0005 wt% for improved press formability (especially hole expansion ratio) by shape control (spheroidization) of sulfide-based inclusions, and the maximum amount thereof to be added is 0.01 wt% in consideration of effect saturation and the adverse effect due to increase in the aforementioned inclusions (reduced hole expansion ratio).
  • REM is added in an amount of from 0.005% to 0.05 wt%.
  • the amount of S is no greater than 0.01 wt%, and preferably no greater than 0.003 wt%, from the standpoint of press formability (especially hole expansion ratio) by sulfide-based inclusions, and reduced spot weldability.
  • the pre-deformation may be press forming for member shaping, or it may be working with a tempering rolling or tension leveler which applied to the steel sheet material prior to its press forming.
  • a tempering roller and tension leveler may be used. That is, the means used may include a tempering rolling, a tension leveler, or a tempering roller and tension leveler.
  • the steel sheet material may also be subjected to press forming after being worked with a tempering rolling or tension leveler.
  • the amount of pre-deformation applied with the tempering rolling and/or tension leveler i.e.
  • the degree of plastic deformation (T) will differ depending on the initial dislocation density, and T should be small if the initial density is large. Also, with few solid solution elements the introduced dislocations cannot be fixed, and high dynamic deformation properties cannot be guaranteed. Consequently, it was found that the plastic deformation (T) is determined based on the ratio between the yield strength YS(0) and the tensile strength TS'(5) in the static tensile test with pre-deformation at 5% of equivalent strain or after further bake hardening treatment (BH treatment), or YS(0)/TS'(5).
  • YS(0)/TS'(5) is an indicator of the sum of the initial dislocation density and the dislocation density introduced by 5% deformation, and the amount of the solid solution elements; it may be concluded that a smaller YS(0)/TS'(5) means a higher initial dislocation density and more of the solid solution elements.
  • YS(0)/TS'(5) is therefore no greater than 0.7, and is preferably provided according to the following equation: 2.5 YS 0 / TS ′ 5 ⁇ 0.5 + 15 ⁇ T ⁇ 2.5 YS 0 / TS ′ 5 ⁇ 0.5 + 0.5 wherein the upper limit for T is determined from the standpoint of press formability including impact absorption property and flexibility.
  • a continuous cast slab is fed directly from casting to a hot rolling step, or is hot rolled upon reheating after momentary cooling.
  • Thin gauge continuous casting and continuous hot rolling techniques may be applied for the hot rolling in addition to normal continuous casting, but in order to avoid a lower ferrite volume fraction and a coarser average grain size of the thin steel sheet microstructure, the bar (cast strip) thickness at the hot rolling approach side (the initial steel bar thickness) is preferred to be at least 25 mm.
  • the mean circle equivalent size of ferrite of the steel sheet is made coarser, while it is also a disadvantage against obtaining the desired martensite.
  • the final pass rolling speed for the hot rolling is preferred to be at least 500 mpm and more preferably at least 600 mpm, in light of the problems described above.
  • the mean circle equivalent diameter of ferrite of the steel sheet is made coarser, while it is also a disadvantage against obtaining the desired martensite.
  • the finishing temperature for the hot rolling is from Ar 3 - 50°C to Ar 3 + 120°C.
  • Ar 3 - 50°C deformed ferrite is produced, with inferior work hardening property and press formability.
  • Ar 3 + 120°C and the mean circle equivalent size of ferrite of the steel sheet is made coarser, while it is also becomes difficult to obtain the desired martensite.
  • the average cooling rate for cooling in the run-out table is at least 5°C/sec. At less than 5°C/sec it becomes difficult to obtain the desired martensite.
  • the coiling temperature is no higher than 350°C. At higher than 350°C it becomes difficult to obtain the desired martensite.
  • the hot rolling is carried out so that when the finishing temperature for hot rolling is in the range of Ar 3 - 50°C to Ar 3 + 120°C, the metallurgy parameter A satisfies inequalities (1) and (2).
  • the average cooling rate on the run-out table is at least 5°C/sec, and the coiling to be carried out under conditions such that the relationship between the metallurgy parameter A and the coiling temperature (CT) satisfies inequality (3).
  • the cold rolled sheet according to the invention is then subjected to the different steps following hot-rolling and coiling and is cold rolled and subjected to annealing.
  • the annealing is ideally continuous annealing through an annealing cycle such as shown in Fig. 12 , and during the annealing of the continuous annealing step, it must be kept for at least 10 seconds in the temperature range of Ac 1 - Ac 3 . At less than Ac 1 austenite will not be produced and it will therefore be impossible to obtain martensite thereafter, while at greater than Ac 3 the austenite monophase structure will be coarse, and it will therefore be impossible to obtain the desired average grain size for the martensite.
  • the maximum residence time is preferably no greater than 200 seconds, from the standpoint of avoiding addition to the equipment and coarsening of the microstructure.
  • the cooling after this annealing must be at an average cooling rate of at least 5°C/sec. At less than 5°C/sec the desired space factor for the martensite cannot be achieved. Although there is no particular upper limit here, it is preferably 300°C/sec when considering temperature control during the cooling.
  • the cooled steel sheet is heated to a temperature To from Ac 1 - Ac 3 in the continuous annealing cycle shown in Fig. 12 , and cooled under cooling conditions provided by a method wherein cooling to a secondary cooling start temperature Tq in the range of 550°C-To at the primary cooling rate of 1 ⁇ 10°C/sec is followed by cooling to a secondary cooling end temperature Te which is no higher than a temperature Tem which is determined by the chemical compositions of the steel and annealing temperature To, at a secondary cooling rate of 10 ⁇ 200°C/sec.
  • This is a method whereby the cooling end temperature Te in the continuous annealing cycle shown in Fig.
  • Te 12 is represented as a function of the chemical compositions and annealing temperature, and is kept under a given critical value. After cooling to Te, the temperature is preferably held in a range of Te - 50°C to 400°C for up to 20 minutes prior to cooling to room temperature.
  • T1 is the temperature calculated from the solid solution element concentration excluding C
  • T2 is the temperature calculated from the C concentration in the retained austenite at Ac 1 and Ac 3 determined by the chemical compositions of the steel and Tq determined by the annealing temperature To.
  • Ceq* represents the carbon equivalents in the retained austenite at the annealing temperature To.
  • Te when Te is equal to or greater than Tem, the desired martensite cannot be obtained. Also, if Toa is 400°C or higher, the martensite obtained by cooling is tempered, making it impossible to achieve satisfactory dynamic properties and press formability. On the other hand, if Toa is less than Te - 50°C, additional cooling equipment is necessary, and greater variation will result in the material due to the difference between the temperature of the continuous annealing furnace and the temperature of the steel sheet; this temperature was therefore determined as the lower limit. Also, the upper limit for the holding time was determined to be 20 minutes, because when it is longer than 20 minutes it becomes necessary to expand the equipment.
  • the microstructure of the steel sheet is a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite at a volume fraction from 3% ⁇ 50% after shaping and working at 5% equivalent strain, and wherein the difference between the quasi-static deformation strength ⁇ s when deformed in a strain rate range of 5 x 10 -4 - 5 x 10 -3 (1/s) after pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and the dynamic deformation strength ⁇ d measured in a strain rate range of 5 x 10 2 - 5 x 10 3 (1/s) after the aforementioned pre-deformation, i.e.
  • the steel sheets according to the invention may be made into any desired product by annealing, tempering rolling, electronic coating or hot-dip coating.
  • the 26 steel materials listed in Table 1 were heated to 1050 ⁇ 1250°C and subjected to hot rolling, cooling and coiling under the production conditions listed in Table 2, to produce hot rolled steel sheets.
  • the steel sheets satisfying the chemical composition conditions and production conditions according to the invention have a dual-phase structure with a martensite volume fraction of at least 3% and no greater than 50%, and as shown in Fig.
  • the mechanical properties of the hot rolled steel sheets indicated excellent impact absorption properties as represented by a work hardening coefficient of at least 0.13 at 5 ⁇ 10% strain, ⁇ d - ⁇ s ⁇ 60 MPa, and ⁇ dyn ⁇ 0.766 x TS + 250, while also having suitable press formability and weldability.
  • Table 2 Production conditions Steel No. Hot rolling conditions Cooling conditions Coiling conditions Finishing temp. °C Initial steel strip thickness (mm) Final pass rolling speed (mpm) Final sheet thickness (mm) Strain rate (/sec) log A calculated ⁇ T °c Inequality (2) Aver. cooling rate (°Clsec) Note Coiling temp.
  • the 22 steel materials listed in Table 5 were heated to 1050 ⁇ 1250°C and subjected to hot rolling, cooling and coiling, followed by acid pickling and then cold rolling under the conditions listed in Table 6 to produce cold rolled steel sheets.
  • Temperatures Ac 1 and Ac 3 were then calculated from the chemical compositions for each steel, and the sheets were subjected to heating, cooling and holding under the annealing conditions listed in Table 6, prior to cooling to room temperature.
  • the steel sheets satisfying the chemical composition conditions and production conditions according to the invention have a dual-phase structure with a martensite volume fraction of at least 3% and no greater than 50% and, as shown in Fig.
  • the mechanical properties of the hot-rolled steel sheets indicated excellent impact absorption properties as represented by a work hardening coefficient of at least 0.13 at 5 ⁇ 10% strain, ⁇ d - ⁇ s ⁇ 60 MPa, and ⁇ dyn ⁇ 0.766 x TS + 250, while also having suitable press formability and weldability.
  • the microstructure was evaluated by the following method.
  • a tensile test was conducted according to JIS5 (gauge mark distance: 50 mm, parallel part width: 25 mm) with a strain rate of 0.001/s and, upon determining the tensile strength (TS), yield strength (YS), total elongation (T. El) and work hardening coefficient (n value for 1% ⁇ 5% strain), the YS x work hardening coefficient and TS x T. E1. were calculated.
  • the stretch flanging property was measured by expanding a 20 mm punched hole from the burrless side with a 30° cone punch, and determining the hole expansion ratio (d/d 0 ) between the hole diameter (d) at the moment at which the crack penetrated the plate thickness and the original hollow diameter (d 0 , 20 mm).
  • the spot weldability was judged to be unsuitable if a spot welding test piece bonded at a current of 0.9 times the expulsion current using an electrode with a tip radius of 5 times the square root of the steel sheet thickness underwent peel fracture when ruptured with a chisel.
  • the present invention makes it possible to provide, in an economical and stable manner, high-strength hot rolled steel sheets and cold rolled steel sheets for automobiles which provide previously unobtainable excellent impact absorption properties and press formability and thus offers a markedly wider range of objects and conditions for uses of high-strength steel sheets.

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EP10181225.3A 1997-03-17 1998-03-16 Feuilles d'acier biphase à haute résistance ayant d'excellentes propriétés de déformation dynamique Expired - Lifetime EP2314729B2 (fr)

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JP8243497 1997-03-17
JP19029997 1997-07-15
JP19029797A JP3530347B2 (ja) 1997-07-15 1997-07-15 動的変形特性に優れた高強度鋼板の選定方法
JP22300897A JP3936440B2 (ja) 1997-08-06 1997-08-06 耐衝突安全性と成形性に優れた自動車用高強度鋼板とその製造方法
JP25893897A JP3839928B2 (ja) 1997-07-15 1997-09-24 動的変形特性に優れたデュアルフェーズ型高強度鋼板
EP98907247.5A EP0969112B2 (fr) 1997-03-17 1998-03-16 Procede de preparation des toles d'acier biphasees a haute resistance mécanique et a haute capacité d'absorption d'energie de chock

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EP98907247.5A Division-Into EP0969112B2 (fr) 1997-03-17 1998-03-16 Procede de preparation des toles d'acier biphasees a haute resistance mécanique et a haute capacité d'absorption d'energie de chock
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CN104281774A (zh) * 2014-09-02 2015-01-14 上海交通大学 Q&p钢在不同应变率单拉后残余奥氏体含量的预测方法
CN104281774B (zh) * 2014-09-02 2017-06-13 上海交通大学 Q&p钢在不同应变率单拉后残余奥氏体含量的预测方法
RU2578618C1 (ru) * 2014-11-18 2016-03-27 Публичное акционерное общество "Северсталь" (ПАО "Северсталь") Способ производства полос из низколегированной свариваемой стали
CN105483530A (zh) * 2015-11-30 2016-04-13 丹阳市宸兴环保设备有限公司 一种石油天然气输送管用热轧宽钢板材料

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CA2283924A1 (fr) 1998-09-24
WO1998041664A1 (fr) 1998-09-24
CA2283924C (fr) 2006-11-28
KR100334949B1 (ko) 2002-05-04
EP0969112A4 (fr) 2003-05-21
EP0969112A1 (fr) 2000-01-05
KR20000076372A (ko) 2000-12-26
EP2314729B2 (fr) 2017-03-08
CN1251140A (zh) 2000-04-19
AU717294B2 (en) 2000-03-23
CN1080321C (zh) 2002-03-06
AU6311898A (en) 1998-10-12
TW426742B (en) 2001-03-21
EP2314729B1 (fr) 2012-02-08
EP0969112B2 (fr) 2017-03-08
EP0969112B1 (fr) 2011-08-17

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