EP3327152B1 - Procédé de formage à chaud d'une ébauche d'acier - Google Patents

Procédé de formage à chaud d'une ébauche d'acier Download PDF

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EP3327152B1
EP3327152B1 EP16201185.2A EP16201185A EP3327152B1 EP 3327152 B1 EP3327152 B1 EP 3327152B1 EP 16201185 A EP16201185 A EP 16201185A EP 3327152 B1 EP3327152 B1 EP 3327152B1
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temperature
steel
range
blank
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EP3327152A1 (fr
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Thomas James Taylor
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Tata Steel UK Ltd
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Tata Steel UK Ltd
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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
    • 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
    • C21D1/62—Quenching devices
    • C21D1/673—Quenching devices for die quenching
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D—WORKING OR PROCESSING OF SHEET METAL OR METAL TUBES, RODS OR PROFILES WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D22/00—Shaping without cutting, by stamping, spinning, or deep-drawing
    • B21D22/02—Stamping using rigid devices or tools
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D—WORKING OR PROCESSING OF SHEET METAL OR METAL TUBES, RODS OR PROFILES WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D22/00—Shaping without cutting, by stamping, spinning, or deep-drawing
    • B21D22/02—Stamping using rigid devices or tools
    • B21D22/022—Stamping using rigid devices or tools by heating the blank or stamping associated with heat treatment
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D—WORKING OR PROCESSING OF SHEET METAL OR METAL TUBES, RODS OR PROFILES WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D37/00—Tools as parts of machines covered by this subclass
    • B21D37/16—Heating or cooling
    • 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
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • 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
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/30—Ferrous alloys, e.g. steel alloys containing chromium with cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
    • C22C38/18—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
    • 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
    • C21D2211/001—Austenite
    • 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

Definitions

  • the invention relates to a method for hot forming a steel blank into an article having enhanced mechanical properties, such as an automotive part with improved ductility and impact toughness, and to a hot formed article obtained by said method.
  • hot forming also known as hot stamping, press-hardening and die-quenching
  • the basics of the hot forming technique and steel compositions adapted to be used were for the first time described in GB1490535 .
  • the blank In a typical hot forming process using a boron steel composition the blank is furnace-heated and austenised at 900-950 °C, transferred from furnace to forming tool, and stamped into the desired part geometry.
  • the blank has an ultimate tensile strength R m ⁇ 200 MPa and a total elongation A>50%.
  • the formed blank is finally die-quenched to 100-200 °C with a mean cooling rate of >30 °C/s and a homogenous martensitic microstructure is obtained, while it is constrained in the forming tool.
  • the final martensitic part typically exhibits a proof strength Rp 0.2 >1100 MPa, an ultimate tensile strength R m >1500 MPa and a total elongation A ⁇ 8%.
  • Hot stamped boron steel parts typically include anti-intrusive structural body parts such as roof pillar-, door beam- and bumper beam-reinforcements that constitute the 'safety-cell'.
  • the relatively soft and ductile high-temperature austenitic microstructure during forming permits down-gauging and lightweighting while not compromising forming limits and moreover, permits part-consolidation and in turn, increased structural strength and increased process efficiency compared to cold forming as joining/welding is reduced.
  • the martensitic transformation which releases forming stresses and geometric constraint on the part during quenching, eliminates springback giving rise to geometric accuracy.
  • the ultra high strength martensitic microstructure resulting in the final part permits down-gauging (lightweighting) simultaneously while improving anti-intrusive crashworthiness compared to the use of cold formed parts.
  • the supremacy of boron steel in hot stamping applications owes to the (quench) hardenability (owing to the boron addition) and in turn ultra high strength martensitic microstructure that can be obtained from the relatively lean chemical composition.
  • hot stamped martensitic boron steel provides excellent mechanical properties for anti-intrusive crashworthiness, the limited ductility and in turn limited toughness give rise to very poor impact-energy absorptive crashworthiness.
  • hot stamped boron steel and the advantages of the hot stamping process cannot be exploited with impact-energy absorptive structural body parts, such as fender and longitudinal beams that constitute 'crumple zones'.
  • Such parts are typically cold formed from ferritic or multiphase steels, such as Transformation Induced Plasticity (TRIP) steels.
  • TRIP Transformation Induced Plasticity
  • EP 2719787 A1 , EP2719786 A1 and US2013/192726 A1 relate to a hot press-formed product that can be provided with a prescribed shape and at the same time heat treated to have the prescribed strength when a preheated steel sheet is formed into the prescribed shape, and a process for producing such a hot press-formed product.
  • the present invention aims to find the right steel composition and an improved method for hot-forming steel blanks into a complex shaped article having ultra high tensile strength combined with high ductility and in turn, excellent impact-energy absorptive crashworthiness and which does not have the drawbacks mentioned above for conventional hot formed boron steels and cold formed multiphase steels.
  • a method of hot-forming a steel blank into an article according to the present invention is defined in claim 1
  • a steel article obtainable by the method according to the invention is defined in claim 14 and preferred embodiments are defined in claims 2-13.
  • a method for hot-forming a steel blank into an article according to the present invention comprises the steps of claim 1.
  • the inventors have found that through forming the heated blank into an article in the sequence as described above, complex shaped articles with enhanced properties can be obtained.
  • the articles exhibit excellent impact-energy absorptive crashworthiness and in turn, down-gauging and lightweighting opportunities based on impact-energy absorptive crashworthiness compared to the use of conventional hot-formed boron steels and cold-formed multiphase steels.
  • the steel blank In the common hot stamping process, the steel blank is simply formed into the desired geometry at and then die-quenched to near-ambient temperature in which the 'complete' transformation to martensite takes place. Thus, the final part exhibits a fully or almost fully martensitic microstructure.
  • removing the tool pressure means that the punch of the hot-forming tool is removed/open, but the hot-formed article still remains in the die.
  • the part remains in the forming tool, but without tool pressure applied and so the cooling rate from the temperature range between Ar1 and Ms down to 200-400 °C is relatively slow. During the slow air cooling down to 200-400 °C, bainitic ferrite and retained austenite are produced in the microstructure.
  • the tool pressure is resumed to continue die-quenching to 100-200 °C and then the part is removed from the tool.
  • resuming the tool pressure means that after the opening of the punch and slow air cooling, the punch of the hot-forming tool is closed and pressure is applied again to the hot-formed article in the hot-forming tool. During the process, rather than complete martensite formation, we achieve a significant volume fraction of retained austenite to provide the TRIP steel like microstructure in the final part.
  • step (c) the heated blank is formed in the forming tool into the desired part geometry having a partially or a fully homogenous austenitic microstructure and simultaneously quenched from a temperature T2 to a temperature T3 in a time period t3 under the application of tool pressure.
  • Temperature T3 is in the chemistry specific Ar1 and Ms temperature range of the steel composition. The inventors have found that when T3 is in the range of 400-600 °C good results have been achieved. When T3 is below 400 °C, atomic diffusion rates are lower and thus the carbon partitioning rate from bainitic ferrite to austenite is reduced. When T3 is above 400 °C, atomic diffusion rates are higher and permit an appreciable degree of carbon partitioning from bainitic ferrite to austenite in order to stabilise the austenite at ambient temperature (and thus retain austenite in the final microstructure). When T3 is above 600 °C, pearlite formation may occur rather than bainitic ferrite and retained austenite. Better results have been achieved when T3 is in the range of 400-580 °C. Advantageously T3 is in the range of 450 - 550 °C.
  • Time t3 is dictated by the quenching rate achieved between temperature T2 and temperature T3.
  • the inventors have found that when t3 is equal or less than 10 s and a quenching rate between T2 and T3 of 50 °C/s or more is applied good results have been achieved.
  • the quenching rate is equal or more than 100, more preferably equal or more than 150, most preferably equal or more than 200 °C/s.
  • Below 50 °C/s austenite may transform to pearlite and proeutectoid ferrite on cooling, preventing martensite, bainite and retained austenite formation in the final microstructure.
  • t3 is equal or less than 8, more preferably equal or less than 5, most preferably equal or less than 4s.
  • Advantageously t3 is equal or less than 2 s.
  • Quenching the article in the forming tool to temperature T3 takes place by the action of conductive heat transfer from the heated blank constrained in the forming tool under tool pressure, to the forming tool itself.
  • the forming tool may or may not exhibit integrated cooling channels to enhance the conductive heat transfer rate.
  • step (d) slow cooling of the article from temperature T3 to temperature T4 takes place with the formed part residing in the forming tool in a time period t4, but without applying forming tool pressure i.e. the punch is open/removed. While and during the punch is removed/open the residual heat collected by the forming tool during the forming and quenching stage from temperature T2 to T3 and moreover, latent heat of phase transformation, is able to circulate the formed part by convective heat transfer allowing slow cooling of the hot-formed article.
  • Slow cooling from T3 to T4 takes place purely through air-cooling, with the cooling rate minimised by the residual heat in the forming tool and the latent heat of phase transformation circulating the formed part by convective heat transfer.
  • Temperature T4 may be above the chemistry specific Ms temperature, between the chemistry specific Ms and Mf temperatures or below the chemistry specific Mf temperature of the steel.
  • the three preferred temperature ranges mentioned above dictate the volume fraction of martensite that forms in the final quench.
  • T4 is in the range of 200-400 °C good results have also been achieved. Slow cooling to below 200 °C may result in austenite decomposing into bainite. The rapid quench from temperatures higher than 200 °C ensures that austenite is retained in the final microstructure. There is no specific upper limit on T4. In practice, it will be challenging to maintain T4 above 400 °C as natural air cooling will drop the temperature below 400 °C before we have formed an appreciable volume fraction of bainitic ferrite.
  • T4 is 300 - 380 °C, more preferably 320 - 370 °C.
  • Time t4 is dictated by the residual heat accumulation and management in the forming tool.
  • the anisothermal bainitic transformation takes place with slow cooling from temperature T3 to temperature T4, marked by austenite transforming to bainitic ferrite, carbon partitioning from the bainitic ferrite to remaining austenite, solute carbon concentration increase of the remaining austenite and stabilisation of the remaining austenite at ambient temperature.
  • the inventors have further found that when t4 is equal or more than 60 s, good results have been achieved. When t4 is less than 60 s it is likely that insufficient retained austenite will be formed. If t4 is longer than 180 s it would be also not beneficial, because forced heating to retard cooling must be applied.
  • t4 is in the range of 60-180 s, preferably is in the range of 100-180 s, more preferably is in the range of 120-180 s.
  • the inventors have further found that when the cooling rate between T3 and T4 is in the range of 2-10 °C/s good results have been achieved.
  • the cooling rate is in the range of 2-8 °C/s.
  • the cooling rate is in the range of 2-5 °C/s.
  • the formed blank is quenched from temperature T4 to temperature T5, wherein austenite may be completely retained during the final quench to temperature T5, or a significant volume fraction may transform to martensite.
  • austenite volume fraction There is no specific lower or upper limit on the retained austenite volume fraction. However, it is preferably to retain not less than 5%, more preferably not less than 10 vol. % austenite in the final part following the final quench. If no retained austenite is present there will be no stress/strain inducted transformation during a crash event. The inventors have found that when 5 % or more retained austenite are present a significant stress/strain induced transformation during a crash event can be achieved. Preferably the retained austenite in the hot-formed part is of at most 20 vol.%. Quenching takes place by the action of conductive heat transfer from said blank constrained in the forming tool under tool pressure, to the forming tool itself. The forming tool may or may not exhibit integrated cooling channels to enhance the conductive heat transfer rate.
  • T5 is in the range of 20 to 200 °C good results have been obtained.
  • austenite may decompose into bainite rather than retained at ambient temperature.
  • any martensite formed may temper above 200 °C, degrading tensile strength.
  • T5 is in the range of 20-100 °C.
  • Time t5 is dictated by the quenching rate achieved between T4 and T5.
  • the inventors have found that when t5 is equal or less than 8 s good results have been obtained. Better results have been achieved when t5 is equal or less than 4s.
  • Advantageously t5 is equal or less than 2 s.
  • the inventors have found that when the cooling rate is at least 50 °C/s good results have been achieved. Better results have been achieved when the cooling rate is at least 100 °C/s.
  • the cooling rate is at least 150 °C/s.
  • the steel blank comprises the elements as indicated in claim 1. Preferred ranges for these elements are given in claim 13. The reason for the amounts of the main constituting elements (in wt%) is provided below.
  • the method according to the invention comprises - prior to the hot forming step (c), the steps of:
  • a steel strip or sheet is provided as an intermediate for the subsequent steps.
  • the steel strip or sheet can be obtained by standard casting processes.
  • the steel strip or sheet is cut to a steel blank and then heated to a temperature T1 for a time period t1.
  • a preformed steel blank may also be used.
  • the preformed blank may be partially or entirely formed into the desired geometry, preferably at ambient temperature.
  • the heating apparatus may be an electric or gas powered furnace, electrical resistance heating device, infra-red induction heating device or any other heating device.
  • Temperature T1 is above the chemistry specific Ac1 temperature of the steel to form a ferritic austenitic microstructure.
  • T1 may be above the chemistry specific Ac3 temperature to produce a fully or almost fully homogenous austenitic microstructure with uniform distribution of carbon.
  • the microstructure is a homogenous austenitic microstructure the formability is enhanced.
  • T1 is at least 710 °C good results have been achieved.
  • T1 is in the range of 850 - 1150, more preferably 980 - 1100, most preferably 1000 -1050 °C.
  • the holding time t1 may be chosen in combination with the temperature T1 and the blank thickness in order to control austenitic grain growth and associated quench hardenability of said steel. This means that a greater blank thickness will require longer time t1 and/or higher temperature T1 in order to achieve the desired (a given) austenitic grain size compared to lower blank thickness.
  • the austenite grain size at the end of time t1 will influence quench hardenability, where larger austenite grain size (corresponding to smaller austenite grain boundary surface area) reduces the number of nucleation sites for proeutectoid ferrite formation on cooling and thereby, increases quench hardenability.
  • t1 is equal or less than 15 min. In a preferred embodiment the inventors have found that when time t1 is equal or less than 10 min. good results have been achieved.
  • Advantageously t1 is in the range of 3-5 minutes in order to increase time and energy efficiency of the process.
  • the heating of the steel blank to temperature T1 in step (a) may be performed with the steel blank residing in the forming tool.
  • This technique of 'in press heating' is applied with electrodes that contact the steel blank in the forming tool providing electrical resistance heating in the forming tool.
  • This adaption permits greater quench hardenability and avoidance of proeutectoid ferrite formation for a leaner steel composition.
  • a leaner chemical composition may be used without forming proeutectoid ferrite.
  • the heated blank is transferred from the heating apparatus to a forming tool in a time period t2 (step (b)).
  • the blank may cool from temperature T1 to temperature T2 by the act of natural air-cooling and/or any other available cooling method.
  • the heated blank may be transferred from the heating apparatus to the forming tool by an automated robotic system or any other transfer method.
  • Temperature T2 is above the chemistry specific Ar1 temperature of the steel to exhibit an austenitic-ferritic microstructure.
  • T2 may be above the chemistry specific Ar3 temperature to exhibit a homogenous austenitic microstructure.
  • a fully or almost fully austenitic microstructure at the commencement of forming and quenching is preferable for optimal formability.
  • a delayed time t2 or a temperature T2 lower than Ar3 permits the formation of proeutectoid ferrite, which is desirable in view of the mechanical properties and product performance of the final article, giving rise to a lower proof strength, higher proof strength to ultimate tensile strength ratio and improved impact-energy absorptive crashworthiness.
  • the inventors have further found that when T2 is above the Ar3 temperature good results have also been achieved.
  • the article has excellent formability associated with the fully or almost fully austenitic microstructure so that very complex shapes can be formed in a single stroke.
  • Time t2 may be also chosen in combination with T1, t1 and T2 in order to control the microstructural evolution of steel at the commencement of forming and quenching.
  • T1 and/or longer time t1 will give rise to greater austenitic grain growth and delay proeutectoid ferrite formation during time t2.
  • a shorter time t2 may be used for a given degree of proeutectoid ferrite formation during time t2.
  • a lower temperature T2 is required (to allow proeutectoid ferrite formation), necessarily a longer time t2 will be required to reach temperature T2.
  • t2 is equal or less than 10 s good results have been also achieved.
  • t2 is equal or less than 8s, more preferably equal or less than 6s.
  • the carbon content is at least 0.05 wt% to provide adequate interstitial solid solution strengthening, adequate quench hardenability and adequate stabilisation of austenite at ambient temperature while maintaining sufficiently low carbon equivalent for automotive resistance spot-welding techniques.
  • the range of preferable carbon contents will provide a range of products exhibiting a range of strength-ductility values according to the invention.
  • the carbon content is in the range of 0.15-0.45.
  • the carbon is in the range of 0.15-0.35.
  • carbon is in the range of 0.15-0.25.
  • Mn 0.05-3.00
  • the manganese content is at least 0.05 to provide adequate substitutional solid solution strengthening, adequate quench hardenability and adequate stabilisation of austenite at ambient temperature, while minimising segregation of Mn during casting and while maintaining sufficiently low carbon equivalent for automotive resistance spot-welding techniques.
  • the inventors have found that the range of manganese contents in combination with the range of carbon contents provides a product exhibiting a range of strength-ductility values.
  • the Mn content is equal to or less than 2.50.
  • manganese is in the range of 1.00-1.50.
  • the silicon, aluminium and phosphorus contents prevent carbide precipitation in the carbon enriched austenite, thus enabling the Ms temperature of the remaining austenite to be depressed below ambient and retention of the austenite at ambient temperature.
  • the silicon, aluminium and phosphorus contents are chosen so to provide the optimal balance of carbide precipitation retardation, kinetics of the anisothermal bainitic transformation, weldability, coatability, economy and manufacture-ability with conventional steelmaking and sheet / strip processing infrastructure.
  • Si ⁇ 2.0 The silicon content is chosen as the dominant carbide precipitation retardant, but is limited to a maximum of in order to minimise formation of surface-bound silicates that impede hot-rolling, cold-rolling and coating of strip steels.
  • silicon is equal or less than 1.0. More preferably silicon is in the range of 0.25-0.75. In a most preferred embodiment silicon is in the range of 0.25-0.50.
  • Al ⁇ 2.0 Aluminium is limited to a maximum of 2.0 in order to preserve weldability and minimise 'nozzle blockage' during steelmaking and casting.
  • the aluminium content is preferably at least 0.25 to increase the kinetics of the anisothermal bainitic transformation so that appreciable volume fractions of austenite can be retained at ambient temperature after limited time t4.
  • Preferably Al is limited to a maximum of 1.0, more preferably aluminium is in the range of 0.25-0.75 wt%.
  • Advantageously aluminium is in the range of 0.25 - 0.50.
  • P ⁇ 0.17 The phosphorus content is limited to a maximum of 0.17 and is present to limit silicon and aluminium contents, while still providing satisfactory carbide precipitation retardation.
  • phosphorus is in the range of 0.02-0.07.
  • Advantageously phosphorus is in the range of 0.02-0.05 in order to provide acceptable weldability by automotive standards.
  • Nb ⁇ 0.1 The niobium content is limited to a maximum of 0.1 to form niobium carbide precipitates and in turn provide austenitic grain size refinement which increases the kinetics of the anisothermal bainitic transformation. Niobium carbides may also provide precipitation strengthening. Preferably niobium is in the range of 0.02 - 0.04. Advantageously niobium is in the range of 0.02 - 0.03.
  • Ti ⁇ 0.1 The titanium content is limited to a maximum of 0.1 in order to bond all nitrogen content and form titanium-nitride precipitates at a titanium: nitrogen ratio of 3.42:1 and in turn to enable the boron content to remain un-bonded in solid solution (in order for hardenability) and moreover, titanium is present to bond all residual sulphur content and form titanium-sulphide precipitates at a ratio of titanium:sulphur 1.5:1.
  • the titanium 'excess' remaining after all nitrogen and sulphur has been bonded will bond with carbon to provide titanium-carbide precipitates and in turn provide precipitation strengthening.
  • titanium is in the range of 0.02 - 0.04.
  • Advantageously titanium is in the range of 0.02-0.03.
  • Cr ⁇ 1.0 The chromium content is limited to at most 1.0 in order to aide hot workability and corrosion resistance.
  • chromium is in the range of 0.2-0.4.
  • chromium is in the range of 0.2-0.3.
  • B ⁇ 0.01 The boron content is limited to a maximum of 0.01 in order to enhance quench hardenability and in turn to avoid proeutectoid ferrite formation during cooling from temperature T1 to temperature T2 and during quenching to temperature T3.
  • boron is in the range of 0.001-0.005.
  • boron is in the range of 0.003-0.005.
  • N ⁇ 0.01: The nitrogen content is limited to a maximum of 0.01 to prevent boron nitride precipitation and minimise the unwanted TiN formation, where TiN has no benefit to mechanical properties. N has no significant effect on the mechanical properties.
  • nitrogen is in the range of 0.001-0.005, more preferably in the range of 0.001-0.003.
  • Sulphur is an impurity and needs to be minimised for minimisation of harmful non-metallic inclusions. Therefore the sulphur content is limited to a maximum of 0.002, preferably to a maximum of 0.001.
  • the steel strip composition may optionally comprise one or more alloying elements in small amounts such as V, Mo, Co, W, or rare earth elements.
  • V, Mo, Co, W, and the other microalloy rare earth elements can be added in a total amount not exceeding 1.00 wt%. in order to benefit from the known property-improving effects of these elements.
  • the steel microstructural composition may be that of a multiphase steel, preferably a TRIP or complex phase (CP) like steel that exhibits the capacity for stress / strain induced transformation effect. More particularly the steel is an advanced high strength TBF (TRIP-aided bainitic ferrite) steel product suitable for hot press forming as a boron steel alternative for automotive applications.
  • TBF TRIP-aided bainitic ferrite
  • the inventors have surprisingly found that applying the present hot-forming method to a TRIP steel like composition, results in hot-formed articles combining the capacity for the stress/strain induced transformation effect with the principles of TRIP steel microstructural evolution (i.e. carbide-free bainitic ferrite formation and austenite retention).
  • the obtained hot-formed TRIP steel articles exhibit significant retained austenite volume fractions in the range of 5 to 20 vol. % and the capacity for stress/strain induced transformation.
  • the optimal austenite volume fraction and capacity for stress/strain induced transformation is established after the forming process and therefore is present in the article, without the austenite volume fraction and capacity for stress/strain induced transformation being used up in the forming process. This enables down-gauging and lightweighting simultaneously while improving impact-energy absorptive crashworthiness.
  • the articles of the present invention show vastly improved mechanical properties, which are comparable to those of the traditional TRIP steel prior to cold forming, and impact-energy absorptive crashworthiness can be achieved.
  • the steel strip, sheet, blank, preformed blank, or article is provided with a coating as claimed in claim 11.
  • a coating may be performed prior to the hot-forming process or after the hot-forming process and has the purpose of minimising oxidation of the steel when exposed to an oxidising atmosphere at high temperature and/or providing cathodic corrosion protection of the final hot formed part.
  • the zinc based coating is a galvanized or galvannealed coating.
  • the coating can be applied in various ways, hot dip galvanising is preferred using a standard GI coating bath.
  • Other Zn coatings may also be applied.
  • An example comprises a Zn alloy coating according to WO 2008102009 , in particular a zinc alloy coating layer as claimed in claim 12.
  • An additional element typically added in a small amount of less than 0.2 wt%, could be selected from the group comprising Pb or Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr or Bi.
  • Pb, Sn, Bi and Sb are usually added to form spangles.
  • the total amount of additional elements in the zinc alloy is at most 0.2%.
  • each additional element is present in an amount ⁇ 0.02 wt%, preferably each is present in an amount ⁇ 0.01 wt%. Additional elements are usually only added to prevent dross forming in the bath with molten zinc alloy for the hot dip galvanising, or to form spangles in the coating layer.
  • the articles according to the present invention exhibit good adhesion to a coating layer, have good surface appearance and superior corrosion resistance after coating.
  • It is a further object of the present invention is to provide a multiphase microstructured steel article having simultaneously improved strength and ductility.
  • the multiphase microstructure according to the invention is obtained by proving a steel composition as described above and by applying careful control of the thermomechanical treatment method as described above.
  • the steel article obtainable by the method according to the first aspect of the invention is claimed in claim 14.
  • the inventors have found that the hot-formed article obtained by the present invention having a total elongation in the range of 25-35% and a R m xA product of 20000-30000 MPa% exhibits improved impact-energy absorptive crashworthiness in the final part and in turn, down-gauging and lightweighting potential based on crashworthiness.
  • R m may be in the range of 750-1500, more preferably in the range of 900-1300, most preferably in the range of 900-1150 MPa.
  • the total elongation may be in the range of 29-35%.
  • the R m xA product may be in the range of 24000-30000 MPa%, more preferably in the range of 26000-28000 MPa%.
  • the final microstructure of the steel comprises bainitic ferrite, (retained) austenite, optionally martensite and/or proeutectoid ferrite.
  • the microstructure may further comprise proeutectoid ferrite.
  • the optimal austenite volume fraction and capacity for stress/strain induced transformation is established after the forming process and therefore is present in the article.
  • the austenite volume fraction is optimal for stress/strain induced transformation to martensite during a crash event, giving rise to optimal impact-energy absorptive crashworthiness.
  • the desired microstructural evolution will not be achieved if the steps are arranged in any other sequence than that disclosed herein.
  • the final microstructure comprises (sum added up to 100) by volume fraction (vol.%) of:
  • Proeutectoid ferrite 0-75. Proeutectoid ferrite may be present. However, in those cases where an elevated yield strength (yield strength to ultimate tensile strength ratio) is aimed for, the fraction of proeutectoid ferrite must be limited. Surprisingly, with the correct balance of other microstructural constituents (most notably retained austenite) high work hardening and tensile ductility can still be attained. Above this limit, the final microstructure will not contain enough bainitic ferrite and/or martensite, and thus ultimate tensile strength will be too low.
  • yield strength yield strength to ultimate tensile strength ratio
  • proeutectoid ferrite Large fractions of proeutectoid ferrite will also lead to low initial yield strength (low yield strength to ultimate tensile strength ratio).
  • the proeutectoid ferrite is present in an amount of at most 30 vol.%.
  • the proeutectoid ferrite is 0 vol.%. This provides optimal volume fractions of martensite (for ultimate tensile strength) and retained austenite (for stress/strain induced transformation effect) and thus, the maximum R m xA product.
  • the absence of proeutectoid ferrite and in turn, a microstructure characterised by microconstituents exhibiting more similar mechanical properties is envisaged to be beneficial for edge ductility / hole-expansion capacity (HEC).
  • HEC edge ductility / hole-expansion capacity
  • Bainitic ferrite 10 - 30.
  • Bainitic ferrite not only provides strength, but the formation thereof is also a prerequisite for retaining austenite. Beyond the upper limit, insufficient martensite will be present, and thus ultimate tensile strength will be too low and moreover, it is unlikely that more than 30 vol.% bainitic ferrite can be produced within the prescribed heat treatment.
  • the formation of bainitic ferrite in the presence of silicon, aluminium and/or phosphorus drives the partition of carbon to the austenite phase, enabling levels of carbon enrichment in the austenite phase allowing formation of a (meta)stable phase at ambient temperature.
  • Bainitic ferrite has also the advantage over martensite as a strengthening phase that it causes less micro-scale localisation of strain to softer phases and consequently improves resistance to fracture with respect to dual phase steels.
  • the bainitic ferrite is in the range of 20-30 in order to achieve a more sufficient austenite stabilization.
  • Martensite 0-75. Martensite may be formed during the final rapid quench of the hot forming process. An optimal balance of ultimate tensile strength and ductility is obtained, when the martensite volume fraction is equal or less than 75 vol.%. The complete absence of martensite is advantageous for maximum ductility, where the microstructure is composed of entirely proeutectoid ferrite, bainitic ferrite and/or retained austenite.
  • the complete absence of martensite can be achieved by sufficient carbon enrichment of austenite during the bainitic ferrite/retained austenite formation, so that the martensite start temperature of the remaining/retained austenite is lowered below ambient temperature and thus, during the final rapid quench of the hot forming process, no austenite transforms to martensite.
  • the martensite is 25-75.
  • Advantageously martensite is in the range 25-50.
  • Retained austenite 5-20.
  • the metastable retained austenite fraction ensures the balanced combination of strength and ductility properties.
  • Retained austenite enhances ductility partly through the stress / strain induced transformation effect, which manifests itself in an observed increase in uniform elongation and total elongation as the work hardening exponent increases to higher plastic strains. Below 5 vol.% the desired level of ductility and/or uniform elongation according to the present invention will not be achieved. The upper limit is set by the composition.
  • the retained austenite is in the range of 10-20.
  • Part or the totality of the process according to the present invention may be conducted in a controlled inert atmosphere of hydrogen, nitrogen, argon or any other inert gas in order to prevent oxidation and/or decarburisation of said steel.
  • Forming the blank into the desired part geometry may take place through a stamping/pressing operation, wherein the forming tool is represented by a mating punch and die with said steel positioned between the mating punch and die before punch and die close/mate so to stamp/press said steel into the desired part geometry.
  • the current process is also applicable to other forming methods such as roll forming.
  • an article produced by the present hot-forming method are characterized by ultra high tensile strength, complex shapes, exhibit minimum or no springback, and achieve a high increment in yield strength, especially for painting. Based on these advantages, excellent impact properties of the steel article are attained.
  • a cold rolled TRIP800 steel strip with gauge of 1.2 mm is provided by the conventional and known processes.
  • the steel strip contains in wt%: C:0.186; Mn:1.330; Si:1.670; P:0.008; Al:0.131; N:0.004; Nb:0.001; Ti:0.014; Cr:0.026, the balance being Fe and inevitable impurities.
  • the steel strip is heated at 1006 °C (T1), hold for 194.4 seconds (t1) at T1, and transferred from the furnace to the hot-forming tool within 11.3 seconds (t2). During the transfer the steel strip is cooled to 820 °C (T2) due to exposure to air cooling.
  • the steel strip is placed into the hot-forming tool having temperature T2, and simultaneously hot-stamped and die-quenched to 451 °C (T3) within 0.2 seconds (t3).
  • T3 the punch of the hot-forming tool is opened, i.e. the forming tool applied no pressure and cooled slowly within 53.2 seconds (t4) to 350 °C (T4).
  • T4 the forming tools is closed again, and the hot-formed article is quenched within 6.8 seconds (t5) to 100 °C (T5).
  • Table 1 gives the temperature values T1, T2, T3, T4 and T5 in °C and table 2 the time periods t1, t2, t3, t4 and t3 in seconds during the process according to the invention.
  • Table 1 T1 T2 T3 T4 T5 1006 820 451 350 100
  • Figures 4a-c illustrate scanning electron microscopy (SEM) images obtained with the JEOL JSM 6100 Scanning Electron Microscope. Samples were prepared by standard metallographic preparation procedures including etching in 2 % nital solution.
  • Figure 4a illustrates a representation of the global microstructure.
  • Figures 4b and 4c illustrate the finer details of the microstructure at higher magnification. Both figures 4b and 4c illustrate the proeutectoid ferrite matrix and the martensitic dispersions (as annotated in the figures).
  • Figure 4b also illustrates what is considered to be a retained austenite dispersion, while figure 4c also illustrates what is considered to be a bainitic ferrite dispersion.
  • the SEM images confirm a multi-phase TRIP steel-like microstructure.

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Claims (14)

  1. Procédé de formage à chaud d'une ébauche en acier en un article comprenant les étapes de :
    (c) formage d'une ébauche en acier chauffée dans un outil de formage à chaud en un article et trempe simultanée de celle-ci d'une température T2 à une température T3 en une période de temps t3, dans lequel l'outil de formage à chaud applique une pression sur l'ébauche en acier, dans lequel T3 est dans la plage de 400 à 600 °C, et dans lequel T2 est au-dessus de la température Ar1 spécifique chimique,
    (d) refroidissement de l'article dans l'outil de formage à chaud de la température T3 à une température T4 en une période de temps t4, dans lequel l'outil de formage à chaud n'applique aucune pression sur l'ébauche en acier,
    (e) trempe de l'article dans l'outil de formage à chaud de la température T4 à une température T5 en une période de temps t5, dans lequel l'outil de formage à chaud applique une pression sur l'ébauche en acier,
    dans lequel t3 est inférieur ou égal à 10 s, et dans lequel la vitesse de trempe entre la température T2 et T3 est supérieure ou égale à 50 °C/s,
    dans lequel T4 est dans la plage de 200 à 400 °C, et dans lequel la période de temps t4 est dans la plage de 60 à 180 s,
    dans lequel la vitesse de refroidissement entre la température T3 et T4 est dans la plage de 2 à 10 °C/s,
    dans lequel T5 est dans la plage de 20 à 200 °C, et la période de temps t5 est inférieure ou égale à 8 s,
    et dans lequel l'acier de l'ébauche en acier comprend, en % en poids, % en poids, les éléments suivants :
    C : 0,05 à 0,50
    Mn : 0,05 à 3,00,
    Si : ≤ 2,0,
    Al : ≤ 2,0,
    P : ≤ 0,17,
    S : ≤ 0,002,
    B : ≤ 0,01,
    N : ≤ 0,01,
    Nb : ≤ 0,1,
    Ti : ≤ 0,1,
    Cr : ≤ 1,0,
    et facultativement un ou plusieurs des éléments choisis parmi V, Mo, Co, W et d'autres éléments de terres rares, REM, de microalliages, en une quantité totale inférieure ou égale à 1,0 % en poids, le reste étant Fe et des impuretés inévitables.
  2. Procédé selon la revendication 1, dans lequel T4 est dans la plage de 300 à 380, de préférence 320 à 370 °C.
  3. Procédé selon la revendication 1 ou 2, dans lequel la période de temps t4 est de 100 à 180 s, de préférence de 120 à 180 s.
  4. Procédé selon l'une quelconque des revendications 1 à 3, dans lequel la vitesse de refroidissement entre la température T3 et T4 est dans la plage de 2 à 8 °C/s, de préférence de 2 à 5 °C/s.
  5. Procédé selon l'une quelconque des revendications 1 à 4, dans lequel T5 est dans la plage de 20 à 100 °C.
  6. Procédé selon l'une quelconque des revendications 1 à 5, dans lequel la période de temps t5 est inférieure ou égale à 4 s, de préférence inférieure ou égale à 2 s.
  7. Procédé selon l'une quelconque des revendications 1 à 6, dans lequel la vitesse de trempe entre la température T4 et T5 est supérieure ou égale à 50 °C/s, de préférence supérieure ou égale à 100 °C/s, plus préférablement supérieure ou égale à 150 °C/s.
  8. Procédé selon l'une quelconque des revendications 1 à 7, comprenant en outre - avant l'étape c) de formage à chaud, les étapes de
    (a) chauffage d'une ébauche en acier à une température T1 et maintien de l'ébauche chauffée à T1 pendant une période de temps t1 ;
    (b) transfert facultatif de l'ébauche chauffée à un outil de formage en un temps t2 pendant lequel la température de l'ébauche chauffée diminue de la température T1 à la température T2.
  9. Procédé selon la revendication 8, dans lequel T1 est d'au moins 730 °C, de préférence est dans la plage de 850 à 1150, plus préférablement de 980 à 1100, le plus préférablement de 1000 à 1050 °C et/ou dans lequel la période de temps t1 est inférieure ou égale à 15 min, de préférence de 3 à 5 min.
  10. Procédé selon l'une quelconque des revendications 1 à 9, dans lequel l'ébauche en acier est partiellement ou entièrement préformée.
  11. Procédé selon l'une quelconque des revendications 1 à 10, dans lequel l'ébauche, l'ébauche préformée ou l'article est pourvu d'un revêtement à base de zinc ou d'un revêtement à base d'aluminium-silicium ou d'un revêtement à base organique ou de tout autre revêtement conçu pour réduire l'oxydation et/ou la décarburation pendant le processus de formage à chaud.
  12. Procédé selon la revendication 11, dans lequel le revêtement à base de zinc est un revêtement contenant 0,5 à 3,8 % en poids d'Al, 0,5 à 3,0 % en poids de Mg, facultativement au plus 0,2 % en poids d'un ou plusieurs éléments supplémentaires, d'impuretés inévitables, le reste étant du zinc.
  13. Procédé selon l'une quelconque des revendications 1 à 12, dans lequel l'acier de l'ébauche en acier comprend, en % en poids, % en poids, les éléments suivants :
    C : 0,15 à 0,45, de préférence 0,15 à 0,35, le plus préférablement 0,15 à 0,25 ;
    Mn : 1,00 à 2,50, de préférence 1,00 à 1,50 ;
    Si : ≤ 1,0, de préférence 0,25 à 0,75, le plus préférablement 0,25 à 0,50 ;
    Al : ≤ 1,0, de préférence 0,25 à 0,75, le plus préférablement 0,25 à 0,50 ;
    P : ≤ 0,1, de préférence 0,02 à 0,07, le plus préférablement 0,02 à 0,05 ;
    S : ≤ 0,001 ;
    B : 0,001 à 0,005, de préférence 0,003 à 0,005 ;
    N : 0,001 à 0,005, de préférence 0,001 à 0,003 ;
    Nb : 0,02 à 0,04, de préférence 0,02 à 0,03 ;
    Ti : 0,02 à 0,04, de préférence 0,02 à 0,03 ;
    Cr : 0,2 à 0,4, de préférence 0,2 à 0,3 ;
    et facultativement un ou plusieurs des éléments choisis parmi V, Mo, Co, W et d'autres éléments de terres rares, REM, de microalliages, en une quantité totale inférieure ou égale à 1,0 % en poids, le reste étant Fe et des impuretés inévitables.
  14. Article en acier pouvant être obtenu par le procédé selon l'une quelconque des revendications 1 à 13, dans lequel l'article en acier a un allongement total compris dans la plage de 25 à 35 %, et dans lequel le produit RmxA de la contrainte à la rupture Rm et de l'allongement total A est compris dans la plage de 20 000 à 30 000 MPa %,
    dans lequel la microstructure de l'article en acier comprend, en somme additionnée jusqu'à 100, en fraction en volume, % en volume :
    • 0 à 75 de ferrite proeutectoïde, de préférence au plus 30, le plus préférablement 0,
    • 10 à 30 de ferrite bainitique, de préférence 20 à 30,
    • 5 à 20 d'austénite, de préférence 10 à 20,
    • 0 à 75 de martensite, de préférence 25 à 75, le plus préférablement 20 à 50.
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