EP4634415A1 - Method for producing a hot-rolled high-strength structural steel with improved formability and a method of producing the same - Google Patents

Method for producing a hot-rolled high-strength structural steel with improved formability and a method of producing the same

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Publication number
EP4634415A1
EP4634415A1 EP23834026.9A EP23834026A EP4634415A1 EP 4634415 A1 EP4634415 A1 EP 4634415A1 EP 23834026 A EP23834026 A EP 23834026A EP 4634415 A1 EP4634415 A1 EP 4634415A1
Authority
EP
European Patent Office
Prior art keywords
hot
steel
temperature
rolled
strip
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Pending
Application number
EP23834026.9A
Other languages
German (de)
French (fr)
Inventor
Shangping Chen
Radhakanta RANA
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Tata Steel Nederland Technology BV
Original Assignee
Tata Steel Nederland Technology BV
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Tata Steel Nederland Technology BV filed Critical Tata Steel Nederland Technology BV
Publication of EP4634415A1 publication Critical patent/EP4634415A1/en
Pending legal-status Critical Current

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Classifications

    • 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
    • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • 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
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • 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
    • 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 invention relates to a method of producing a hot-rolled high strength structural steel having improved formability and toughness and to a method for producing such a steel strip.
  • High strength steels with improved toughness are achieved advantageously by direct quenching.
  • the uniform elongation in uniaxial tensile testing i.e. work hardening capacity and formability are limited although the total elongation is generally acceptable.
  • This deficiency is an important factor limiting the wider and more demanding application of high strength steels because strain localization during fabrication or because of overloading in the final application can be detrimental to the integrity of the structure.
  • these steels are obtained by means of applying high cooling rates to a quenching temperature after hot rolling to obtain a microstructure consisting of partitioned martensite and retained austenite.
  • the intensive cooling may cause stress heterogeneity on a steel strip, which further generates great residual stress and influences steel strip shape.
  • US20060011274A1 discloses a two-step heat treatment, called quenching and partitioning (Q&P) which enables the production of steels with microstructures containing retained austenite.
  • Q&P quenching and partitioning
  • the steel is quenched from an austenization temperature around the Ac3 point of the steel to a quench stop temperature (QT) between the martensite start (Ms) and finish (Mf) temperatures.
  • QT quench stop temperature
  • Ms martensite start
  • Mf martensite start
  • PT partitioning temperature
  • the aim of the second step is to enrich the untransformed austenite with carbon through depletion of the carbon-super saturated martensite.
  • the invented two step Q&P is intended to improve the mechanical and forming related properties of thin sheet steels to be used in automotive applications. Such a two-step Q&P process is difficult to achieve for a production process of hot rolling as the process requires a rapid temperature increase to a PT temperature and maintaining for a period of time.
  • a one-step Q&P process can be used, that is, after the completion of the final rolling, coiling is carried out after on-line quenching to a certain temperature not greater than Ms.
  • US20140299237A1 discloses a one-step Q&P process for hot rolling product.
  • the steel has a composition C: 0.17 to 0.23%, Si: 1.4 to 2.0% or Si + Al: 1.2 to 2.0%, where Si is at least 0.4% and Al is at least 0.1%, preferably at least 0.8%, Mn: 1.4 to 2.3%, and Cr: 0.4 to 2.0% Mo: 0 to 0.7%, preferably 0.1 to 0.7%.
  • the hot-rolled steel strip is quenched at cooling rate of at least 20 °C/s to a quenching-stop temperature (QT), which QT is between Ms and Mf, temperatures, a partitioning treatment step for partitioning the hot-rolled steel in the temperature range from 250 to 500 °C and a cooling step for cooling the hot-rolled steel to room temperature by forced or natural cooling.
  • QT quenching-stop temperature
  • EP3235920A1 discloses a steel composition, C: 0.2-0.3%, Si: 1.0-2.0%, Mn: 1.5- 2.5%, P: ⁇ 0.015%, S: ⁇ 0.005%, Al: 0.5-1.0%, N: ⁇ 0.006%, Nb: 0.02-0.06%, Ti: ⁇ 0.03%, O: ⁇ 0.003%, and the balance being Fe and inevitable impurities.
  • the manufacture method comprises stepped cooling process to finally obtain a three-phase structure containing 0-25% of proeutectoid ferrite + 65-85% of (martensite + bainite) + 5-10% of residual austenite, with a yield strength of > 600 MPa, a tensile strength of > 1300 MPa, a yield ratio of ⁇ 0.5, and an elongation of > 10%.
  • US2016017449A1 discloses a hot rolled Q&P steel having a composition, C: 0.15% to 0.40%; Si: 1.0% to 2.0%; Mn: 1.5% to 3.0%; P: 0.015% or less; S: 0.005% or less; Al: 0.3% to 1.0%; N : 0.006% or less; Ti: 0.005% to 0.015%, and the balance being Fe and inevitable impurities.
  • the steel has a yield strength of > 700 MPa, a tensile strength of > 1300 MPa, and an elongation of > 10%.
  • a two-stepped cooling is applied to obtain a microstructure containing ferrite, martensite and retained austenite.
  • US20150203946A1 discloses a hot-rolled flat steel product having a product of Rm (tensile strength) and A80 (total elongation) of >18,000 MPa%, a composition including (in wt.) C: 0.10-0.60%, Si: 0.4-2.0%, Al: ⁇ 2.0%, Mn: 0.4-2.5%, Ni: ⁇ 1%, Cu: ⁇ 2.0%, Mo: ⁇ 0.4%, Cr: ⁇ 2%, Ti: ⁇ 0.2%, Nb: ⁇ 0.2%, V: ⁇ 0.5%, remainder Fe and unavoidable impurities, and a microstructure of bainite and residual austenite, wherein the microstructure includes > 60% bainite.
  • the steel is hot-rolled at a hot-rolling end temperature of > 880 °C, cooled with a cooling rate of > 5 °C/s to a coiling temperature between the martensite start temperature and 600 °C, coiled, and cooled in the coil while being held between the bainite start temperature and the martensite start temperature until > 60 % of the hot strip microstructure is bainite.
  • the steel according to the invention will possess high strength in combination with good elongation and good formability.
  • the high strength steel according to the invention has a tensile strength of at least 1300 MPa and a total elongation A50 of at least 7.0% and has a hole expansion capacity (HEC) of at least 20%.
  • HEC hole expansion capacity
  • composition range and role of the alloying elements for the present invention are as follows. All amounts of elements in compositions are given in weight % (wt.%), unless specified otherwise.
  • Carbon (C) is a necessary element in the steel according to the invention for strength and hardenability and enables the stabilization of retained austenite.
  • the strength of the steel, the amount of the retained austenite and also the average C content in the retained austenite increase as the carbon content is increased.
  • the carbon content is from 0.18 to 0.45. Carbon content lower than 0.18 makes it difficult to ensure required high strength. Carbon content exceeding 0.45 will impair the formability and the toughness of the steel.
  • the carbon content is preferably at least 0.19 and at most 0.40, more preferably at most 0.35 and even more preferably at most 0.32.
  • Si Silicon (Si) and Aluminium (Al) are compulsory elements in the steel.
  • Si has the main function to prevent carbon from precipitating in the form of iron carbides, most commonly cementite and to suppress decomposition of retained austenite.
  • Si is a ferrite stabilizer and is known to be quite effective for the retention of austenite during cooling after hot rolling.
  • a minimal amount of 0.50 Si is needed to suppress the carbide formation to an extent, but preferably 0.80 Si is needed to effectively suppress the carbide formation.
  • Si is at least 0.85. Addition of Si in excess of 2.50 is undesirable, as it deteriorates the formability and also causes Si-rich scales to result in poor surface quality. Therefore, the Si content is limited to 2.50 or less.
  • Si is at most 2.00, more preferably 1.80 and even more preferably at most 1.60.
  • Al is controlled to a value between 0.010 to 1.50.
  • a primary function of Al is to deoxidise the liquid steel before casting, therefore at least 0.01 is needed for this function.
  • Preferably at least 0.020 is added.
  • Al has a similar function as Si to prevent the formation of carbides and to stabilize the retained austenite, although Al is not as effective as Si and has no significant effect on strengthening. Too high levels of Al increase the austenite to ferrite transformation temperature Ar3 and therefore facilitate the formation of pro-eutectoid ferrite during cooling after hot rolling. In addition, the risk of cracking during casting increases as the Al content is increased. Therefore, the upper limit is set to 1.50.
  • Al is at most 1.30, more preferably 1.20 and even more preferably at most 1.15.
  • the effects of Al and of Si on the retardation of carbides and on the stabilization of austenite are very similar, they can be replaced by each other.
  • the sum total of Si and Al should be at least 1.20 % to reach the effects.
  • Al is less effective as Si, Al can be added to partially replace Si for this function. Therefore, the ratio of Si: Al is preferably higher than 2: 1, more preferably 3: 1 when they are used together.
  • the sum total of Si and Al should be at a value between 1.00 and 2.50, preferably between 1.20 and 2.20, more preferably between 1.20 and 2.00.
  • Mn Manganese
  • Mn is required to stabilize the retained austenite and to obtain hardenability of the steel. Mn is also added to balance the elevated phase transformation point as a result of high amounts of Si and Al. In order to obtain sufficient hardenability, Mn should be 1.20 or more, preferably 1.50 or more, more preferably 1.60 or more. A Mn content in excess of 3.50 will result in a decrease in the Ms point of the steel, thus the QT temperature is too low. In addition, a large amount of Mn promotes macro segregation, which results in unfavourable band structure formation, which is detrimental to ductility and formability. Therefore, the Mn content is limited to 3.50 or less, preferably to 3.00 or less, more preferably at most 2.80 and even more preferably at most 2.50.
  • Chromium (Cr) is a compulsory element in the steel, which increases the strength of the steel by increasing the hardenability, retards the formation of carbides and stabilizes the retained austenite. Cr may also change the state of carbide formation which results in good formability in the hard, high strength partitioned martensite. This effect is provided at a Cr content of 0.50 or more. Meanwhile, when Cr is added excessively, the effect is saturated. The upper limit of 2.50. Preferably Cr is at most 2.30, more preferably at most 2.00.
  • the sum of Cr and Mn should be 2.20 or more.
  • the upper limit of the sum of Cr and Mn is set to be 5.00 or less, preferably 4.50 or less for cost reasons.
  • Cr + Mn is at most 4.25.
  • the composition contains one or two elements selected from Ca and a rare earth metal (REM), in an amount consistent with a treatment for MnS inclusion control.
  • the rare earth element include scandium, yttrium, and lanthanide. It is recommended that for these elements to be useful they have to be present in amounts of 0.0003 or higher. However, when added excessively, the effect is saturated and the economic efficiency is reduced. Therefore, it is better to limit the amount to 0.0100 or less.
  • the following elements may be present as residual elements or optionally added as alloying elements to the composition of the steel according to the invention.
  • Residual elements are defined as elements which are not added on purpose to steel and which cannot be removed by simple metallurgical processes. Residual elements enter steel from impurities in ore, coke, flux and scrap; from these, scrap is considered to be the main source of residuals. Consequently, the level of residual elements in the Electric Arc Furnace process route (100% scrap based) is significantly higher than in the Basic Oxygen Steelmaking (BOS) process route. The most commonly found residuals are Cu, Ni, Cr, Mo, and Sn. The acceptance limits of these residuals depend mainly on product requirements.
  • the steel according to the invention is produced in a BOS-process route, in which case the allowable levels of residual elements for Cu, Ni, Mo and Sn are 0.040, 0.040, 0.020 and 0.020 wt.% respectively.
  • Copper may optionally be added to the steel according to the invention to facilitate the removal of high-Si scales formed in the hot rolling stage.
  • Cu can promote bainitic structures and may cause solid solution hardening.
  • Cu also reduces the amount of hydrogen penetrating into the steel and thus improves the delayed fracture characteristic.
  • the minimal content should be 0.05 or more. However, Cu causes hot shortness if an excess amount is added. Therefore, the Cu content should be 0.50 or less.
  • Nickel (Ni) and/or Molybdenum (Mo) may optionally be added to the steel according to the invention. When added, these elements improve the hardenability of the steel and facilitate the formation of bainite ferrite, and at the same time may stabilize retained austenite. Ni and/or Mo, when added, should be present at 0.05 or more to sufficiently obtain this effect. However, when added excessively, the cost of the steel is significantly increased. Therefore, Mo is set to be at most 0.50, and the amount of Ni is set to be at most 1.00. Ni may also be used to reduce the tendency of hot shortness when a high amount of Cu is added. For this purpose, the amount of Ni is preferably at least limited to 1/3 of the Cu content.
  • Niobium (Nb), Vanadium (V) and/or Titanium (Ti) may also be optionally added in the steel.
  • These elements can be used to refine the microstructure in the hot rolled intermediate product and the finished products. They possess a precipitation strengthening effect and may change the size or morphology of the bainitic ferrite and the partitioned martensite.
  • the allowable levels of Nb, Ti and V as a residual element is 0.005 for each. When these are added excessively, too much carbo-nitride is precipitated which result in deteriorating the workability. Therefore, they should be limited at 0.10 or less each, to prevent deterioration of the formability of the steel.
  • the sum of Ti + Nb + V should not exceed 0.250, preferably 0.100 for workability and cost.
  • B Boron
  • B is an optional element in the invented steel. It is a useful element in terms of suppressing formation and growth of polygonal ferrite from austenite grain boundaries. If present, the B content should be at least 0.0003, whereas the upper limit is set to 0.0050. A B content in steel exceeding 0.0050 deteriorates formability of a steel strip. For B to be able to perform this role, it is essential that no free nitrogen is present so that the formation of BN is avoided. This is where the nitrogen scavenging effect of certain elements such as titanium or aluminium plays a role.
  • Phosphorus (P) is an impurity element in steels. P segregates the grain boundaries and therefore decreases the workability and deteriorate impact resistance. Therefore, the amount of P should be suppressed to 0.050 or less.
  • the content of P is preferably minimized as far as possible.
  • the lower limit of the P content in steel is preferably around 0.005 because decreasing P below 0.005 significantly increases production costs.
  • S Sulphur
  • MnS sulphide-based inclusions
  • S it is desirable to reduce the amount of S as much as possible. Accordingly, S is 0.020 or less.
  • the lower limit of the S content in steel is preferably around 0.0005 because decreasing S content below 0.0005 significantly increases production costs.
  • N Nitrogen
  • the N content in steel is to be 0.010 or less, preferably 0.006 or less.
  • the practical lower limit of nitrogen content in steel is around 0.001 in view of production costs because decreasing nitrogen content in steel below 0.001 significantly increases production costs.
  • a suitable and practical minimum N content is between 0.002 to 0.005.
  • the manganese content is between 1.80 and 2.20 wt.%. At these values the properties were found to be well balanced.
  • the steel according to the invention has a complex phase microstructure, comprising in volume percent (vol.%): (partitioned martensite + bainitic ferrite) 85-95% and retained austenite 5-15%, wherein the retained austenite comprises an average C content of 0.90 % or more, the rest being S(pro-eutectoid ferrite + fresh martensite + carbides) 0-10%.
  • S(pro-eutectoid ferrite + fresh martensite + carbides) is at most 5%, more preferably at most 3%.
  • the matrix of a mixture of partitioned martensite (PM) and bainitic ferrite (BF) ensures a good balance in the strength and ductility of the steel.
  • the BF in these steels is generally present in the form of plates with an ultrafine grain size. These plates are typically about 5 pm long and about 200 nm thick.
  • the BF can be formed during cooling to the QT or formed during the cooling of the coil.
  • the PM has a similar substructure to BF but with a finer size of substructure as the formation temperature is lower.
  • the size of the ferrite plate and retained austenite is smaller in PM than in BF, thereby contributing to strength and the stability of the retained austenite (D. De Knijf, C. Fbjer, L. A. I. Kestens and R.
  • the sum of PM and BF should be at least 85%. If the sum of PM and BF is too low, too much fresh martensite or retained austenite may be produced which makes it impossible to simultaneously improve strength and formability. On the other hand, when the amount of PM and BF is too large, the amount of the retained austenite is not enough and the elongation of the steel is remarkably reduced. Therefore, the total fraction of PM and BF should be from 85 to 95%.
  • Retained austenite enhances ductility partly through the TRIP effect, which manifests itself in an increase in uniform elongation.
  • the volume fraction of retained austenite is controlled to be 5-15%. Below 5% the desired level of elongation will not be achieved. However, when the volume fraction is greater than 15%, the average carbon content in the retained austenite becomes too low. The retained austenite is not stable enough to contribute the TRIP effect and also the stretch flangeability will deteriorate.
  • the microstructure comprises 0 - 5% of proeutectoid ferrite.
  • Proeutectoid ferrite may form during cooling process if the cooling rate is too low.
  • the presence of the proeutectoid ferrite increases the total elongation but decreases strength and formability.
  • the amount of proeutectoid ferrite should preferably be 5% or less.
  • the microstructure comprises 0 - 5% of fresh martensite. Some amount of fresh martensite may be produced during final cooling if the bainitic transformation is incomplete during coiling. The presence of the fresh martensite further increases the strength of the steel. However, when the volume fraction of martensite is greater than 5%, the hole expansion capacity is remarkably reduced. Therefore, the volume fraction of martensite is preferably limited to a maximum amount of 5%.
  • the microstructure comprises 0 - 1.5% of carbides.
  • Carbides may be present in addition to proeutectoid ferrite, partitioned martensite, bainitic ferrite and fresh martensite and retained austenite when the coiling temperature is too high.
  • the formation of carbides reduces the amount of the retained austenite and the C content in the retained austenite, which deteriorate both the strength and ductility. If the amount of the carbides is higher than 1.5%, the minimum amount of retained austenite and the stability of the retained austenite cannot be ensured. Therefore, the amount of the carbide should preferably be limited to be less than 1.5%.
  • a slab having the above-described compositions is casted. It is possible to use a continuously cast slab or a slab produced by a thin slab caster or the like. It is also possible to cast a classical steel ingot which is converted into a slab by hot rolling.
  • the cast slab is reheated to a reheating temperature above the Ac3 point (the temperature at which the microstructure has become completely austenitic upon heating) and preferably above 1150 °C to completely dissolve the added alloying elements in the austenite phase to form a homogeneous solid solution prior to hot rolling and to enable a finish rolling temperature of Ar3 (the temperature at which austenite begins to transform into ferrite upon cooling a steel) or higher.
  • a reheating temperature above the Ac3 point (the temperature at which the microstructure has become completely austenitic upon heating) and preferably above 1150 °C to completely dissolve the added alloying elements in the austenite phase to form a homogeneous solid solution prior to hot rolling and to enable a finish rolling temperature of Ar3 (the temperature at which austenite begins to transform into ferrite upon cooling a steel) or higher.
  • the upper limit of the slab heating temperature is desirably 1350 °C or lower.
  • the slab or strip usually needs to be reheating or homogenised in temperature.
  • the reheating or homogenisation annealing is also referred to as reheating to a reheating temperature.
  • Hot rolling is completed by finishing hot rolling at a temperature being above the Ar3 point of the steel, preferably in the range of 800-950 °C, whereby the austenite phase is refined into fine grains.
  • the finish rolling temperature is lower than the Ar3 point, rolling with a two-phase region of ferrite and austenite is performed. The premature formation of ferrite can damage the formability of the hot rolled steel strip.
  • the finish rolling temperature is above 950 °C, the austenite grain size may become too large, which affect the size of the bainitic ferrite and martensite during cooling, thereby deteriorating the strength and ductility.
  • the Ar3 transformation point is estimated by the following expression using the content (wt.%) of each element and the strip thickness t (mm).
  • Ar3 901- 325C + 33Si - 92 (Mn + Ni/2 + Cr/2 + Cu/2 + Mo/2) + 52AI +0.35(t-8) (1)
  • the thickness of the hot rolled strips according to the invention is in the range from 2 to 15 mm. In a preferable embodiment the range is from 3 to 10 mm.
  • the hot rolled strip is then cooled at a cooling rate of 3 to 20 °C/s from the finish rolling temperature to a quenching temperature QT between Ms and Mf.
  • the cooling rate in the hot rolling strip mill can be adjusted by the number of water cooling nozzles in the run-out table (ROT), by forced air or by a mist spray.
  • the cooling rate in this context is the average cooling rate on the basis of surface strip temperature measurements and strip speed measurements.
  • the cooling rate can be controlled by controlling the cooling power of the ROT-cooling and the line speed.
  • the cooling rate is lower than 3 °C/s, pro-eutectoid ferrite and/or pearlite may form. The presence of the ferrite can increase the elongation but significantly reduce the strength and the formability.
  • the cooling rate is higher than 20 °C/s, the difference in temperature throughout the hot rolled strip becomes too large as the cooling rate gradient across the strip thickness, which leads to the difficulty in the QT control and which may produce distortion of the strip as the formation of internal stress.
  • the cooling rate is therefore in the range of 3 to 20 °C/s, and preferably of 5 to 20 °C/s. The gentler the cooling rate, the smaller the cooling rate gradient across the strip thickness.
  • the cooling rate is at most 18 °C/s, more preferably 17 °C/s or even 16 °C/s.
  • a preferred range of cooling rate is from 3 to 15 °C/s.
  • the quenching temperature (QT) should be below Ms, more preferably Ms - 50 °C, in the range of 150 to 350 °C. If the QT is too high, not enough partitioned martensite forms which will significantly reduce the yield strength and formability.
  • the QT should be above Mf, more preferably above Mf + 50 °C to ensure the presence of a sufficient amount of retained austenite, which is necessary to obtain the TRIP effect.
  • Coiling is then carried out at a coiling temperature ranging from 150 to 450 °C.
  • the coiling temperature is at most 400 °C and even more preferably at most 350 °C.
  • the CT is at least 175 °C and even more preferably at least 200°C.
  • partitioning of C from BF or martensite to the untransformed austenite occurs and some austenite may continue to transform to bainitic ferrite or martensite. If the CT is too high then carbides might be formed along the original austenite grain boundaries, which damage the formality. If the coiling temperature is too low then the carbon enrichment of the austenite may be insufficient during coil cooling process.
  • the coil cooling process is complete when the surface temperature of the out layer of the coil reaches ambient temperature.
  • the average cooling rate during coil cooling which is estimated by dividing the difference between the starting coiling temperature and ambient temperature by the coil cooling time, is preferably less than 0.02 °C/s, more preferably 0.01 °C/s to allow the retained austenite to stabilize and bainitic transformation to complete.
  • this cooling rate can be realized by transferring the coil to a furnace with a pre-set temperature or a programmed cooling or putting the coil to a hot water pool.
  • the coil size is adjusted to meet the cooling rate by natural cooling a coil in air.
  • a furnace with a programmed cooling means that the temperature in the furnace is programmed todecrease according to a pre-set program, e.g. to realise an average cooling rate during coil cooling of 0.02 °C/s.
  • Dilatometry was performed on the rolled samples of 10 mm (length) x 5 mm (width) x 5 mm (thickness) dimensions (length along the rolling direction). The samples were heated at a rate of 10 °C/s, held at the austenizing temperature of 950 °C for 2 minutes and then quenched at 60 °C/s to room temperature. The critical phase transformation points were determined from the dilatometry curves and are given in Table 1.
  • Room temperature tensile tests were performed in a Schenk TREBEL testing machine following NEN-EN10002-l :2001 standard to determine tensile properties (yield strength YS (MPa), ultimate tensile strength UTS (MPa), total elongation TE (%)). For each condition, three tensile tests were performed and the average values of mechanical properties are reported.
  • the microstructure was determined by optical microscopy (OM) using a commercially available image-processing program.
  • the polished cross section was etched by using a 3%-nital solution.
  • ferrite is revealed as white
  • pearlite and carbide are revealed as black
  • fresh martensite is revealed as straw-coloured
  • a mixture of bainitic ferrite and partitioned martensite is revealed as grey.
  • the volume fraction of retained austenite and carbides of the base microstructures were determined by X-ray diffraction (XRD) according to DIN EN 13925 on a D8 Discover GADDS (Bruker AXS).
  • XRD X-ray diffraction
  • the XRD measurements were conducted on a plane parallel to the street surface at 1 /4 thickness of the steel sheet.
  • the steel sheet was mechanically and chemically polished and was then analyzed by measuring the integral intensity of each of the (200) plane, (220) plane, and (311) plane of fee iron and that of the (200) plane, (211) plane, and (220) plane of bcc iron with an X-ray diffractometer using Co-Ka radiation.
  • the amount of retained austenite (RA) and the lattice parameter in the retained austenite were determined using Rietveld analysis.
  • the properties and/or the microstructures that do not meet the requirements are underlined. It is clear that a steel heat treated according to the invention has excellent balanced properties both in terms of strength, elongation and HEC value.
  • Figure 1 shows the fraction of martensite (f) transformed as a function of temperature. The fraction is determined from the dilatation test. The results that are presented apply to alloy 4 and show the determination of Ms and M90.
  • M5 is the temperatures at which 5% of the martensite has formed
  • M90 is the temperature at which 90% of the martensite has formed. In the context of this invention for practical reasons M5 is equalled to Ms (martensite start temperature) and M90 is equalled to Mf (martensite finishing temperature).
  • V2 cooling rate during coil cooling
  • C of RA carbon content of retained austenite

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Abstract

The invention relates to a method of producing a hot-rolled high strength structural steel having improved formability and toughness and to a method for producing such a steel strip.

Description

METHOD FOR PRODUCING A HOT-ROLLED HIGH-STRENGTH STRUCTURAL STEEL WITH IMPROVED FORMABILITY AND A METHOD OF PRODUCING THE SAME
Field of the invention
The invention relates to a method of producing a hot-rolled high strength structural steel having improved formability and toughness and to a method for producing such a steel strip.
Background of the invention
High strength steels with improved toughness are achieved advantageously by direct quenching. However, the uniform elongation in uniaxial tensile testing, i.e. work hardening capacity and formability are limited although the total elongation is generally acceptable. This deficiency is an important factor limiting the wider and more demanding application of high strength steels because strain localization during fabrication or because of overloading in the final application can be detrimental to the integrity of the structure.
According to the state of the art these steels are obtained by means of applying high cooling rates to a quenching temperature after hot rolling to obtain a microstructure consisting of partitioned martensite and retained austenite. The intensive cooling may cause stress heterogeneity on a steel strip, which further generates great residual stress and influences steel strip shape.
US20060011274A1 discloses a two-step heat treatment, called quenching and partitioning (Q&P) which enables the production of steels with microstructures containing retained austenite. Firstly, the steel is quenched from an austenization temperature around the Ac3 point of the steel to a quench stop temperature (QT) between the martensite start (Ms) and finish (Mf) temperatures. Then, the steel is either held at the QT or heated to a higher temperature, the so-called partitioning temperature (PT). The aim of the second step is to enrich the untransformed austenite with carbon through depletion of the carbon-super saturated martensite. In the Q&P process, the formation of carbides or bainite is intentionally suppressed, and the retained austenite is stabilized. The microstructure after Q&P treatment consists of ferrite, partitioned martensite and retained austenite or partitioned martensite and retained austenite. The invented two step Q&P is intended to improve the mechanical and forming related properties of thin sheet steels to be used in automotive applications. Such a two-step Q&P process is difficult to achieve for a production process of hot rolling as the process requires a rapid temperature increase to a PT temperature and maintaining for a period of time.
In the hot rolling process, a one-step Q&P process can be used, that is, after the completion of the final rolling, coiling is carried out after on-line quenching to a certain temperature not greater than Ms. US20140299237A1 discloses a one-step Q&P process for hot rolling product. The steel has a composition C: 0.17 to 0.23%, Si: 1.4 to 2.0% or Si + Al: 1.2 to 2.0%, where Si is at least 0.4% and Al is at least 0.1%, preferably at least 0.8%, Mn: 1.4 to 2.3%, and Cr: 0.4 to 2.0% Mo: 0 to 0.7%, preferably 0.1 to 0.7%. The hot-rolled steel strip is quenched at cooling rate of at least 20 °C/s to a quenching-stop temperature (QT), which QT is between Ms and Mf, temperatures, a partitioning treatment step for partitioning the hot-rolled steel in the temperature range from 250 to 500 °C and a cooling step for cooling the hot-rolled steel to room temperature by forced or natural cooling.
EP3235920A1 discloses a steel composition, C: 0.2-0.3%, Si: 1.0-2.0%, Mn: 1.5- 2.5%, P: < 0.015%, S: < 0.005%, Al: 0.5-1.0%, N: < 0.006%, Nb: 0.02-0.06%, Ti: < 0.03%, O: < 0.003%, and the balance being Fe and inevitable impurities. The manufacture method comprises stepped cooling process to finally obtain a three-phase structure containing 0-25% of proeutectoid ferrite + 65-85% of (martensite + bainite) + 5-10% of residual austenite, with a yield strength of > 600 MPa, a tensile strength of > 1300 MPa, a yield ratio of < 0.5, and an elongation of > 10%.
US2016017449A1 discloses a hot rolled Q&P steel having a composition, C: 0.15% to 0.40%; Si: 1.0% to 2.0%; Mn: 1.5% to 3.0%; P: 0.015% or less; S: 0.005% or less; Al: 0.3% to 1.0%; N : 0.006% or less; Ti: 0.005% to 0.015%, and the balance being Fe and inevitable impurities. The steel has a yield strength of > 700 MPa, a tensile strength of > 1300 MPa, and an elongation of > 10%. A two-stepped cooling is applied to obtain a microstructure containing ferrite, martensite and retained austenite.
US20150203946A1 discloses a hot-rolled flat steel product having a product of Rm (tensile strength) and A80 (total elongation) of >18,000 MPa%, a composition including (in wt.) C: 0.10-0.60%, Si: 0.4-2.0%, Al: < 2.0%, Mn: 0.4-2.5%, Ni: < 1%, Cu: < 2.0%, Mo: < 0.4%, Cr: < 2%, Ti: < 0.2%, Nb: < 0.2%, V: < 0.5%, remainder Fe and unavoidable impurities, and a microstructure of bainite and residual austenite, wherein the microstructure includes > 60% bainite. The steel is hot-rolled at a hot-rolling end temperature of > 880 °C, cooled with a cooling rate of > 5 °C/s to a coiling temperature between the martensite start temperature and 600 °C, coiled, and cooled in the coil while being held between the bainite start temperature and the martensite start temperature until > 60 % of the hot strip microstructure is bainite.
Objectives of the invention
It is an object of the invention to provide a hot rolled high strength steel sheet having improved formability.
It is also an object of the invention to provide a hot rolled high strength steel sheet having improved toughness.
It is also an object of the invention to provide a hot rolled high strength steel sheet that is suitable for engineering structural applications. Description of the invention
One or more of the objects is reached with the method according to claim 1. Preferred embodiments are provided by the dependent method claims. The invention is also embodied in the product according to claim 7. Preferred embodiments are provided by the dependent product claims.
Properties
The steel according to the invention will possess high strength in combination with good elongation and good formability. The high strength steel according to the invention has a tensile strength of at least 1300 MPa and a total elongation A50 of at least 7.0% and has a hole expansion capacity (HEC) of at least 20%.
Composition
The composition range and role of the alloying elements for the present invention are as follows. All amounts of elements in compositions are given in weight % (wt.%), unless specified otherwise.
C: 0.18-0.45
Carbon (C) is a necessary element in the steel according to the invention for strength and hardenability and enables the stabilization of retained austenite. The strength of the steel, the amount of the retained austenite and also the average C content in the retained austenite increase as the carbon content is increased. For the steels according to the invention the carbon content is from 0.18 to 0.45. Carbon content lower than 0.18 makes it difficult to ensure required high strength. Carbon content exceeding 0.45 will impair the formability and the toughness of the steel. The carbon content is preferably at least 0.19 and at most 0.40, more preferably at most 0.35 and even more preferably at most 0.32.
Si: 0.50 - 2.50;
Al: 0.01 - 1.50;
(Si + Al): 1.00 - 2.50;
Silicon (Si) and Aluminium (Al) are compulsory elements in the steel. Si has the main function to prevent carbon from precipitating in the form of iron carbides, most commonly cementite and to suppress decomposition of retained austenite. Si is a ferrite stabilizer and is known to be quite effective for the retention of austenite during cooling after hot rolling. A minimal amount of 0.50 Si is needed to suppress the carbide formation to an extent, but preferably 0.80 Si is needed to effectively suppress the carbide formation. Preferably Si is at least 0.85. Addition of Si in excess of 2.50 is undesirable, as it deteriorates the formability and also causes Si-rich scales to result in poor surface quality. Therefore, the Si content is limited to 2.50 or less. Preferably Si is at most 2.00, more preferably 1.80 and even more preferably at most 1.60. Al is controlled to a value between 0.010 to 1.50. A primary function of Al is to deoxidise the liquid steel before casting, therefore at least 0.01 is needed for this function. Preferably at least 0.020 is added. In addition, Al has a similar function as Si to prevent the formation of carbides and to stabilize the retained austenite, although Al is not as effective as Si and has no significant effect on strengthening. Too high levels of Al increase the austenite to ferrite transformation temperature Ar3 and therefore facilitate the formation of pro-eutectoid ferrite during cooling after hot rolling. In addition, the risk of cracking during casting increases as the Al content is increased. Therefore, the upper limit is set to 1.50. Preferably Al is at most 1.30, more preferably 1.20 and even more preferably at most 1.15.
As the effects of Al and of Si on the retardation of carbides and on the stabilization of austenite are very similar, they can be replaced by each other. The sum total of Si and Al should be at least 1.20 % to reach the effects. As Al is less effective as Si, Al can be added to partially replace Si for this function. Therefore, the ratio of Si: Al is preferably higher than 2: 1, more preferably 3: 1 when they are used together. The sum total of Si and Al should be at a value between 1.00 and 2.50, preferably between 1.20 and 2.20, more preferably between 1.20 and 2.00.
Mn: 1.20-3.50
Manganese (Mn) is required to stabilize the retained austenite and to obtain hardenability of the steel. Mn is also added to balance the elevated phase transformation point as a result of high amounts of Si and Al. In order to obtain sufficient hardenability, Mn should be 1.20 or more, preferably 1.50 or more, more preferably 1.60 or more. A Mn content in excess of 3.50 will result in a decrease in the Ms point of the steel, thus the QT temperature is too low. In addition, a large amount of Mn promotes macro segregation, which results in unfavourable band structure formation, which is detrimental to ductility and formability. Therefore, the Mn content is limited to 3.50 or less, preferably to 3.00 or less, more preferably at most 2.80 and even more preferably at most 2.50.
Cr: 0.40-2.50
Chromium (Cr) is a compulsory element in the steel, which increases the strength of the steel by increasing the hardenability, retards the formation of carbides and stabilizes the retained austenite. Cr may also change the state of carbide formation which results in good formability in the hard, high strength partitioned martensite. This effect is provided at a Cr content of 0.50 or more. Meanwhile, when Cr is added excessively, the effect is saturated. The upper limit of 2.50. Preferably Cr is at most 2.30, more preferably at most 2.00.
To ensure the hardenability and high strength, the sum of Cr and Mn should be 2.20 or more. The upper limit of the sum of Cr and Mn is set to be 5.00 or less, preferably 4.50 or less for cost reasons. Preferably Cr + Mn is at most 4.25. Ca + REM: 0.0003-0.0500
The composition contains one or two elements selected from Ca and a rare earth metal (REM), in an amount consistent with a treatment for MnS inclusion control. Examples of the rare earth element include scandium, yttrium, and lanthanide. It is recommended that for these elements to be useful they have to be present in amounts of 0.0003 or higher. However, when added excessively, the effect is saturated and the economic efficiency is reduced. Therefore, it is better to limit the amount to 0.0100 or less.
The following elements may be present as residual elements or optionally added as alloying elements to the composition of the steel according to the invention.
Residual elements (aka inevitable impurities, such as Cu, Ni, As, Pb, Sn, Mo, etc.) are defined as elements which are not added on purpose to steel and which cannot be removed by simple metallurgical processes. Residual elements enter steel from impurities in ore, coke, flux and scrap; from these, scrap is considered to be the main source of residuals. Consequently, the level of residual elements in the Electric Arc Furnace process route (100% scrap based) is significantly higher than in the Basic Oxygen Steelmaking (BOS) process route. The most commonly found residuals are Cu, Ni, Cr, Mo, and Sn. The acceptance limits of these residuals depend mainly on product requirements. Preferably the steel according to the invention is produced in a BOS-process route, in which case the allowable levels of residual elements for Cu, Ni, Mo and Sn are 0.040, 0.040, 0.020 and 0.020 wt.% respectively.
Copper (Cu) may optionally be added to the steel according to the invention to facilitate the removal of high-Si scales formed in the hot rolling stage. Cu can promote bainitic structures and may cause solid solution hardening. Cu also reduces the amount of hydrogen penetrating into the steel and thus improves the delayed fracture characteristic. When Cu is added as alloying element, the minimal content should be 0.05 or more. However, Cu causes hot shortness if an excess amount is added. Therefore, the Cu content should be 0.50 or less.
Nickel (Ni) and/or Molybdenum (Mo) may optionally be added to the steel according to the invention. When added, these elements improve the hardenability of the steel and facilitate the formation of bainite ferrite, and at the same time may stabilize retained austenite. Ni and/or Mo, when added, should be present at 0.05 or more to sufficiently obtain this effect. However, when added excessively, the cost of the steel is significantly increased. Therefore, Mo is set to be at most 0.50, and the amount of Ni is set to be at most 1.00. Ni may also be used to reduce the tendency of hot shortness when a high amount of Cu is added. For this purpose, the amount of Ni is preferably at least limited to 1/3 of the Cu content.
Niobium (Nb), Vanadium (V) and/or Titanium (Ti) may also be optionally added in the steel. These elements can be used to refine the microstructure in the hot rolled intermediate product and the finished products. They possess a precipitation strengthening effect and may change the size or morphology of the bainitic ferrite and the partitioned martensite. The allowable levels of Nb, Ti and V as a residual element is 0.005 for each. When these are added excessively, too much carbo-nitride is precipitated which result in deteriorating the workability. Therefore, they should be limited at 0.10 or less each, to prevent deterioration of the formability of the steel. In addition the sum of Ti + Nb + V should not exceed 0.250, preferably 0.100 for workability and cost.
Boron (B) is an optional element in the invented steel. It is a useful element in terms of suppressing formation and growth of polygonal ferrite from austenite grain boundaries. If present, the B content should be at least 0.0003, whereas the upper limit is set to 0.0050. A B content in steel exceeding 0.0050 deteriorates formability of a steel strip. For B to be able to perform this role, it is essential that no free nitrogen is present so that the formation of BN is avoided. This is where the nitrogen scavenging effect of certain elements such as titanium or aluminium plays a role.
Phosphorus (P) is an impurity element in steels. P segregates the grain boundaries and therefore decreases the workability and deteriorate impact resistance. Therefore, the amount of P should be suppressed to 0.050 or less. For the properties of the steel, the content of P is preferably minimized as far as possible. For production cost, the lower limit of the P content in steel is preferably around 0.005 because decreasing P below 0.005 significantly increases production costs.
Sulphur (S) is a harmful impurity in the steel which forms sulphide-based inclusions such as MnS, which may serve as a crack initiator, thereby deteriorating processability, impact resistance and formability of the steel. Therefore, it is desirable to reduce the amount of S as much as possible. Accordingly, S is 0.020 or less. For production cost, the lower limit of the S content in steel is preferably around 0.0005 because decreasing S content below 0.0005 significantly increases production costs.
Nitrogen (N) is inevitably present in the steel making process and allowable in the steels according to the invention in amounts up to 0.010 (=100 ppm). N in solid solution can markedly increases hardness and yield strength and decreases the tensile elongation. However, N deteriorates the toughness and formability of steel when an excess amount of 0.010 is added due to the formation of coarse nitride compounds or blowholes. Accordingly, the N content in steel is to be 0.010 or less, preferably 0.006 or less. The practical lower limit of nitrogen content in steel is around 0.001 in view of production costs because decreasing nitrogen content in steel below 0.001 significantly increases production costs. A suitable and practical minimum N content is between 0.002 to 0.005.
In an embodiment the hot-rolled high strength structural steel the composition of the hot rolled steel comprises one or more of the following:
- C between 0.19 and 0.40, and/or - Si between 0.80 and 2.00, and/or
- Al between 0.020 to 1.30, and/or
- Mn between 1.50 and 3.00, and/or
- Cr between 0.50 to 2.30.
The inventors found that these elements could be amended independent of each other, hence the "and/or".
In an embodiment the manganese content is between 1.80 and 2.20 wt.%. At these values the properties were found to be well balanced.
Microstructure
The steel according to the invention has a complex phase microstructure, comprising in volume percent (vol.%): (partitioned martensite + bainitic ferrite) 85-95% and retained austenite 5-15%, wherein the retained austenite comprises an average C content of 0.90 % or more, the rest being S(pro-eutectoid ferrite + fresh martensite + carbides) 0-10%. Preferably S(pro-eutectoid ferrite + fresh martensite + carbides) is at most 5%, more preferably at most 3%.
The matrix of a mixture of partitioned martensite (PM) and bainitic ferrite (BF) ensures a good balance in the strength and ductility of the steel. The BF in these steels is generally present in the form of plates with an ultrafine grain size. These plates are typically about 5 pm long and about 200 nm thick. The BF can be formed during cooling to the QT or formed during the cooling of the coil. The PM has a similar substructure to BF but with a finer size of substructure as the formation temperature is lower. The size of the ferrite plate and retained austenite is smaller in PM than in BF, thereby contributing to strength and the stability of the retained austenite (D. De Knijf, C. Fbjer, L. A. I. Kestens and R. Petrov, Factors influencing the austenite stability during tensile testing of Quenching and Partitioning steel determined via in-situ Electron Backscatter Diffraction, Mater. Sci. Eng. A 638 (2015) 219-227). To ensure the required mechanical properties as well as to reduce the amount of the fresh martensite formed during final cooling the sum of PM and BF should be at least 85%. If the sum of PM and BF is too low, too much fresh martensite or retained austenite may be produced which makes it impossible to simultaneously improve strength and formability. On the other hand, when the amount of PM and BF is too large, the amount of the retained austenite is not enough and the elongation of the steel is remarkably reduced. Therefore, the total fraction of PM and BF should be from 85 to 95%.
Retained austenite (RA) enhances ductility partly through the TRIP effect, which manifests itself in an increase in uniform elongation. The volume fraction of retained austenite is controlled to be 5-15%. Below 5% the desired level of elongation will not be achieved. However, when the volume fraction is greater than 15%, the average carbon content in the retained austenite becomes too low. The retained austenite is not stable enough to contribute the TRIP effect and also the stretch flangeability will deteriorate.
The microstructure comprises 0 - 5% of proeutectoid ferrite. Proeutectoid ferrite may form during cooling process if the cooling rate is too low. The presence of the proeutectoid ferrite increases the total elongation but decreases strength and formability. When the amount of proeutectoid ferrite is higher than 5%, it is difficult to achieve both a high tensile strength and high formability. Therefore, the amount of proeutectoid ferrite should preferably be 5% or less.
The microstructure comprises 0 - 5% of fresh martensite. Some amount of fresh martensite may be produced during final cooling if the bainitic transformation is incomplete during coiling. The presence of the fresh martensite further increases the strength of the steel. However, when the volume fraction of martensite is greater than 5%, the hole expansion capacity is remarkably reduced. Therefore, the volume fraction of martensite is preferably limited to a maximum amount of 5%.
The microstructure comprises 0 - 1.5% of carbides. Carbides may be present in addition to proeutectoid ferrite, partitioned martensite, bainitic ferrite and fresh martensite and retained austenite when the coiling temperature is too high. The formation of carbides reduces the amount of the retained austenite and the C content in the retained austenite, which deteriorate both the strength and ductility. If the amount of the carbides is higher than 1.5%, the minimum amount of retained austenite and the stability of the retained austenite cannot be ensured. Therefore, the amount of the carbide should preferably be limited to be less than 1.5%.
Process steps
First, a slab having the above-described compositions is casted. It is possible to use a continuously cast slab or a slab produced by a thin slab caster or the like. It is also possible to cast a classical steel ingot which is converted into a slab by hot rolling.
The cast slab is reheated to a reheating temperature above the Ac3 point (the temperature at which the microstructure has become completely austenitic upon heating) and preferably above 1150 °C to completely dissolve the added alloying elements in the austenite phase to form a homogeneous solid solution prior to hot rolling and to enable a finish rolling temperature of Ar3 (the temperature at which austenite begins to transform into ferrite upon cooling a steel) or higher. When the slab heating temperature is too low, a decrease in slab heating temperature leads to excessive increase in rolling load, and there are concerns of difficulty in rolling or causing a defective shape of the base steel sheet after rolling, and the like. Setting an excessively high heating temperature is not preferable in terms of being economical, and thus, the upper limit of the slab heating temperature is desirably 1350 °C or lower. In the case of hot rolling a thin slab or a cast strip immediately after casting then the slab or strip usually needs to be reheating or homogenised in temperature. However, in the context of this invention the reheating or homogenisation annealing is also referred to as reheating to a reheating temperature.
Hot rolling is completed by finishing hot rolling at a temperature being above the Ar3 point of the steel, preferably in the range of 800-950 °C, whereby the austenite phase is refined into fine grains. When the finish rolling temperature is lower than the Ar3 point, rolling with a two-phase region of ferrite and austenite is performed. The premature formation of ferrite can damage the formability of the hot rolled steel strip. On the other hand, if the finish rolling temperature is above 950 °C, the austenite grain size may become too large, which affect the size of the bainitic ferrite and martensite during cooling, thereby deteriorating the strength and ductility.
The Ar3 transformation point is estimated by the following expression using the content (wt.%) of each element and the strip thickness t (mm).
Ar3 = 901- 325C + 33Si - 92 (Mn + Ni/2 + Cr/2 + Cu/2 + Mo/2) + 52AI +0.35(t-8) (1)
The thickness of the hot rolled strips according to the invention is in the range from 2 to 15 mm. In a preferable embodiment the range is from 3 to 10 mm.
The hot rolled strip is then cooled at a cooling rate of 3 to 20 °C/s from the finish rolling temperature to a quenching temperature QT between Ms and Mf. The cooling rate in the hot rolling strip mill can be adjusted by the number of water cooling nozzles in the run-out table (ROT), by forced air or by a mist spray. The cooling rate in this context is the average cooling rate on the basis of surface strip temperature measurements and strip speed measurements. The cooling rate can be controlled by controlling the cooling power of the ROT-cooling and the line speed.
During said cooling, some bainitic ferrite and martensite forms. If the cooling rate is lower than 3 °C/s, pro-eutectoid ferrite and/or pearlite may form. The presence of the ferrite can increase the elongation but significantly reduce the strength and the formability. If the cooling rate is higher than 20 °C/s, the difference in temperature throughout the hot rolled strip becomes too large as the cooling rate gradient across the strip thickness, which leads to the difficulty in the QT control and which may produce distortion of the strip as the formation of internal stress. The cooling rate is therefore in the range of 3 to 20 °C/s, and preferably of 5 to 20 °C/s. The gentler the cooling rate, the smaller the cooling rate gradient across the strip thickness. Preferably the cooling rate is at most 18 °C/s, more preferably 17 °C/s or even 16 °C/s. A preferred range of cooling rate is from 3 to 15 °C/s.
The quenching temperature (QT) should be below Ms, more preferably Ms - 50 °C, in the range of 150 to 350 °C. If the QT is too high, not enough partitioned martensite forms which will significantly reduce the yield strength and formability. The QT should be above Mf, more preferably above Mf + 50 °C to ensure the presence of a sufficient amount of retained austenite, which is necessary to obtain the TRIP effect.
Coiling is then carried out at a coiling temperature ranging from 150 to 450 °C. Preferably the coiling temperature (CT) is at most 400 °C and even more preferably at most 350 °C. Preferably the CT is at least 175 °C and even more preferably at least 200°C. After coiling the coil cools further. During coil cooling, partitioning of C from BF or martensite to the untransformed austenite occurs and some austenite may continue to transform to bainitic ferrite or martensite. If the CT is too high then carbides might be formed along the original austenite grain boundaries, which damage the formality. If the coiling temperature is too low then the carbon enrichment of the austenite may be insufficient during coil cooling process. The coil cooling process is complete when the surface temperature of the out layer of the coil reaches ambient temperature. The average cooling rate during coil cooling, which is estimated by dividing the difference between the starting coiling temperature and ambient temperature by the coil cooling time, is preferably less than 0.02 °C/s, more preferably 0.01 °C/s to allow the retained austenite to stabilize and bainitic transformation to complete. Depending on the coil size, this cooling rate can be realized by transferring the coil to a furnace with a pre-set temperature or a programmed cooling or putting the coil to a hot water pool. Preferably the coil size is adjusted to meet the cooling rate by natural cooling a coil in air. A furnace with a programmed cooling means that the temperature in the furnace is programmed todecrease according to a pre-set program, e.g. to realise an average cooling rate during coil cooling of 0.02 °C/s.
Examples
6 steels having compositions according to the invention, as shown in Table 1, were cast into 25 kg ingots of 200 mm x 110 mm x 110 mm using vacuum induction. Then, the following process schedule was used to manufacture hot rolled strips of 5 mm thickness.
• Reheating of the ingots at 1225 °C for 2 hours;
• Rough rolling of the ingots from 110 mm to 35 mm;
• Reheating of the rough-rolled ingots at 1200 °C for 30 min;
• Hot rolling from 35 mm to 5 mm (35 - 27 - 19 - 11 - 7 - 5 mm) with the finish rolling temperature (FRT) at about 880 ± 40 °C;
• Run out table cooling according to the process parameters as given in Table 2;
• Transfer the strip to a preheated furnace at a chosen coiling temperature;
• Cooling the strip with controlled cooling rate to room temperature to simulate the coiling process;
• Pickling the hot rolled sheets in HCI at 85 °C to remove the oxide layers; Samples for microstructure observations, tensile tests and hole expansion tests were machined from hot rolled strips.
Dilatometry was performed on the rolled samples of 10 mm (length) x 5 mm (width) x 5 mm (thickness) dimensions (length along the rolling direction). The samples were heated at a rate of 10 °C/s, held at the austenizing temperature of 950 °C for 2 minutes and then quenched at 60 °C/s to room temperature. The critical phase transformation points were determined from the dilatometry curves and are given in Table 1.
Tensile tests - JIS5 test pieces (gauge length = 50 mm; width = 25 mm) were machined from the obtained hot-rolled sheets such that the tensile direction was parallel to the rolling direction. Room temperature tensile tests were performed in a Schenk TREBEL testing machine following NEN-EN10002-l :2001 standard to determine tensile properties (yield strength YS (MPa), ultimate tensile strength UTS (MPa), total elongation TE (%)). For each condition, three tensile tests were performed and the average values of mechanical properties are reported.
Hole Expansion Test (Stretch Flangeability Evaluation Test) - Test pieces for testing hole expandability (size: 90x90 mm) were sampled from the obtained rolled sheet. In accordance with The Japan Iron and Steel Federation Standards JFS T 1001, a 10 mm diameter punch hole was punched in the centre of the test piece and a 60° conical punch was pushed up and inserted into the hole. When a crack penetrated the sheet thickness, the hole diameter d (mm) was measured. The hole expansion ratio HEC (%) was calculated by the following equation: HEC(%) = {(d-do)/do}xlOO, with do being 10 mm.
The microstructure was determined by optical microscopy (OM) using a commercially available image-processing program. The polished cross section was etched by using a 3%-nital solution. Under optical microscopy, ferrite is revealed as white, pearlite and carbide are revealed as black, fresh martensite is revealed as straw-coloured, a mixture of bainitic ferrite and partitioned martensite is revealed as grey.
The volume fraction of retained austenite and carbides of the base microstructures were determined by X-ray diffraction (XRD) according to DIN EN 13925 on a D8 Discover GADDS (Bruker AXS). The XRD measurements were conducted on a plane parallel to the street surface at 1/4 thickness of the steel sheet. The steel sheet was mechanically and chemically polished and was then analyzed by measuring the integral intensity of each of the (200) plane, (220) plane, and (311) plane of fee iron and that of the (200) plane, (211) plane, and (220) plane of bcc iron with an X-ray diffractometer using Co-Ka radiation. The amount of retained austenite (RA) and the lattice parameter in the retained austenite were determined using Rietveld analysis. The C-content in the retained austenite is calculated using the formula (D. Dyson and B. Holmes, Effect of alloying additions on the lattice parameter austenite, J. Iron Steel Inst. 208 (1970) 469-474): C =(a[A] - 3.572 - 0.0012 Mn + 0.00157 Si - 0.0056 Al) / 0.033 (2) where a is the lattice parameter of the retained austenite in angstrom, and C, Mn, Si and Al are in wt.%.
The processing parameters, some material properties and the resulting tensile properties and HEC value are given in Table 2.
For the steels that are not processed according to the invention, the properties and/or the microstructures that do not meet the requirements are underlined. It is clear that a steel heat treated according to the invention has excellent balanced properties both in terms of strength, elongation and HEC value.
Brief description of the figure
The invention will now be explained by means of the following, non-limiting figure.
Figure 1 shows the fraction of martensite (f) transformed as a function of temperature. The fraction is determined from the dilatation test. The results that are presented apply to alloy 4 and show the determination of Ms and M90. M5 is the temperatures at which 5% of the martensite has formed, and M90 is the temperature at which 90% of the martensite has formed. In the context of this invention for practical reasons M5 is equalled to Ms (martensite start temperature) and M90 is equalled to Mf (martensite finishing temperature).
Table 1. The composition (in wt%) and critical phase transformation points of the tested steels
* al steels were treated with Ca 0.0005
Table 2. Process, microstructures and properties of the hot rolled steel strips
FRT: Finish Rolling Temperature Rm: ultimate tensile strength
VI: Cooling rate from FRT to QT A50: total elongation
QT: Quench Temperature HEC: hole expansion capacity
CT: Coiling Temperature RA: retained austenite
V2: cooling rate during coil cooling C of RA: carbon content of retained austenite
Rp: yield strength

Claims

1. A method for manufacturing a hot-rolled high-strength structural steel strip comprising the following steps:
• providing a steel ingot, slab, thin slab or a cast strip with a specific composition;
• heating or reheating the steel ingot, slab or strip to a reheating temperature in the range 950 to 1300 °C;
• optionally hot rolling the steel ingot to produce a slab;
• hot rolling the slab or strip in a temperature range in which the structure of the steel is entirely austenitic and finish rolling at a finish rolling temperature (FRT) above the Ar3- temperature;
• cooling the hot-rolled steel strip at a cooling rate of 3 - 20 °C/s to a quenching temperature (QT), which is between Ms and Mf;
• coiling the hot-rolled steel strip at a coiling temperature (CT) in the range 150 to 450 °C;
• allowing the coil to cool to room temperature by transferring the coil to a furnace with a pre-set temperature or a programmed cooling or by natural cooling, wherein the specific composition of the ingot, slab or strip comprises
• 0.18 - 0.45 wt.% C;
• 1.20 - 3.50 wt.% Mn;
• 0.40 - 2.50 wt.% Cr;
• 0.50 - 2.00 wt.% Si;
• 0.010 - 1.50 wt.% Al;
• (Si + Al) > 1.00%;
• (Mn + Cr) > 2.20% and (Mn + Cr) < 5.00%;
• at most 0.050 wt.% P;
• at most 0.020 wt.% S;
• at most 0.010 wt.% N;
• 0.0003 - 0.0500 wt.% of Z (Ca + REM); and optionally also comprising one or more of
• 0 - 0.50 wt.% Cu;
• 0 - 1.00 wt.% Ni;
• 0 - 0.50 wt.% Mo;
• 0 - 0.10 wt.% Nb;
• 0 - 0.10 wt.% Ti;
• 0 - 0.10 wt.% V;
• 0 - 0.0050 wt.% B the remainder being Fe and unavoidable impurities, and wherein the hot-rolled strip at room temperature has a complex phase microstructure, comprising in volume percent (vol.%): (partitioned martensite + bainitic ferrite) 85-95% and retained austenite 5-15%, wherein the retained austenite comprises an average C content of 0.90 % or more, at most 5% of proeutectoid ferrite, at most 5% of fresh martensite and at most 1.5% of carbides, wherein the tensile strength (Rm) of the hot-rolled high-strength structural steel is at least 1300 MPa, a total elongation (A50) of at least 7.0% and a hole expansion capacity (HEC) of at least 20%, wherein the tensile strength, the total elongation and the hole expansion capacity is measured as defined in the description.
2. Method according to claim 1 wherein coil is cooled to room temperature by transferring the coil to a furnace with a programmed cooling, preferably wherein the average cooling rate during coil cooling is less than 0.02 °C/s.
3. Method according to claim 1 or 2 wherein the composition of the hot rolled steel comprises one or more of the following
C between 0.19 and 0.40, and/or
Si between 0.80 and 2.00, and/or
Al between 0.020 to 1.30, and/or Mn between 1.50 and 3.00, and/or Cr between 0.50 to 2.30.
4. Method according to any one of claims 1 to 3 wherein the composition of the hot rolled steel comprises 1.8 to 2.20 % of Mn.
5. Method according to claims 1 to 4 wherein the thickness of the hot-rolled strip is between 2-15 mm, preferably between 3 to 10 mm.
6. Method according to any one of claims 1 to 5 wherein the quenching temperature is below (Ms - 50) °C.
7. Method according to any one of claims 1 to 6 wherein the quenching temperature is above (Mf + 50) °C.
8. Method according to any one of claims 1 to 7 wherein the coiling temperature is above 175 °C, preferably above 200 °C.
9. Method according to any one of claims 1 to 8 wherein the coiling temperature is below 400 °C, preferably below 350 °C.
10. A hot-rolled high strength structural steel strip, produced by hot-rolling an ingot, slab or strip comprising:
• 0.18 - 0.45 wt.% C;
• 1.20 - 3.50 wt.% Mn;
• 0.40 - 2.50 wt.% Cr;
• 0.50 - 2.00 wt.% Si;
• 0.010 - 1.50 wt.% Al;
• (Si + Al) > 1.00%;
• (Mn + Cr) > 2.20%;
• at most 0.050 wt.% P;
• at most 0.020 wt.% S;
• at most 0.010 wt.% N;
• 0.0003 - 0.0500 wt.% of Z (Ca + REM); and optionally also comprising one or more of
• 0 - 0.50 wt.% Cu;
• 0 - 1.00 wt.% Ni;
• 0 - 0.50 wt.% Mo;
• 0 - 0.10 wt.% Nb;
• 0 - 0.10 wt.% Ti;
• 0 - 0.10 wt.% V;
• 0 - 0.0050 wt.% B, the remainder being Fe and unavoidable impurities, wherein the hot-rolled strip has a complex phase microstructure, comprising in volume percent (vol.%): (partitioned martensite + bainitic ferrite) 85-95% and retained austenite 5-15%, wherein the retained austenite comprises an average C content of 0.90 % or more, at most 5% of proeutectoid ferrite, at most 5% of fresh martensite and at most 1.5% of carbides, wherein the tensile strength (Rm) of the hot-rolled high-strength structural steel is at least 1300 MPa, a total elongation (A50) of at least 7.0% and a hole expansion capacity (HEC) of at least 20%.
11. The hot-rolled high strength structural steel according to claim 10 wherein the S(pro- eutectoid ferrite + fresh martensite + carbides) is at most 5%, preferably at most 3%.
12. The hot-rolled high strength structural steel according to claim 10 or 11 wherein the composition of the hot rolled steel comprises one or more of the following C between 0.19 and 0.40, and/or
Si between 0.80 and 2.00, and/or
Al between 0.020 to 1.30, and/or
Mn between 1.50 and 3.00, and/or - Cr between 0.50 to 2.30.
13. The hot-rolled high strength structural steel according to any one of claims 10 to 121 wherein the composition of the hot rolled steel comprises 1.8 to 2.20 % of Mn.
14. The hot-rolled high strength structural steel according to any one of claims 10 to 13 wherein the allowable levels Cu, Ni, Mo and Sn respectively are 0.040, 0.040, 0.020 and 0.020 %.
15. Use of the hot-rolled steel according to any one of claims 10 to 114 for application in agricultural equipment, automotive components , construction and building components , infrastructure and infrastructure equipment, heavy vehicles equipment .
EP23834026.9A 2022-12-18 2023-12-16 Method for producing a hot-rolled high-strength structural steel with improved formability and a method of producing the same Pending EP4634415A1 (en)

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