WO2012120020A1

WO2012120020A1 - Process for producing high strength formable steel and high strength formable steel produced therewith

Info

Publication number: WO2012120020A1
Application number: PCT/EP2012/053856
Authority: WO
Inventors: David Neal Hanlon; Stefanus Matheus Cornelis VAN BOHEMEN
Original assignee: Tata Steel Nederland Technology Bv
Priority date: 2011-03-07
Filing date: 2012-03-07
Publication date: 2012-09-13
Also published as: EP2683839A1; EP2683839B1; ES2535420T3

Abstract

The invention relates to a method for producing a steel strip composite and to a steel strip composite produced by said method.

Description

Process for producing high strength formable steel and high strength formable steel produced therewith.

FIELD OF THE INVENTION

BACKGROUND OF THE INVENTION

Enhancement of ductility at elevated strength is desirable for widespread markets. In the automotive industry in particular, where legislation is driving improvements in fuel economy and safety, there is a move towa rds stronger, formable high strength steels. Hig h strength a nd ultrahig h strength strip steel provides a utomotive man ufacturers potentia l for down weig hting of the body in white and the opportunity for countering weight increases arising from the move to electric and hybrid vehicles. In addition high and ultrahigh strength steels play a critical role in determining the crash worthiness of modern passenger vehicles.

Extensive a pplication of high strength a nd ultra high strength steel requires, in many cases, levels of formability that are higher than can be expected for conventional carbon-manganese steels. Enhancement of residual ductility in formed parts is beneficial for integrity in crash. As a first approximation tensile elongation may be considered a simple measure of both formability and impact integrity. Consequently, such effort has been given to developing advanced high strength steels (AHSS) with optimised tensile ductility.

In the case of current commercia l AHSS, enha ncement of d uctility is realised either by using composite effects (Dua l Phase or DP steels) or by using deformation induced transformation of retained austenite (TRIP steels). Both composite hardening and transformation induced plasticity mechanisms can measurably enhance the ductility of predominantly ferritic steels. However, the enha ncement of ductility which may be achieved i n th is way is l imited to a rou nd 1.5 (for DP) to 2 (for TRIP) times that of conventional C-Mn steels at equivalent strength . Furthermore, the strength level that may be achieved in commercial, ferritic-matrix DP and TRIP steels is also limited (TRIP to around 800MPa and DP to around lOOOMPa).

Extreme enhancement of tensile ductility may also be achieved by utilising high a lloy (predomina ntly manga nese) additions to sta bilise austenite (Twinning Induced Plasticity or TWIP steels) . Austenite is inherently more ductile than ferrite and twinning provides a very effective work hardening mechanism. Such steels may achieve very high elongations (typically 50%) at very high strengths (typically lOOOMPa). The improvement in elongation may be typically 5x that of a conventional C-Mn steel. However, the yield strength is comparatively low and large strains need to be uniformly imposed to achieve high strength in the formed component. Furthermore, the extremely high levels of alloy make large scale production over conventional process problematic.

Practical experience gained during the implementation of these first and second generations of AHSS have revealed that forming and performance parameters beyond those measured in the tensile test can represent significant barriers to implementation. In particular, although exhibiting high uniform and total elongations, and consequently high resistance to necking during stretch forming, mixed microstructures comprising a distribution of hard phase in a matrix of soft ferrite may be highly susceptible to cracking at stretched edges. Low stretched edge formability is known to impose a practical limit to the formability of otherwise highly ductile AHSS.

A new generation of AHSS that exhibits an advantage in strength ductility-balance with respect to carbon manganese steels but at lower levels of alloy, and consequently lower cost and greater processability, is emerging. For this generation of steels more attention is a lso bei ng pa id to other forma bility pa ra meters such as stretched edge d ucti lity (hole expa nsitivity) a nd benda bility. In most cases, mixed microstructures comprising hard majority phases such as martensite or bainite are used to develop high to ultrahigh strengths. Such hard, uniformly fine, microstructures tend to exhibit good hole expansitivity. Alloy compositions fall into two basic categories: Medium Manganese compositions (Mn typically 7wt%) and Medium Carbon, carbon-manganese steels with additions of Si and or Al.

In the case of the high Mn TRIP steels austenite is simply chemically stabilised by large additions of Mn. Although clearly differentiated strength-ductility combinations can be ach ieved i n th is way, processi ng ha s proven to be difficult since hard, brittle martensitic phases are developed in intermediate product. This renders further processi ng , such a s cold rol l i ng , d ifficu lt at d i mensions releva nt to com mercia l processing. Furthermore, alloy costs are high due to the high manganese content.

Therefore the problem arises that it is difficult to provide a steel strip having improved strength and edge ductility.

It is an object of this invention to provide a method for improving the strength and edge ductility of steels.

It is a further object of this invention to provide a steel having improved strength and ductility in combination with an low allow content.

SUMMARY OF THE INVENTION According to a first aspect of the invention there is provided a method for producing a quenched and pa rtitioned steel by provid ing a cold rolled a nd a nnea led steel strip containing (in weight %) :

0.18 - 0.4% C

1.5 - 4.0% Mn

0.5 - 2.0% Si

0 - 1.5% Al

0 - 0.5% Mo

0 - 0.5% Ti

• 0 - 0.4% V

0 - 0.010% Nb

0 - 0.005% B

0 - 0.015% N

0 - 0.08% P

. 0 - 0.01% S

0 - 0.06 % Sb

0 - 0.05 % Ca

0 - 1.0%Cr

0 - 1.0% Ni

• the remainder being iron and unavoidable impurities,

wherein the annealing process comprises the following steps:

(i) reheating the cold rolled strip to an annealing temperature Ti of between Ac₃-40 and Ac₃+80;

(ii) holding the strip at Ti for an annealing time ti of between 10 and 200 seconds;

(iii) cooling the annealed strip at a cooling rate CR1 to a quench temperature T₂ for producing a microstructure in the strip comprising a martensite fraction and a retained austenite fraction;

(iv) partitioning annealing the cooled strip at a temperature T₃ for enriching the austenite in carbon by repartitioning the carbon from the martensite to the austenite fraction for a repartitioning time t₂ of between 20 and 500 seconds;

(v) cooling the strip at a cooling rate CR2 to ambient temperature.

According to the invention a cold-rolled strip is provided by the conventional and known processes of casting, hot-rolling and cold-rolling. The casting process may involve thick slab casting (sla b thickness between 150 a nd 350 mm), thin sla b-casting (slab thickness below 150 mm, usually between 50 and 100 mm) or even strip casting. Cold rolling is also a conventional and known process. The composition of the steel according to the invention is a balanced composition comprising medium carbon and moderate manganese additions. The cold rolled steel is heated to form austenite, either partially or fully, followed by quenching to a temperature between the ma rtensite start temperature (M_s) a nd the ma rtensite finish temperature (M_f), thereby creating a controlled amount of martensite and retained austenite. The manganese additions result in stabilisation of an austenite fraction during cooling from the annea ling temperature and the subsequent ca rbon enrichment further sta bilises the austenite fraction. Combined stabilisation via C and Mn additions enables alloying with either element to be restricted to reasonable limits lead ing to cost and processability advantages. The steel is then subjected to a thermal treatment to partition carbon from the martensite into the austenite. The carbon enrichment of the austenite fraction is achieved by partitioning from martensite or low temperature bainitic transformation or preferably a combination of both. By this pa rtitioning, the formation of ca rbides is su pp ressed a nd the a uste n ite is sta b i l ised rather tha n decom posed . Combined stabilisation with partitioning and bainitic transformation enables the amount of austenite and the microstructure in which it is embedded to be optimised. The bainitic transformation also leads to ca rbon enrichment of the remaining a ustenite because the formation of carbides is suppressed. All compositions are given in weight percentages, unless otherwise indicated. The final microstructure of the steel comprises martensite, bainite and carbon- enriched austenite and, if Ti < Ac₃, equiaxed ferrite. The stabilisation of austenite results in the steel exhibiting improved ductility relative to traditional high strength steels.

Ca rbon (C) provides solid solution strengthening, enhances hardenability (thus enabling avoidance of high temperature transformations at cooling rates achievable in conventional annealing lines) and, when dissolved in austenite, promotes the retention of austenite at room temperature. Above 0.4wt% C the propensity for formation of brittle high carbon martensite increases.

Ma nga nese (M n) del ivers substa ntia l solid sol ution strengthen ing, sta bil ises austenite, thus promoting its retention at room temperature, and enhances hardenability promoting the formation of hard transformation products at cooling rates achievable in conventional annealing lines. A preferable upper limit for the manganese content is 3.5%.

Sil icon (Si) add ition provides sol id solution strengthen ing th us ena bl ing the attainment of high strength and promotes the stabilisation of austenite. Si very effectively retards the formation of carbides during overaging thus keeping carbon in solution for stabilisation of austenite. Ferrite and hard phases such as bainite and martensite exhibit improved ductility in the absence of carbides. Free carbon, not trapped in carbides, may be partitioned to austenite. For acceptable coatability the imposed addition of Si should be below a certain level : Si may be added i n the ra nge of between 0.5 to lwt% when combined with Al addition . In the absence of Al additions beyond the level needed for deoxidation Si should be maintained in the range of between 1 to 2wt%.

Aluminium (Al) is usually already added in small quantities of at most 0.1% and preferably at most 0.05% to liquid steel for the purpose of deoxidation by forming alumina. Ideally, the total aluminium content in the steel is between 0.01 and 0.08% if aluminium is only added for desoxidation. In the right quantity it also provides an acceleration of the bainite transformation. Al also retards the formation of carbides thus keeping carbon in solution for pa rtitioning to a ustenite a nd promoting the sta bi lisation of a ustenite. A maximum value of 1.5wt% is imposed for castability purposes because higher Al-contents lead to poisoning of casting mould slag and consequently an increase in mould slag viscosity leading to incorrect heat transfer and lubrication during casting. Aluminium alone delivers low strength . Therefore, if Al is used a bove levels required for deoxidation it should always be in combination with Si.

Niobium (Nb), if added, is added in small amounts of up to 0.1% or more preferably of u p to 0.05wt%. It is added for austenitic grain refinement during hot rolling. If combined with a suitable rolling schedule, it promotes transformation on the run-out-table, and thus finer and a more homogeneous grain size in the hot-rolled intermediate product which is then subsequently cold-rolled.

Titanium (Ti) ca n be used to form fine precipitates in the ferritic component of mixed microstructures thus increasing strength and promoting uniformity of strength at the microstructural scale and in turn good stretched edge ductility. Preferably Ti is below 0.1%.

Molybdenum (Mo) is used to increase hardenability thus retarding the formation of high temperature transformation products on cooling to the quench temperature. Additions of Mo may therefore allow the utilisation of lower cooling rates more readily achieved in full-scale production lines or to achieve the desired structure and properties in heavier gauge strip. The use of Mo is to be preferred over the use of Cr when UHS is required since in addition to enhancing hardenability Mo provides additional solid solution strengthening. Mo is also known to retard coarsening of fine strengthening precipitates thus promoting thermal stability precipitation strengthening variants. For reasons of cost less than 0.5wt% is preferable.

Vanadium (V) is used to increase hardenability thus retarding the formation of high temperature transformation products on cooling to the quench temperature. Furthermore V may combine with C or N or both to form fine strengthening precipitates thus increasing strength and promoting uniformity in strength at the microstructural sca le i n mixed microstructures and in turn improved stretched edge ductility. V addition up to 0.4wt% is effective. Higher additions are undesirable for reasons of cost and because excessive levels of precipitation tie up high amounts of C. Since free C is required for austenite stabilisation too high V requires increased C addition. Preferably V is below 0.1%.

Phosphorus (P) is used to suppress the formation of pearlite d uring cooling, to suppress carbide formation and thereby promote the partitioning of carbon to austenite resulting in austenite stabilisation. However, too high a P addition is known to cause embrittlement at hot-rolling temperatures and to lead to reduced toughness in martensitic UHSS. P may also lead to problems in spot welding of the final product. For these reasons P is limited to a maximum of 0.08wt% and more preferably to a maximum of 0.02wt%.

Sulphur (S) is an impurity which may embrittle the intermediate or final product. S should be limited to a maximum level of 0.01wt% and more preferably to at most 0.005wt%.

Chromium (Cr) and nickel (Ni) may be used to increase hardenability thus retarding the formation of high temperature transformation products on cooling to the quench temperature. Additions of Cr and/or Ni may therefore allow the utilisation of lower cooling rates more readily achieved in full-scale production lines or to achieve the desired structure and properties in heavier gauge strip. Cr and/or Ni should be limited to a level of less than lwt% for reasons of cost and preferably to levels of 0.5wt% or less and more preferably below 0.1%.

Boron (B) may be used to improve hardenability and, in particular, to prevent the formation of ferrite on cooling from a fully austenitic soaking temperature. B should be limited to 50ppm because above these levels further addition is ineffective.

Antimony (Sb) may be used to enhance the wettability of zinc during hot dip galvanising. Addition should be limited to 0.06wt% or less.

Calcium (Ca) may be required to avoid clogging during casting and may be beneficial for modifying the morphology of MnS inclusion. Globular inclusions are known to be less detrimental to stretched edge ductility than highly elongated inclusions. If used, then addition should preferably be made to the level of 30ppm or less. Additions above 0.05% lead to the formation of coarse inclusions detrimental to ductility and toughness.

N itrogen (N) may be used, in combination with V, to form fine strengthening precipitates imparting strength and promoting uniformity in strength at the microstructural scale and in turn good stretched edge ductility. Additions should be limited to 150ppm or less.

The annealing step starts with reheating to an annealing temperature (Ti) . Ti may be above or below Ac₃. If Ti is below Ac₃ the resulting ferrite fraction leads to a lower strength and may introduce heterogeneity in strength at a microstructural scale. Local differences in strength lower the ratio of yield to tensile strength and lead to poor stretched edge ductility and bendability. It is therefore preferable that Ti is above Ac₃. In the context of this invention the phrase "above Ac₃" means that the microstructure is austenitic at TV Although the equilibrium transformation temperature Ae₃ is only determined by the composition, the value of the corresponding Ac₃ temperature is not a constant value as its value depends among others on the heating rate during which Ac₃ is measured and the starting microstructure of the steel. Usually Ac₃ is determined using dilatometry. When the heating rate used during dilatometry and the microstructure of the test specimen are those used in the process according to the invention, the value of Ac₃ is easy to determine. In determining Ti and the allowable variation around Ti the narrowing of the transformation temperature range with increasing C content should preferably be taken into account as follows:

(Ac₃-40) < Ti < (Ac₃+50) for carbon between 0.18 - 0.3% C or

(Ac₃-20) < Ti < (Ac₃+30) for carbon between 0.3 to 0.4% C

The rapid cooling to T₂ (CR1) is required to avoid the formation of high temperature transformation phases. The specific rate required depends upon the steel chemistry and corresponds to the critical cooling rate for avoidance of ferrite and pearlite noses in the relevant CCT diagram. In addition, the critical rate decreases with increasing Ti above Ac₃. Preferably CR1, i.e. the cooling rate over the temperature interval 800-500°C from Ti to T₂ is between 30 to 80 °C/s. T₂ should be chosen low enough to deliver partial transformation to martensite, but not so low as to cause complete transformation to martensite. T₂ is chosen to deliver a volume fraction of martensite of between 50 to 90% (in volume) and preferably an a ustenite fraction of at least vol.5%. The dependency of the martensitic transformation start temperature (M_s) on composition means that T₂ will also depend upon chemistry. Preferably a martensite fraction of 60 - 85 vol.% is chosen.

Holding at T₃ is needed to enrich the remaining austenite in carbon via a bainitic transformation or carbon partitioning or both. This temperature may be the same as the quench stop temperature (T₂=T₃) or may be higher (T₃>T₂). Higher isothermal holding temperatures may be advantageous since increased rates of carbon diffusion may make feasible shorter isothermal holds. Several processes occur d uring isothermal holding including tempering of martensite, diffusion of carbon from the martensite fraction to the rema ining a ustenite, the precipitation of ca rbides a nd the formation of ba inite. The p resence of a sma l l fra ctio n of ma rte nsite is known to a cce lerate su bseq uent transformation to bainite so, for some combinations of alloy composition and processing capability, quench stop temperatures under Ms are desirable. T₃ is chosen so as to give a suitable rate of transformation to bainite or rate of partitioning or both. The specific temperature will be dependent upon alloy composition and will preferably fall in a range expressed by: (Ms-70) < T₃ < (Ms+ 150), and preferably T₃ < (Ms+50)

The correct balance of isothermal holding temperature and isothermal holding time must be chosen for each composition. These can be determined by means of dilatometry as described hereinbelow.

The formation of carbides may lock up carbon which would otherwise be available for stabilisation of austenite and should therefore preferably be avoided. Furthermore, coarse Fe₃C carbides may lead to a deterioration in tensile ductility and/or stretched edge ductility. The levels of Si or Si/AI must be suitably tuned to retard carbide formation for the duration of the isothermal hold.

Since both the carbon partitioning accompanying the martensite tempering and bainitic transformation, in the presence of sufficient Si or Si/AI, each deliver an enrichment of carbon in austenite, exploitation of both mechanisms in tandem is highly beneficial leading to a greater total degree of austenite stabilisation.

It should also be noted that the optimum properties are not delivered by simply tuning process to deliver the maximum stabilisation of austenite. Levels of carbon enrichment too low to completely stabilise all austenite but sufficient to deliver high carbon austenite prior to the final cool can lead to the formation of brittle martensite. Excessive tempering of martensite can lead to significant loss in strength and, in combination with insufficient stabilisation of austenite and the formation of new hard martensite on final cooling, to low ratios of the yield to tensile strength.

Therefore, it is preferred that the partitioning temperature and time are chosen such as to optimise the enrichment of ca rbon in the austenite but without creation of deleterious microstructures during the isothermal hold.

Subsequent to the hold at T₃ the strip is cooled to ambient temperature. The strip may also be coated with zinc or other such metallic layers using a suita ble method of deposition either in-line or in a following process step.

In a preferable embodiment the cold rolled strip contains at least 0.25% C and at least 0.03% Al. A lower limit of 0.25wt% is placed on C because below this level the desired combinations of strength and ductility may not be achieved.

In an embodiment the use of Si is preferred to that of Al such that a minimum silicon content of lwt% and a max maximum aluminium content of 0.5 wt% is defined. Si provides substantial strengthening allowing the achievement of ultra high strength, more effectively su ppresses ca rbide formation ena bl ing longer isotherma l holds without formation of large volumes of coarse iron carbides, and because it does not accelerate bainite formation to the same extent as Al thus preventing excessive formation of bainite and enabling higher strengths to be achieved. Strength-Ductility data for a range of production C-Mn steels including ferritic forming steels and quenched martensitic steels have been used to generate a base-line strength-ductility decay for conventional strip steels. The data conform to the expression :

1

e calculated k (1)

where e_caicuiated is the total elongation (expressed as % engineering strain), UTS is the ultimate tensile strength and k is a constant which for tensile test pieces with 80 mm gauge and thickness 1 mm is 250000.

If elongations are measured at d ifferent ga uge or thickness then they must be converted to an equ iva lent elongation at 80mm ga uge a nd 1 mm thick or the a bove expression must be fit to base-line data measured at that alternative gauge/thickness combination using appropriate values of the constants. Conversion of tensile ductility can be performed using accepted procedures (ISO Norm 2566/1-2) when the geometries corresponding to the measured and to be calculated elongations are know :

Where e₂ is the required elongation for a gauge length of L₂ with a cross section of A₂, ei is the known elongation measured for a gauge length of U with a cross section of Ai and the exponent m is a material constant here assumed to be equal to 0.4.

In an embodiment the quenched and partitioned steel has an e-ratio of at least 1.8 wherein the e-ratio is defined as emeasured/e_Caicuiated a nd wherein e_caicuiated is calculated according to equation (1) and wherein e_measured is the elongation measured from an 80mm gauge length sample at 1mm thick (or measured at some other geometry and converted to an equivalent elongation on an 80mm gauge at 1mm thick using expression (2)).

In an embodiment the tensile strength of the steel according to the invention is at least 900 MPa. This strength regime is of interest since it provides significant opportunity for down-gauging and is a strength regime for which formability is most limited. Ductility levels a re at least 1.8x or more tha n that of conventiona l C-Mn steels at equivalent strength.

Preferably the yield to tensile strength ratio is 0.6 or higher. More preferably the ratio is at least 0.65 or even higher. Low yield to tensile strength ratios are associated with poor bendability and edge cracking sensitivity. Performance is often dependent on yield strength, anti-intrusion components for instance require high yield strength. High yield to tensile strength ratios ensure strength uniformity in the formed part, especially in forming operations which apply localised strain such as bending, or hole expansion.

In a preferred embodiment of the invention the metal or metal alloy coating is zinc, aluminium, magnesium or alloys thereof. Hereby the steel is afforded sacrificial corrosion protection since the zinc and aluminium will oxidise in preference to iron in the steel.

It was stated above that it is preferred that the partitioning temperature and time are chosen such as to optimise the enrichment of carbon in the austenite but without creation of deleterious microstructures during the isothermal hold. The temperature and time can be determined using dilatometry as follows:

First, fully austenitise a sample and apply a fast cool to room temperature to determine the fraction of martensite as a function of temperature below Ms. This allows to select an appropriate fraction of martensite as a function of the quench temperature. For this selected quench temperature the range of time-temperature partitioning combinations that deliver the correct stabilisation of austenite can be determined as follows. Cool a fully austenitized specimen to the selected quench temperature T₂ to produce the desired fraction of austenite and martensite and subsequently reheat the sample to a selected partitioning temperature T₃ and hold the sample at this temperature for a certain partitioning time t₂. The value of t₂ may be between 10 and 500 seconds, but for practical purposes in commercial annealing lines t₂ is preferably in the range of 20 to 180 s or even 20 to 100 s. After the partitioning step the specimen can be cooled naturally or acceleratedly to room temperature; no fast quench is needed. If during this cooling the dilatation curve shows that the martensite formation re-starts at a temperature in the range from (Ms - 20) to 120°C, then a correct degree of stabilisation has been achieved. In the case that the martensite formation recommences at a temperature < 120°C, it means that the stabilisation is too strong and the martensite has a very high carbon content compared to the bulk concentration. In the other case that the martensite formation recommences at a temperature less than 20°C below Ms the stabilisation is not sufficient. This experiments needs to be repeated for different partitioning times to find the times corresponding the two limiting boundary conditions for the chosen partitioning temperature.

By varying the partitioning temperature T₃ and repeat above the required partitioning time will a lso va ry. Increasing the partitioning temperature resu lts in a decrease in partitioning time t₂ and to a higher degree of tempering of the martensite formed during the quench.

In a preferred embodiment of the invention the metallic coating is provided by hot- dip galvanising or by electro-galvanising. BRIEF DESCRIPTION OF THE DRAWINGS

The invention will now be elucidated by way of example making reference to the accompanying figures and examples.

Figure 1 shows the schematic annealing schedule indicating the meaning of Ti, T₂ and T₃, ti and t₂, and of CR1 and CR2.

Figure 2a and 2b show a set of result of the dilatometric experiments to determine the quench temperature and the partitioning time. Figure 2a shows the temperature as a fu nctio n of ti me for a steel havi ng 3.5% M n w h ich was q uenched to a quench temperature of 280°C and reheated to a partitioning temperature of 330°C (triangle) and 440°C (circle). The sample was held at the partitioning temperature for 20 seconds. The quenching temperature resulting in the required amount of martensite is determined on the basis of the base curve (NC-III, sq ua re). Fig ure 2b shows the d ilatation of the samples for these conditions. The base curve, with a full quench to room temperature allows to determine Ms (about 315°C). The partitioning step for 20s at 330°C shows no dilation of the sample, which means that no bainite is formed, carbon is partitioned and the martensite is only marginally tempered. The transformation to martensite re-starts at temperatures lower tha n the quench temperature of 280°C, namely at 250°C, which ind icates that the austenite has been stabilized d ue to carbon partitioning. Increased partition times at 330°C show that the transformation to martensite re-starts at lower temperatures than 250°C. The dilation observed during annealing of the sample for 20s at 440°C mea ns that ba in ite is formed. Concurrently, carbon is partitioned and the martensite is tempered. The consequence of this high partitioning temperature is that the martensite is severely tempered . Due to the ba in itic tra nsformation a nd the ca rbon pa rtitioning the austenite is strong ly en riched in ca rbon, wh ich is reflected i n the relatively low start temperature of the martensite of approximately 150°C during the second quench after the partitioning at 440°C. The latter martensite is very ha rd and brittle and therefore undesirable. By va rying the partitioning time and the partitioning temperature the optimum combination can be determined.

Table 2 shows the results of various thermal cycles with the steels of Table 1. These results show that (I-VI) :

I. The desired property balance is not achieved for processing with C and Mn additions at the level of current commercia l TRIP Steels (composition G) . G delivers a very la rge extension of ductility (typically 2.2 times that of a conventional C-Mn steel) at strengths ra ng i ng fro m 850 to 1050 M Pa . H oweve r, th is is o n ly true w he n the a n nea ling temperature Ti is chosen below the preferred range, (Ac₃-40 to Ac₃+40 such that a high fraction of ferrite is retained in the final structure. From table 2 it is apparent that, in these cases (cycles 31-32, although ductility is at the desired level, the ratio of YS to UTS drops below the desired level to approximately 0.4.

When annealing is performed with Ti within the claimed range, both direct quenching to room temperature and direct quenching to room temperature followed by an isothermal hold at a higher temperature (i.e. conventional quenching and tempering as in cycle 25) deliver strengths in the desired range but do not deliver ductility above the desired minimum level.

When annealing is performed with Ti within the claimed range (cycles 26-30 and 33) strengths in the range 1000-1300 MPa may be achieved but the desired level of ductility is not achieved.

Processing variants annealed at Ti within the claimed range, subsequently quenched to a low T₂ to give high fractions of martensite (in examples 85-93%) after the quench and followed by an isotherm at some higher temperature (table 3 cycles 28-30) to not deliver the desired level of ductility.

Processing variants annealed at Ti within the claimed range, subsequently quenched to a low T₂ to give high fractions of martensite after the quench followed by an isotherm at the same temperature (cycle 26) do not deliver the desired level of ductility.

Processing variants annealed at Ti within the claimed range, subsequently quenched to a high T₂ to give low fractions of martensite (in example 50%) after the quench followed by an isotherm at the same temperature (cycle 27) to not deliver the desired level of ductility.

II. Complete replacement of Si with Al delivers strength and ductility below the desired minima : Composition A when directly quenched to room temperature delivers strength in the desired range but ductility below the desired minimum (cycle 4). Composition A when quenched to a T₂ above the M_s temperature and subjected to isothermal holding at the same temperature delivers strength greatly below the desired range and ductility below the desired minimum (cycle 3). Composition A when quenched to a T₂ below the M_s temperature and subjected to isothermal holding at some higher temperature delivers strengths in the approximate range 950 to lOOOMPa and ductilities below the desired minimum (thermal cycles 1-2).

III. Addition of C and/or Mn to levels beyond that found in current commercial trip steels does not enable the desired strength ductility balance to be achieved with conventional quench a nd temper process. For compositions B, C, D, E , F, a nd H utilising a direct quench or a direct quench followed by isothermal holding at some higher temperature (i.e. conventional quench and temper) delivers strengths in the desired range but ductility lower than the required minimum (cycles 9, 14, 25, 19, 24, 38). IV. Addition of C and/or Mn enables the desired strength ductility balance to be achieved when T₂ is above the preferred range but only when the duration of the isothermal hold is unacceptably long for conventional process. Compositions B, C, D, F and H each enable the desired property range to be achieved even if T₂ is set such that no martensite is formed during the initial quench (cycles 8, 12, 13, 16, 17, 36 and 37) but in all cases the isothermal holding time at T₃ is unacceptably long to be practical or economical in a continuous annealing process.

V. Addition of C and/or Mn enables the desired strength ductility balance to be achieved when processing is carried out in the preferred manner using a quench temperature (T₂) below Ms and a suitable combination of isothermal holding temperature (T₃) and time. Compositions B, D and F each enable the desired property range to be achieved if T₂ is set such that the desired fraction of martensite is formed during the initial quench (cycles 6, 15, 22) and if an isothermal holding temperature (T₃) and the holding time at this temperature are set at suitable levels.

VI. C addition of higher than 0.4wt% leads to embrittlement when processing is carried out according to the preferred route. Compositions C and H each return effectively zero ductility results when subjected to processing involving a deep quench (low T₂) a nd isothermal hold times at T₃ sufficiently short for conventional CA process (cycles 10-11 and 34-35) due to the formation of brittle, high carbon martensite in the final structure.

Table 1 - Alloy Compositions (Ca-treated).

(black cells not according to invention)

Table 2 - Processing parameters and mechanical properties

Comparative

Ac3 Ms Tl T2 T3 YS UTS YS/UTS e ratio

Cycle t2 (s) Comment

(°C) (°C) (°C) (°C) (°C) (MPa) (MPa) /

(-) (-) Inventive

A 1 1000 420 1080 240 400 60 766 952 0.80 0.6 C Ductility low

A 2 1000 420 1080 320 400 60 809 1004 0.81 0.6 C Ductility low

A 3 1000 420 1080 430 430 60 434 573 0.76 0.7 c Ductility & Strengt

A 4 1000 420 1080 20 NA NA 925 1335 0.69 1.1 c Ductility low

B 5 855 330 880 200 350 60 734 0.55 1.9 c YS/UTS Low

B 6 855 330 880 220 450 60 849 1137 0.75 1.9 I Inventive Exam

B 7 855 330 920 370 370 480 543 1018 0.53 2.2 c YS/UTS Low

B 8 855 300 820 370 370 600 588 926 0.63 2.1 c t₂ too long

B 9 855 330 920 20 450 300 1317 1363 0.97 1.1 c Ductility low

C 10 825 300 860 180 320 60 879 1344 0.65 0.2 c Ductility low

C 11 825 260 800 200 400 60 606 1098 0.55 0.3 c Ductility low

C 12 825 300 860 340 340 900 1012 1414 0.72 2.4 I t₂ too long

C 13 825 260 800 350 350 900 975 1305 0.75 3.5 I t₂ too long

C 14 825 300 860 20 NA NA 833 954 0.87 0.0 c Ductility Low

D 15 855 325 900 200 350 60 1196 1576 0.76 1.8 I Inventive Exam

Comparative

Ac3 Ms Tl T2 T3 YS UTS YS/UTS e ratio

Cycle t2 (s)

(°C) (°C) (°C) (°C) (°C) (MPa) (MPa) / Comment

(-) (-) Inventive

D 16 855 300 810 370 370 900 860 1273 0.68 2.5 I t₂ too long

D 17 855 325 880 360 360 900 770 1179 0.65 1.9 I t₂ too long

D 18 855 325 900 20 450 300 1306 1420 0.92 1.2 c Ductility low

E 19 895 370 930 20 450 300 1153 1193 0.97 1.0 c Ductility low

E 20 895 370 930 410 410 180 809 1169 0.69 1.5 c Ductility low

F 21 840 315 870 270 450 20 780 1426 0.55 1.4 c Ductility low

F 22 840 315 870 270 350 20 1057 1565 0.68 1.8 I Inventive Exam

F 23 840 315 870 360 360 900 790 1500 0.53 1.7 c Ductility low

F 24 840 315 870 20 NA NA 1220 1709 0.71 1.6 c Ductility low

G 25 895 380 900 20 350 60 987 1269 0.78 1.1 c Ductility low

G 26 895 380 900 250 250 60 934 1461 0.64 1.5 c Ductility low

G 27 895 380 900 350 350 60 929 1335 0.70 0.7 c Ductility low

G 28 895 380 900 240 350 60 1078 1298 0.83 1.0 c Ductility low

G 29 895 380 900 240 450 60 1063 1176 0.90 1.2 c Ductility low

G 30 895 380 900 280 400 60 1105 1246 0.89 1.0 c Ductility low

G 31 895 350 850 240 350 60 431 1020 0.42 2.2 c YS/UTS Low

G 32 895 350 850 280 400 60 421 876 0.48 2.2 c YS/UTS Low

G 33 895 380 900 420 420 60 754 986 0.76 1.4 c Ductility low

H 34 795 230 830 100 300 60 1095 1208 0.91 0.1 c Ductility low

H 35 795 230 780 120 300 60 763 854 0.89 0.0 c Ductility low

H 36 795 230 830 320 320 1800 1126 1524 0.74 2.9 I t₂ too long

H 37 795 230 780 320 330 1800 1145 1437 0.80 3.4 I t₂ too long

H 38 795 230 830 20 450 300 1591 1645 0.97 1.4 c Ductility low

(black cells not according to invention)

Claims

1. A method for producing a quenched and partitioned steel by providing a cold rolled and annealed steel strip containing (in weight %) :

. 0.18 - 0.4% C

• 1.5 - 4.0% Mn

. 0.5 - 2.0% Si

• 0 - 1.5% AI

• 0 - 0.5% Mo

• 0 - 0.5% Ti

• 0 - 0.4% V

• 0 - 0.010% Nb

• 0 - 0.005% B

• 0 - 0.015% N

• 0 - 0.08% P

. 0 - 0.01% S

• 0 - 0.06 % Sb

• 0 - 0.05 % Ca

• 0 - 1.0%Cr

• 0 - 1.0% Ni

• the remainder being iron and unavoidable impurities,

wherein the annealing process comprises the following steps:

(i) reheating the cold rolled strip to an annea ling temperature Ti of between Ac₃-40 and Ac₃+80;

(ii) hold ing the strip at Ti for a n a nnea li ng ti me ti of between 10 and 200 seconds;

(iv) repartitioning annealing the cooled strip at a temperature T₃ for enriching the austenite in carbon by repartitioning the carbon from the martensite to the austenite fraction for a repartitioning time t₂ of between 20 and 500 seconds wherein the martensite fraction is between 50 and 90% of the microstructure at the start of the repartitioning annealing;

(v) cooling the strip at a cooling rate CR2 to ambient temperature.

2. A method according to claim 1 wherein the martensite fraction is at least 60 and/or at most 85% of the microstructure at the start of the repartitioning annealing.

3. A method according to any one of the preceding claims wherein (Ac₃-40) < Ti < (Ac₃+50) for carbon between 0.18 - 0.3% C or

(AC3-2O) < Ti < (Ac₃+30) for carbon between 0.3 to 0.4% C

4. A method according to any one of the preceding claims wherein TI is above Ac3.

5. A method according to any one of the preceding claims wherein the cold rolled strip contains at most 0.5% Al.

6. A method according to any one of the preceding claims wherein the cold rolled strip contains at least 0.25% C and at least 0.01% Al.

7. A method according to any one of the preceding claims wherein T₂ equals T₃.

8. A method according to any one of claims 1 to 6 wherein T₂ is lower than T₃.

9. A method according to any one of the preceding claims wherein (Ms-70) < T₃ < (Ms+ 150), preferably wherein (Ms-70) < T₃ < (Ms+50)

10. A method according to any one of the preceding claims wherein the cold-rolled and annealed strip is coated with one or more metallic layer(s), preferably wherein the step of provid i ng the meta l l ic coating is by hot-dip galvanising or electro- galvanising.

11. Steel stri p prod uced by a ny one of the cla ims 1 to 10 wherein the steel has a microstructure containing at least 5% austenite an the e-ratio of at least 1.8 and a the tensile strength (UTS) of at least 900 MPa.

12. Steel strip according to claim 11 wherein the microstructure further contains one or more of bainite, martensite, tempered martensite, ferrite, fine carbides.

13. Steel strip according to claim 11 or 12 wherein the microstructure does not contain ferrite and/or coarse cementite.

14. Steel strip according to any one of claims 11 to 13 wherein the yield strength to tensile strength ratio is at least 0.6.