US20110168300A1 - Manufacturing method for very high-strength cold-rolled dual-phase steel sheets and sheets so produced - Google Patents

Manufacturing method for very high-strength cold-rolled dual-phase steel sheets and sheets so produced Download PDF

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US20110168300A1
US20110168300A1 US12/993,498 US99349809A US2011168300A1 US 20110168300 A1 US20110168300 A1 US 20110168300A1 US 99349809 A US99349809 A US 99349809A US 2011168300 A1 US2011168300 A1 US 2011168300A1
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product
steel sheet
temperature
rolled
cold
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Antoine Moulin
Veronique Sardoy
Catherine Vinci
Gloria Restrepo Garces
Tom Waterschoot
Mohamed Goune
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ArcelorMittal Investigacion y Desarrollo SL
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ArcelorMittal Investigacion y Desarrollo SL
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D11/00—Continuous casting of metals, i.e. casting in indefinite lengths
    • B22D11/001—Continuous casting of metals, i.e. casting in indefinite lengths of specific alloys
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22D—CASTING OF METALS; CASTING OF OTHER SUBSTANCES BY THE SAME PROCESSES OR DEVICES
    • B22D7/00—Casting ingots, e.g. from ferrous metals
    • C—CHEMISTRY; METALLURGY
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/84—Controlled slow cooling
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
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    • C21D8/0236—Cold rolling
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
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    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the invention relates to the manufacture of cold-rolled and annealed sheets from steels known as “dual-phase” which have a very high strength and ductility for the manufacture of parts by shaping, in particular in the automobile industry.
  • Dual-Phase steels the structure of which comprises martensite, and possibly some bainite, in a ferritic matrix, have become widely used because they combine a high strength with high deformation capacity.
  • their yield strength is relatively low compared with their fracture strength, which gives them a very favorable yield strength/strength ratio during forming operations.
  • Their work-hardening ability is very high, which allows good deformation distribution in a collision and produces a much higher yield strength in a part after forming.
  • parts as complicated as those produced with conventional steels can be made, but with better mechanical properties, which enables a reduction in thickness to meet the same functional specification. In that way, these steels are an effective answer to the requirements of vehicle lightening and safety.
  • this type of steel In the field of hot-rolled (with a thickness for example of 1 to 10 mm) or cold-rolled (thickness for example of 0.5 to 3 mm) sheets, this type of steel especially finds applications for structural and safety parts for motor vehicles, such as crossmembers, side members, reinforcing parts, or even pressed steel wheels.
  • the document EP 0796928 A1 also describes cold-rolled Dual-Phase steels of which the strength is greater than 550 MPa, having the composition 0.05-0.3% C, 0.8-3% Mn, 0.4-2.5% Al, and 0.01-0.2% Si.
  • the ferritic matrix contains martensite, bainite and/or retained austenite.
  • the examples presented show that the strength does not exceed 660 MPa, even with a high carbon content (0.20-0.21%).
  • the document JP 11350038 describes Dual-Phase steels of which the strength is greater than 980 MPa, having the composition 0.10-0.15% C, 0.8-1.5% Si, 1.5-2.0% Mn, 0.01-0.05% P, less than 0.005% S, 0.01-0.07%Al in solution, and less than 0.01% N, also containing one or more of the following elements: 0.001-0.02% Nb, 0.001-0.02% V, 0.001-0.02% Ti.
  • This high strength is obtained however at the expense of a large addition of silicon which of course allows martensite to form, but can nevertheless result in the formation of surface oxides which negatively affect the dip coatability.
  • the object of the present invention is to provide a manufacturing method for very high-strength dual-phase steel sheets, cold-rolled, bare or coated, not having the disadvantages mentioned above.
  • the invention aims to provide Dual-Phase steel sheets having a mechanical strength between 980 and 1100 MPa together with a breaking elongation greater than 9% and good forming capacity, especially good bending capacity.
  • the invention also aims to provide a manufacturing method of which small variations of the parameters do not cause major changes to the microstructure or the mechanical properties.
  • the invention also aims to provide a steel sheet easily manufactured by cold-rolling, that is to say of which the hardness after the hot-rolling step is limited in such a way that the rolling strains remain moderate during the cold-rolling step.
  • the invention also aims to provide a steel sheet on which a metallic coating can be deposited, in particular by hot-dip galvanizing according to the usual methods.
  • the invention also aims to provide a steel having good weldability by means of the usual methods of assembling such as by resistance spot welding.
  • the invention also aims to provide an economical manufacturing method by avoiding the addition of costly alloying elements.
  • the subject of the invention is a cold-rolled and annealed Dual-Phase steel sheet having a strength between 980 and 1100 MPa, and a breaking elongation greater than 9%, of which the composition comprises, the contents being expressed by weight: 0.055% ⁇ C ⁇ 0.095%, 2% ⁇ Mn ⁇ 2.6%, 0.005% ⁇ Si ⁇ 0.35%, S ⁇ 0.005%, P ⁇ 0.050%, 0.1 ⁇ Al ⁇ 0.3%, 0.05% ⁇ Mo ⁇ 0.25%, 0.2% ⁇ Cr ⁇ 0.5%, it being understood that Cr+2Mo ⁇ 0.6%, Ni ⁇ 0.1%, 0.010 ⁇ Nb ⁇ 0.040%, 0.010 ⁇ Ti ⁇ 0.050%, 0.0005 ⁇ B ⁇ 0.0025%, and 0.002% ⁇ N ⁇ 0.007%, the remainder of the composition consisting of iron and the inevitable impurities resulting from smelting.
  • the composition of the steel contains, the content being expressed by weight: 0.12% ⁇ Al ⁇ 0.25%.
  • the composition of the steel contains, the content being expressed by weight: 0.10% ⁇ Si ⁇ 0.30%.
  • the composition of the steel preferably contains: 0.15% ⁇ Si ⁇ 0.28%. According to a preferred embodiment, the composition contains: P ⁇ 0.015%.
  • the microstructure of the steel sheet preferably contains a surface area fraction of 35 to 50% martensite.
  • the complement of the microstructure consists of a surface area fraction of 50 to 65% ferrite.
  • the complement of the microstructure consists of surface area fractions of 1 to 10% bainite and 40 to 64% ferrite.
  • the non-recrystallized ferrite surface area fraction compared to the whole of the ferritic phase is preferably less than or equal to 15%.
  • the steel sheet preferably has a ratio of its yield strength Re to its strength R m such that: 0.6 ⁇ Re/R m ⁇ 0.8.
  • the sheet is continuously galvanized. According to another particular embodiment, the sheet includes a galvannealed coating.
  • Another subject of the invention is a manufacturing method for a cold-rolled and annealed Dual-Phase steel sheet characterized in that a steel having a composition according to any one of the above specifications is supplied, then:
  • Another subject of the invention is a manufacturing method for a cold-rolled, annealed and galvanized Dual-Phase steel sheet characterized in that the heated and annealed product with a structure comprising austenite according to the above specification is supplied, then:
  • Another subject of the invention is a manufacturing method for a cold-rolled and galvannealed Dual-Phase steel sheet, characterized in that the heated and annealed product with a structure comprising austenite according to the above specification is supplied, then:
  • Another subject of the invention is a manufacturing method according to one of the above specifications, characterized in that the temperature T M is between 760 and 830° C.
  • the rate of cooling V R is greater than or equal to 15° C./s.
  • Another subject of the invention is the use of a steel sheet according to any one of the above specifications, or manufactured by a method according to any one of the above specifications, for the manufacture of structural or safety parts for motor vehicles.
  • FIG. 1 shows an example of a microstructure of a steel sheet according to the invention.
  • FIGS. 2 and 3 show examples of microstructures of steel sheets which are not according to the invention.
  • carbon plays an important part in the formation of the microstructure and affects the mechanical properties: below 0.055% by weight, the strength is unsatisfactory. Above 0.095%, an elongation of 9% cannot be guaranteed. The weldability is also reduced.
  • manganese is an element which increases the hardenability and reduces the precipitation of carbides.
  • a minimum content of 2% by weight is required to obtain the desired mechanical properties.
  • gamma-iron-forming quality results in the formation of a band structure which is too pronounced.
  • Silicon is an element which contributes to the deoxidizing of the liquid steel and the hardening in solid solution. This element also plays an important part in the formation of the microstructure by preventing the precipitation of carbides and by promoting the formation of martensite which is a component of the structure of Dual-Phase steels. It has a significant effect above 0.005%.
  • an increase in the silicon content reduces the dip-coating capacity by promoting the formation of oxides adhering to the surface of the products: its content must be limited to 0.35% by weight, and preferably 0.30%, to obtain good coatability. Silicon also reduces the weldability: a content less than 0.28% provides very good weldability as well as good coatability at the same time.
  • the ductility is reduced due to the presence of excess sulfides such as MnS which reduce the ductility, in particular during hole expansion tests.
  • Phosphorus is an element which hardens in solid solution but which reduces the spot weldability and the hot ductility, particularly due to its tendency to segregation at the grain boundaries or co-segregation with manganese. For these reasons, its content must be limited to 0.050%, and preferably 0.015%, in order to obtain good spot weldability.
  • Aluminum plays an important part in the invention by preventing the precipitation of carbides and by promoting the formation of martensitic components on cooling. These effects are obtained when the aluminum content is greater than 0.1%, and preferably when the aluminum content is greater than 0.12%.
  • AlN aluminum limits the grain growth during annealing after cold-rolling. This element is also used for deoxidizing the liquid steel in a quantity usually less than approximately 0.050%. In fact it is generally thought that higher contents increase the erosion of the refractories and the risk of blocking the nozzles. In excessive amounts, aluminum reduces the hot ductility and increases the risk of defects appearing in continuous casting. An effort is also made to limit inclusions of alumina, in particular in the form of clusters, with the aim of ensuring satisfactory elongation properties.
  • the inventors have demonstrated that, in combination with the other elements of the composition, a quantity of aluminum up to 0.3% by weight could be added without any negative effect on the other properties required, in particular with regard to the ductility, and would also make it possible to obtain the microstructural and mechanical properties sought. Above 0.3%, there is a risk of interaction between the liquid metal and the slag during continuous casting, which may result in the appearance of defects.
  • An aluminum content up to 0.25% by weight ensures the formation of a fine microstructure without large martensitic islands which would have a negative effect on the ductility.
  • the inventors have shown that, surprisingly, it was possible to obtain a high level of strength, between 980 and 1100 MPa, even in spite of limiting additions of aluminum and silicon. This is obtained by the particular combination of alloying or micro-alloying elements according to the invention, in particular by means of additions of Mo, Cr, Nb, Ti, and B.
  • molybdenum has a positive effect on the hardenability and retards the growth of ferrite and the appearance of bainite.
  • a content greater than 0.25% excessively increases the cost of the additions.
  • chromium due to its effect on the hardenability, also contributes to retarding the formation of proeutectoid ferrite. Above 0.5%, the cost of the addition is once again excessive.
  • chromium and molybdenum contents are such that: Cr+(2 ⁇ Mo) ⁇ 0.6%.
  • the coefficients in this relationship indicate the respective influences of these two elements on the hardenability for the purpose of promoting the production of a fine ferritic structure.
  • Titanium and niobium are micro-alloying elements used together according to the invention:
  • the above titanium and niobium contents make it possible to arrange that nitrogen is completely trapped as nitrides or carbonitrides, so much so that boron occurs in the free state and can have a positive effect on the hardenability.
  • the effect of boron on hardenability is crucial.
  • boron in fact makes it possible to control and limit the diffusive phase transformations (ferrite or pearlite transformation during cooling) and to form the hardening phases (bainite or martensite) required for obtaining high mechanical strength characteristics.
  • the addition of boron is therefore an important component of the present invention, and it also makes it possible to limit the addition of hardening elements such as Mn, Mo, and Cr and reduce the cost of the steel grade.
  • the minimum boron content to provide useful hardenability is 0.0005%. Above 0.0025%, the effect on the hardenability peaks and a negative effect on the coatability and the hot ductility is observed.
  • a minimum nitrogen content of 0.002% is required.
  • the nitrogen content is limited to 0.007% to prevent the formation of BN which would reduce the quantity of free boron required for the hardening of the ferrite.
  • An optional addition of nickel can be made so as to obtain extra hardening of the ferrite. This addition is however limited to 0.1% for cost reasons.
  • the implementation of the manufacturing method for a rolled sheet according to the invention includes the following successive steps:
  • This casting can be made in ingots or continuously as slabs having a thickness of the order of 200 mm.
  • the casting can also be carried out as thin slabs a few tens of millimeters thick or in thin strips between contra-rotating steel cylinders.
  • the cast semi-finished products are first brought to a temperature T R greater than 1150° C. so that at every point they reach a favorable temperature for the large deformations that the steel will undergo during rolling.
  • the austenite grains grow in an undesirable manner.
  • the only precipitates that can effectively control the austenite grain size are the nitrides of titanium, and the heating temperature should be limited to 1250° C. in order to maintain a fine austenite grain size at this stage.
  • the hot-rolling step for these semi-finished products starting at more than 1150° C. can be done directly after casting so that an intermediate heating step is not required in this case.
  • the semi-finished product is hot-rolled in a temperature range in which the structure of the steel is fully austenitic: if T FL is less than the start temperature of austenite transformation on cooling A r3 , the ferrite grains are work-hardened by the rolling and the ductility is reduced.
  • T FL is less than the start temperature of austenite transformation on cooling A r3
  • the ferrite grains are work-hardened by the rolling and the ductility is reduced.
  • an end-of-rolling temperature greater than 850° C. will be selected.
  • the hot-rolled product is next coiled at a temperature T bob between 500 and 570° C.: this temperature range makes it possible to obtain a complete bainite transformation during the nearly isothermal holding time associated with coiling. This range results in a morphology of Ti and Nb precipitates which is fine enough to make use of their hardening power during later steps of the manufacturing method.
  • a coiling temperature greater than 570° C. results in the formation of coarser precipitates, of which the coalescence during continuous annealing significantly reduces the effectiveness.
  • the hot-rolled product is descaled using a method known in its own right, then a cold-rolling is carried out with a reduction of preferably between 30 and 80%.
  • the cold-rolled product is heated, preferably in a continuous annealing plant, at an average rate of heating V C between 1 and 5° C./s. Combined with the annealing temperature T M below, this rate of heating range produces a non-recrystallized ferrite fraction less than or equal to 15%.
  • the heating is carried out at an annealing temperature T M between the temperature A c1 (start temperature of allotropic transformation on heating)+40° C., and A c3 (end temperature of allotropic transformation on heating) ⁇ 30° C., that is to say in a specific temperature range within the intercritical range: when T M is less than (A c1 +40° C.), the structure can also include zones of non-recrystallized ferrite of which the surface area fraction can reach 15%. This non-recrystallized ferrite fraction is calculated in the following manner: having identified the ferritic phase in the microstructure, the non-recrystallized ferrite surface area percentage compared with the whole of the ferritic phase is quantified.
  • An annealing temperature T M according to the invention produces enough austenite to form martensite later on cooling in such a quantity that the desired characteristics are achieved.
  • a temperature T M less than (A c3 ⁇ 30° C.) also ensures that the carbon content of the islands of austenite formed at the temperature T M does in fact result in a later martensite transformation: when the annealing temperature is too high, the carbon content of the islands of austenite becomes too low, which results in a later unfavorable transformation to bainite or pearlite. What is more, too high a temperature results in an increase in the size of the niobium precipitates which lose part of their hardening capacity. The final mechanical strength is then reduced.
  • a temperature T M between 760° C. and 830° C. will preferably be selected.
  • a minimum holding time t M of 30 s at the temperature T M allows the carbides to dissolve, and a partial transformation to austenite occurs. After a time of 300 s the effect peaks.
  • a holding time greater than 300 s is also hardly compatible with the productivity requirements of continuous annealing plants, in particular the pass speed.
  • the holding time t M is between 30 and 300 s.
  • the following steps of the method differ according to whether uncoated steel sheet, or continuous hot-dip galvanized steel sheet, or galvannealed steel sheet is being manufactured:
  • This cooling can be carried out starting from the temperature T M in one or more steps and can use in the latter case various cooling methods such as cold or boiling water baths, water or gas jets. These possible accelerated cooling methods can be combined so as to obtain a complete transformation of austenite to martensite. After this martensite transformation, the steel sheet is cooled to the ambient temperature.
  • the microstructure of the cooled bare sheet then consists of a ferritic matrix with islands of martensite of which the surface area fraction is between 35 and 50%, and which is free of bainite.
  • the inventors have also observed that small variations of the manufacturing parameters, in the conditions defined according to the invention, do not cause major changes to the microstructure or the mechanical properties, which is an advantage for the stability of the characteristics of the industrial products manufactured.
  • Cast semi-finished products corresponding to the compositions above were heated to 1230° C. then hot-rolled to a thickness of 2.8-4 mm in a temperature range in which the structure is entirely austenitic.
  • the manufacturing conditions of these hot-rolled products (end-of-rolling temperature T FL , coiling temperature T bob ) are shown in table 2.
  • the hot-rolled products were next descaled then cold-rolled to a thickness of 1.4 to 2 mm which is a reduction of 50%.
  • some steels were subjected to different manufacturing conditions.
  • the references IX 1 , IX 2 and IX 3 designate for example three steel sheets manufactured under different conditions starting with the steel composition IX.
  • the sheets were hot-dip galvanized in a bath of zinc at a temperature T Zn of 460° C., others were also subjected to galvannealing treatment.
  • Table 3 shows the manufacturing conditions of the sheets annealed after cold-rolling:
  • the microstructure of the steels, of which the matrix is ferritic, has also been determined.
  • the surface area fractions of bainite and martensite have been quantified after attack with Picral and LePera reagents respectively, followed by image analysis using AphelionTM software.
  • the surface area fraction of non-recrystallized ferrite was also determined using optical microscopy and scanning electronic microscopy observations in which the ferritic phase was identified, then the recrystallized fraction in this ferritic phase was quantified.
  • the non-recrystallized ferrite occurs generally in the form of islands elongated by the rolling.
  • the bending capacity was quantified in the following manner: sheets were bent back on themselves several times. In this way, the bending radius gets smaller each time. The bending capacity is then evaluated by noting the presence of cracks at the surface of the folded block, the score being expressed from 1 (low bending capacity) to 5 (very good capacity). Results which scored 1-2 are considered unsatisfactory.
  • the steel sheets according to the invention have a set of microstructural and mechanical characteristics which enable the advantageous manufacture of parts, especially for structural applications: strength between 980 and 1100 MPa, ratio R e /R m between 0.6 and 0.8, breaking elongation greater than 9%, good bending capacity.
  • FIG. 1 illustrates the morphology of the steel sheet IX 1 , in which all the ferrite is recrystallized.
  • the sheets according to the invention have good weldability, especially by resistance spot welding, the carbon equivalent being less than 0.25.
  • spot-welding weldability current range as defined by the IS018278-2 standard, is very wide, of the order of 3500A. It is increased compared with a reference steel of the same grade.
  • cross-tensile tests or shear-tensile tests carried out on spot welds on sheets according to the invention reveal that the strength of these spot welds is very high in terms of mechanical properties.
  • FIG. 2 illustrates the microstructure of the steel sheet IX 3 : note the presence of non-recrystallized ferrite in the form of elongated islands (marked (A)) coexisting with recrystallized ferrite and martensite, the latter component appearing darker in the micrograph.
  • a Scanning Electronic Microscopy micrograph FIG. 3 ) clearly differentiates the zones of non-recrystallized ferrite (A) from the recrystallized ones (B).
  • Sheet IX 5 is a galvannealed sheet annealed at too high a temperature T M : the carbon content of the austenite at high temperature is then too low and the appearance of bainite is promoted to the detriment of the formation of martensite. There is also coalescence of the niobium precipitates, which causes a loss of hardening. The strength is then unsatisfactory, the ratio Re/R m being too high.
  • the galvannealed sheet IX 7 was cooled at too slow a rate V R after the annealing step: the transformation of the austenite formed to ferrite during this cooling step is then excessive, the steel sheet containing in the final stage too high a bainite fraction and too low a martensite fraction, which results in unsatisfactory strength.
  • the composition of the steel sheet R does not correspond to the invention, its carbon content being too high, and its manganese, aluminum, niobium, titanium, and boron contents being too low. Consequently, the martensite fraction is so low that the mechanical strength is unsatisfactory.
  • the steel sheets according to the invention will be beneficially used for the manufacture of structural or safety parts in the automobile industry.

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RU2470087C2 (ru) 2012-12-20
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US20160222486A1 (en) 2016-08-04

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