MEDIUM MANGANESE STEEL PRODUCT AND METHOD OF MANUFACTURING THE SAME

FIELD OF INVENTION

The present invention relates to medium manganese steels with excellent strength-ductility combination, and a method of manufacturing the same. In particular, the invention relates to medium manganese steels having a product of ultimate tensile strength (R _m ) and total elongation (A80), i.e. R _m x A80, up to 50 GPa %.

BACKGROUND

In view of the climate change due to greenhouse gas emission, the long-term strategy of automotive industry is to minimize vehicle weight without compromising crash performance, affordability and product reliability. Hot-dip galvanized (HDG) or galvannealed advanced high strength steel (AHSS) products with ultimate tensile strength extending from 500 to 1800 MPa are desired due to the advantages of lightweight and high strength-ductility combination which allow for considerable reduction in automobile body weight without compromising safety requirements.

The AHSS grades can be used to make safety relevant body in white (BIW) components such as side impact protection beams in which characteristics such as capacity for energy absorption and anti-intrusion are needed in the event of a crash. The ability to absorb energy subject to loading is an essential factor to improve the impact safety of the automotive vehicles. The product of ultimate tensile strength (R _m ) and total elongation (A), i.e., R _m ^x A, is an index of the ability to absorb impact energy. The higher the value of R _m ^x A, the better ability to absorb impact energy is.

The AHSS grades are divided into three generations according to the values of R _m ^x A.

The R _m ^x A is typically 10 - 25 GPa % for the first- generation AHSS including dual phase (DP) and transformation induced plasticity (TRIP) multiphase steels. The TRIP steels contain retained austenite that transforms into martensite during straining, giving the steel a combination of strength and ductility. Steels with high volume fraction of austenite have also improved wear and abrasion resistance. The R _m ^x A is typically above 50 GPa % for the second- generation AHSS which is mainly represented by the high manganese twinning induced plasticity (TWIP) steels. Manganese is well known for stabilizing austenite and is therefore alloyed to produce steels with metastable retained austenite. The TWIP steels with a fully austenitic microstructure comprise manganese in the range of 15 - 25 wt. %, which inevitably results in high alloying costs and a challenging production process.

The third generation AHSS is a group of medium manganese steels with Mn contents typically in the range of 3 - 10 wt. % filling the gap between the first and second generation AHSS and exhibits R _m ^x A in the range of 25 - 50 GPa %.

The transformation induced plasticity (TRIP) phenomenon is described as the mechanically induced martensitic transformation of metastable retained austenite. This phenomenon is important in applications where energy absorption is required, such as in vehicle crash situations. Upon straining the fraction of retained austenite in the microstructure activates the TRIP phenomenon during plastic deformation by transforming into martensite. When the thermal driving force, i.e. free energy for martensite transformation is insufficient, an externally applied force can complement and trigger the transformation. The TRIP phenomenon is activated even at high deformation rates as martensite formation occurs at a rapid rate. In the absence of a competing phase transformation, retained austenite will remain in the metastable state.

Previous studies show that a heat treatment called quenching and partitioning (Q&P) can be applied to medium manganese steels resulting in steels with high strength and ductility, properties that were for a long time thought to exclude one another. Partitioning refers to the movement of carbon from supersaturated martensite to untransformed austenite resulting in carbon depleted martensite and carbon-enriched austenite. The microstructure of Q&P steels is martensitic with a fraction of retained austenite.

In order to obtain a microstructure consisting of martensite and retained austenite by Q&P, i. the formation of ferrite and/or pearlite during quenching should be avoided; ii. the bainite formation should be inhibited or retarded so that the partitioning of carbon is maximized; iii. the carbide precipitation should be inhibited or retarded so that carbon is available for partitioning to austenite; iv. the carbon content of the steel should be optimized for thermal stabilization of desired fraction of retained austenite. However, conventional Q&P steels contain maximum 3 wt. % Mn which is less than medium manganese steels. Thus, medium manganese alloying involves technological and economical challenges.

The stability of austenite increases by i. increasing carbon and alloying element (e.g. Mn) content; ii. reducing the austenite grain size; iii. segregation of austenite stabilizing elements to austenite during intercritical annealing.

Austenite Reverted Transformation (ART) annealing is a two-step heat treatment for medium manganese steels. The ART annealing enables the production of steels with ultra- fine intercritical ferrite (IF), metastable retained austenite and precipitated phases. The ultra-fine intercritical ferrite (IF) is also known as high-temperature tempered martensite.

The first step of the two-step heat treatment - “austenitization” - includes complete or partial austenitization followed by quenching to room temperature, resulting in a martensitic microstructure. The second step - “intercritical annealing” - reverts martensite to austenite y), after which the steel is quenched and in some cases also tempered. The reversion y) in the two-phase region of the initial lath martensitic microstructure allows alloying elements, such as carbon and manganese, to redistribute between intercritical ferrite (IF) and reverted austenite. It was found that carbon and substantially manganese enrichment within the volume fraction of retained austenite contribute to the chemical stabilization of retained austenite.

SUMMARY OF INVENTION

The present invention is intended to develop a relatively low-alloy third generation AHSS product that has excellent combination of high strength and ductility. The present invention is further intended to develop a method of manufacturing the third generation AHSS product in a continuous industrial line.

In view of the state of art, the object of the present invention is to solve the problem of providing a medium manganese steel product with a high volume fraction of retained austenite and an improved value of R _m x A80. The problem is solved by the combination of specific alloy designs and optimized heat treatment parameters which enhances stabilization of retained austenite to room temperature. In a first aspect, the present invention provides a steel sheet product having a product of ultimate tensile strength (R _m ) and total elongation (A80), i.e. R _m x A £ 50 GPa %, preferably in the range of 10 GPa % to 50 GPa %, which steel sheet product comprises a composition consisting of, in terms of weight percentages (wt. %):

C 0.05 - 0.30

Mn 0.9 - 6.0, preferably 2.5 - 6.0

Si < 5, preferably £ 2, more preferably 0.1 - 1.5

Al £ 3, preferably 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1

Nb £ 0.1 , preferably £ 0.05

Ti £ 0.2, preferably £ 0.05

B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S < 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities.

The steel sheet product has a tensile strength up to 1350 MPa and a yield strength up to 90 % of the tensile strength. Preferably, the yield strength is lower than 90 % of the tensile strength.

The steel sheet product may comprise a zinc coating for corrosion protection.

Preferably, the total alloying content other than iron (Fe) is in the range of 4 wt. % to 8 wt. %.

The steel sheet product has a microstructure comprising a matrix consisting of, in terms of volume percentages (vol. %): martensite 45 - 95 retained austeinite 4 - 55 bainite £ 10

The 45 - 95 vol. % martensite consists of tempered martensite such as high-temperature tempered martensite, and untempered martensite. Preferably, the matrix comprises £ 40 vol. % untempered martensite. It is preferable that the steel sheet product has an average Mn content of £ 10 wt. %, and more preferably in the range of 5 wt. % to 10 wt. %, within the retained austenite microstructure.

The steel sheet product has a thickness of 50 mm or less, preferably 10 mm or less, and more preferably 6 mm or less.

In some embodiments, the steel sheet product is coated by hot-dip galvanizing or galvannealing.

In a second aspect, the present invention provides a method for manufacturing the steel sheet product comprising the following steps of

- providing a hot-rolled or cold-rolled steel sheet with a composition consisting of, in terms of weight percentages (wt. %):

C 0.05 - 0.30

Mn 0.9 - 6.0, preferably 2.5 - 6.0

Si < 5, preferably £ 2, more preferably 0.1 - 1.5

Al £ 3, preferably 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1

Nb £ 0.1 , preferably £ 0.05

Ti £ 0.2, preferably £ 0.05

B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S < 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities, and, which hot-rolled or cold-rolled steel sheet has an essentially martensitic microstructure;

- heating to an intercritical annealing temperature in the range of Ai to A ₃ ;

- annealing at the intercritical annealing temperature for a time period in the range of 30 seconds to 24 hours; and

- cooling to ambient temperature.

Essentially martensitic structure means here that more than 90%, preferably more than 95% and more preferably more than 99% of the microstructure comprises martensite. Therefore, the hot rolled or cold rolled steel sheet may have more than 90%, preferably more than 95% and more preferably more than 99% martensitic structure before intercritical annealing.

Preferably, the step of intercritical annealing is performed by batch annealing or continuous annealing.

In some embodiments, the step of cooling to ambient temperature is continuous cooling.

Optionally, the step of cooling to ambient temperature is followed by a heat treatment at a temperature in the range of 200 °C to 400 °C for a time period in the range of 30 seconds to 200 seconds, and preferably around 100 seconds.

Optionally, the manufacturing method comprises a further step of hot-dip galvanizing or hot- dip galvannealing.

In some embodiments, the step of intercritical annealing is performed by continuous annealing in combination with hot-dip galvanizing or galvannealing.

In some embodiments, the step of cooling to ambient temperature comprises the following stages of

- a first cooling to an annealing temperature in the range of 430 °C to 530 °C;

- annealing at the temperature in the range of 430 °C to 530 °C for a time period in the range of 30 seconds to 200 seconds; and

- a second cooling to ambient temperature.

Preferably, the time period of annealing at the temperature in the range of 430 °C to 530 °C is around 100 seconds.

In some embodiments, the hot-rolled steel sheet having an essentially martensitic microstructure is produced by the following steps of

- providing a steel slab with a composition consisting of, in terms of weight percentages (wt. %):

C 0.05 - 0.30

Mn 0.9 - 6.0, preferably 2.5 - 6.0

Si < 5, preferably £ 2, more preferably 0.1 - 1.5

Al £ 3, preferably 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1 Nb £ 0.1, preferably £ 0.05

Ti £ 0.2, preferably £ 0.05

B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S £ 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities, and ;

- heating to a temperature in the range of 1100 °C to 1250 °C, and soaking for a time period of at least 1 hour;

- hot rolling, wherein the finish rolling temperature is above Ar ₃ and below 1000 °C, preferably in the range of 800 °C to 1000 °C; and

- cooling to ambient temperature.

In some embodiments, the step of cooling to ambient temperature is performed by direct quenching to a temperature of 150 °C or below at a cooling rate of at least 10 °C/s.

In some embodiments, the hot-rolled and cold-rolled steel sheet having an essentially martensitic microstructure is manufactured by the following steps of

- providing a hot-rolled steel sheet;

- heating to a sub-critical annealing temperature below Aci and above AcHOO °C;

- annealing at the sub-critical annealing temperature for a time period of at least 4 hours;

- pickling and cold rolling;

- reheating to an austenitization temperature of A ₃ or above, and annealing for at least 30 seconds; and

- cooling to a temperature of 150 °C or below.

Preferably, the sub-critical annealing temperature is in the range of 650 °C to 670 °C.

In a third aspect, the present invention provides an intermediate hot-rolled steel sheet having an average hardness of 400HV1 or less that can be achieved after the step of sub- critical annealing and before the step of pickling and cold rolling.

Accordingly, a hot-rolled steel sheet having an average hardness of 400HV1 or less can be manufactured by a method comprising the following steps of

- providing a hot-rolled steel sheet which comprises a composition consisting of, in terms of weight percentages (wt. %): C 0.05 - 0.30 Mn 0.9 - 6.0, preferably 2.5 - 6.0

Si < 5, preferably £ 2, more preferably 0.1 - 1.5

Al £ 3, preferably 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1

Nb £ 0.1 , preferably £ 0.05

Ti £ 0.2, preferably £ 0.05

B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S < 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities, and ;

- heating to a sub-critical annealing temperature below Aci and above AcHOO °C;

- annealing at the sub-critical annealing temperature for a time period of at least 4 hours; and

- cooling to ambient temperature.

The hot-rolled steel sheet having an average hardness of 400HV1 can be used for manufacturing the hot-rolled and cold-rolled steel sheet having an essentially martensitic microstructure.

In a fourth aspect, the present invention provides use of the hot-rolled steel sheet having an average hardness of 400HV1 or less for manufacturing the steel sheet product according to any one of the claims 12 to 14.

The method may comprise to manufacturing a steel sheet product as disclosed herein.

BRIEF DESCRIPTION OF DRAWINGS

Fig. 1 is a graph showing a heat treatment cycle for austenite reverted transformation of the steel Examples 1 to 20.

Fig. 2 is a graph showing a heat treatment cycle for austenite reverted transformation of the steel Example 21. Fig. 3 is a graph showing the microstructures of the steel Example 21, at (a) magnification x5000 and (b) magnification x20000.

Fig. 4 is a graph showing the average hardness vs. time of annealing at a sub-critical temperature of 670°C for the steel Example 22.

Fig. 5 is a graph showing the microstructures of the steel Example 1, at magnification x20000.

DETAILED DESCRIPTION OF THE INVENTION

The term “steel” is defined as an iron alloy containing carbon (C).

When manufacturing a steel sheet product and the method comprises the step of cold rolling a sheet, the cold-rolling step is always preceded by a step of hot-rolling the sheet.

The symbol “Ai” is used to denote the eutectoid temperature at which austenite begins to form during heating or transformation of austenite to ferrite (plus cementite) is completed during cooling, where lower indices “c”, “r” and “e” denote to heating, cooling and equilibrium, respectively.

The symbol “A ₃ ” is used to denote the temperature at which transformation of ferrite to austenite is completed during heating or austenite begins to transform to ferrite during cooling, where lower indices “c”, “r” and “e” denote to heating, cooling and equilibrium, respectively.

The symbol “M _s ” is used to denote the temperature at which austenite begins to transform to martensite during cooling.

The symbol “ART” denotes austenite reverted transformation.

The symbol “HR” denotes hot rolling.

The symbol “CR” denotes cold rolling.

The term “heat-affected zone (HAZ)” refers to a non-melted area of a metal material that has experienced changes in its material properties as a result of exposure to high temperatures. The alterations in material properties are usually a result of welding or high- heat cutting procedures. The HAZ is identified as the area between the weld or cut and the base metal material. These areas can vary in size and severity depending on the properties of the materials involved, the intensity and concentration of heat, and the process employed.

The term “total elongation (A)” refers to the percentage by which the material can be stretched before it breaks; a rough indicator of formability, usually expressed as a percentage over a fixed gauge length of the measuring extensometer. Two common gauge lengths are 50 mm (A50) and 80 mm (A80).

The term “ultimate tensile strength (UTS, R _m )” refers to the limit, at which the steel fractures under tension, thus the maximum tensile stress.

The term “yield strength (YS, R _p 0.2)” refers to 0.2 % offset yield strength defined as the amount of stress that will result in a plastic strain of 0.2 %.

Alloying elements have various effects on the microstructure and mechanical properties of a steel product. Thus, alloy design is one of the first issues to be considered when developing a steel product with targeted mechanical properties. Alloying elements also affect the cost, processability, recyclability, weldability and coatability of steel, all of which are requirements set by the automotive industry.

Alloying elements can affect the transformation temperatures such as Ai, A ₃ and M _s .

Carbon (C), manganese (Mn) and nickel (Ni) decrease Aci while silicon (Si), chromium (Cr) and molybdenum (Mo) increase Aci. Chromium (Cr) and copper (Cu) decrease AC ₃ while aluminum (AI) and microalloying elements, such as niobium (Nb), titanium (Ti) and vanadium (V), increase AC ₃ . In commercial practice phase transformations occur usually at temperatures above the theoretical transformation temperatures determined in laboratory conditions.

Next, the chemical composition is described in more details, wherein % of each alloying element refers to weight percentage (wt. %). Preferably, the total alloying content other than iron (Fe) is in the range of 6 wt. % to 8 wt. %.

Carbon C is used in the range of 0.05 % to 0.30 %.

C alloying increases strength of steel by solid solution strengthening, and hence C content determines the strength level. C content less than 0.05 % may lead to insufficient strength. C depresses pearlitic transformation. C also functions as an austenite stabilizer and the austenite fraction increases by the diffusion of C into the austenite phase. C decreases A thereby delaying transformation of austenite, which inhibits the formation of ferritic-pearlitic microstructures.

However, C has detrimental effects on weldability, weld toughness and impact toughness of steel. Therefore, C content is set to not more than 0.30 %.

Manganese Mn is used in the range of 0.9 % to 6.0 %.

Mn is a substitutional element in the iron lattice and acts thus as a solid solution strengthener. There seems to be a rough relation between higher Mn and higher strength level. Mn increases hardenability and hardness of steel and decreases the brittleness by binding to sulfur and forming MnS. Mn is a weak carbide former but a strong sulphide and oxide former. Mn depresses pearlitic transformation.

Mn is alloyed to stabilize austenite to room temperature, i.e. the stability of austenite increases by the diffusion of Mn into the austenite phase. In addition, Mn decreases the M _s temperature thereby delaying transformation of austenite to form martensite.

Mn segregates during solidification and causes concentration gradients i.e. Mn-rich and Mn-poor regions in the microstructure. As Mn stabilizes austenite, a larger fraction of other phases than austenite is formed in the Mn-poor regions during quenching. Also, manganese lowers the chemical potential of carbon in austenite meaning that carbon diffuses from Mn-poor regions to Mn-rich regions. Consequently, austenite in Mn-rich regions receives more of stabilizing carbon during partitioning. If the carbon partitioning of austenite has been insufficient, fresh martensite forms in the Mn-rich regions during final quenching.

Alloying with Mn more than 6.0% unnecessarily increases the CE value thereby weakening the weldability. If the Mn content is too high, hardenability of the steel increases such that the heat-affect zone (HAZ) toughness is deteriorated. Large additions of Mn may also result in economic losses as manganese oxides diffuse to slag. In addition, Mn reduces wettability of molten Zn during galvanizing.

Preferably, Mn is used in the range of 2.5 % to 6.0 %.

Furthermore, it is preferable that the steel sheet product according to the present invention has a Mn content of £ 10 wt. %, and more preferably in the range of 5 wt. % to 10 wt. %, within the retained austenite microstructure.

Silicon Si is used in the range of 5 % or less. Si is effective as a deoxidizing or killing agent that can remove oxygen from the melt during a steelmaking process. Si alloying enhances strength by solid solution strengthening, and enhances hardness by increasing austenite hardenability. Also the presence of Si can stabilize retained austenite. Si is an effective cementite suppressor and does not form carbides. Si increases the A3 temperature of steel, meaning that the austenitization temperature increases. The presence of Si enhances stabilization of retained austenite by suppressing the formation of carbon-consuming cementite. However, austenite in Si- containing steels can decompose to carbide-free bainitic ferrite interwoven with retained austenite.

A Si content of higher than 5 % may pose problems during galvanizing operations, as Si may react with iron oxide and forms fayalite on the hot-rolled steel surface which is difficult to remove in a pickling treatment. The rolls pick up the fayalite causing indentations on the surface of the sheet metal. Furthermore, excessive content of Si unnecessarily increases carbon equivalent (CE) value thereby weakening the weldability. Silicon also promotes pearlitic transformation. In addition, Si reduces wettability of molten Zn during galvanizing.

Si is preferably used in the amount of 2 % or less, and more preferably in the range of 0.1 % to 1.5 %.

Aluminum Al is used in the range of 3 % or less.

Al is effective as a deoxidizing or killing agent that can remove oxygen from the melt during a steelmaking process. Al increases solid solution strength and hardness. Al also removes nitrogen N by forming stable AIN particles and provides grain refinement, which effects promote high toughness. Aluminum is preferentially oxidized internally, thereby improving the galvanizability of the steel. Al increases the A3 temperature of steel thereby increasing the austenitization temperature. Al is an effective cementite suppressor and does not form carbides. Al stabilizes ferrite and depresses pearlitic transformation.

Al inhibits the carbon-consuming cementite precipitation thereby promoting the carbon enrichment of austenite. The stability of austenite is determined by the balance of C and Mn partitioned to austenite and the stabilizing effect of C is more dominant. Thus, the presence of Al can stabilize retained austenite. Al alloying also allows wider processing windows for the annealing temperature for obtaining near maximum fractions of retained austenite.

However, excess Al may increase non-metallic inclusions thereby deteriorating cleanliness. High-aluminum additions in steel deteriorate lubrication quality and pose problems related to the function of mould flux during continuous casting of steel. Preferably, Al is used in the range of 0.01 % to 2.5 %.

Nickel Ni is used in the range of 2 % or less.

Ni is an alloying element that decreases Aci thereby delaying transformation of austenite, which inhibits the formation of ferritic-pearlitic microstructures. Ni improves austenite hardenability thereby increasing strength without any loss of toughness.

However nickel contents of above 2 % would increase alloying costs too much without significant technical improvement. Excess Ni may produce high viscosity iron oxide scales which deteriorate surface quality of the steel product. Higher Ni contents also have negative impacts on weldability due to increased CE value and cracking sensitivity coefficient.

Preferably, Ni is used in the range of £1.6%.

Elevated Mn levels may cause excessive surface oxidation, which may interfere with the hot-dip galvanizing and galvannealing processes. In some embodiments, it may be therefore beneficial to replace some of the Mn with Ni. The combined levels of Mn and Ni may be: Mn+Ni > 3.0, and preferably Mn+Ni > 3.5.

Chromium Cr is used in the range of 0.8 % or less.

Cr increases Aci, but decreases Ac3. Cr alloying enhances strength and hardness by increasing austenite hardenability. Cr forms mid-strength carbides and increases the strength of both the base steel and the HAZ.

However, if Cr is used in an amount above 0.8 % the HAZ toughness as well as field weldability may be adversely affected. Preferably, Cr is used in the range of £0.65%.

Vanadium V is used in the range of 0.3 % or less.

V has substantially the same but smaller effects as Nb. V is a strong carbide and nitride former, but V(C,N) can also form and its solubility in austenite is higher than that of Nb or Ti. Thus, V alloying has potential for precipitation strengthening, because large quantities of

V are dissolved and available for precipitation in ferrite. V alloying retards recrystallization of austenite, thereby impeding the grain growth, which achieves a fine-grained microstructure with enhanced strength. V is also applied as a pearlite suppressor. Vanadium precipitates V4C3 in the form of plates enhance the resistant to hydrogen embrittlement of steel by trapping hydrogen.

However, The precipitation strengthening of V is decreased by the formation of coarse mixed compounds of microalloying elements such as (Ti,V)N. The driving force for such precipitation increases with increasing nitrogen content. Furthermore, V is a strong carbide stabilizing agent whereby less carbon may be available for stabilizing austenite. In case V is used excessively the carbon content of steel may need to be increased in order to compensate the carbon depletion due to the formation of vanadium carbides. Excessive amounts of V also have negative effects on weldability and hardenability.

Therefore, the upper limit of V is set to 0.3 %. Preferably, V is used in the range of £01%.

Niobium Nb is used in the range of 0.1 % or less.

Nb forms carbides NbC and carbonitrides Nb(C,N). Nb is applied as a pearlite suppressor and a major grain refining element. Nb alloying retards recrystallization of austenite thereby impeding the grain growth. Nb alloying further stabilizes the formation of carbides thereby enabling the strengthening effects of precipitation and grain refinement.

However, only the formation of fine precipitates contribute to further grain refinement, from which precipitation strengthening may follow. Thus, large amounts of Nb do not necessarily increase the strength of the material if only higher amounts of large precipitates are formed. Excessive content of Nb is harmful for the HAZ toughness since Nb may promote the formation of coarse upper bainite structure by forming relatively unstable TiNbN or TiNb(C,N) precipitates. Furthermore, Nb is a strong carbide stabilizing agent whereby less carbon may be available for stabilizing austenite. In case Nb is used excessively the carbon content of steel may need to be increased in order to compensate the carbon depletion due to the formation of carbides.

Therefore, the upper limit of Nb is set to 0.1 %. Preferably, Nb is used in the range of £0.05%.

Titanium Ti is used in the range of 0.2 % or less.

Ti is applied as a grain refiner, precipitation strengthener and pearlite suppressor. Ti binds the harmful free N by forming stable TiN. TiN together with NbC can efficiently impede grain growth by retarding recrystallization of austenite, which achieves a fine-grained microstructure with enhanced strength. TiN precipitates can further prevent grain coarsening in the HAZ during welding thereby improving toughness. TiN formation also suppresses BN precipitation, thereby leaving B free to make its contribution to hardenability. For this purpose, the ratio of Ti/N is at least 3.4.

The grain coarsening resistance is similar for steels with only TiN and Ti-rich nitrides with vanadium and niobium. Therefore, small amounts (0.01 %) of Ti are usually added to vanadium containing commercial steels. However, if Ti content is too high, coarsening of TiN and precipitation hardening due to TiC develop and the low temperature toughness may be deteriorated. Therefore, it is necessary to restrict titanium so that it is less than 0.2%. Preferably, Ti is used in the range of £0.05%.

Boron B is used in the range of 0.005 % or less.

B is a well-established microalloying element to increase hardenability. Boron can be added to retard phosphorus segregation to grain boundaries thereby reducing embrittlement during welding in the HAZ. Effective B alloying requires the presence of Ti to prevent formation of BN. In the presence of B, Ti content can be lowered to be less than 0.02%, which is beneficial for toughness. However, hardenability and weldability deteriorate if the B content exceeds 0.005 %.

Preferably, B is used in the range of £0.002%.

Calcium Ca is used in a content of 0.006 % or less.

Ca is not used as alloying element due to its low solubility in steel and high vapor pressure. The optional Ca addition during a steelmaking process is for refining, deoxidation, desulphurization, and control of shape, size and distribution of oxide and sulphide inclusions.

Unavoidable impurities can be phosphor P, sulfur S and nitrogen N. Their content in terms of weight percentages (wt. %) is preferably defined as follows:

P 0 - 0.05, preferably 0-0.02,

S 0 - 0.01, preferably, 0-0.006,

N 0 - 0.01.

Phosphor P suppresses the formation of cementite and acts as a solid solution strengthener. However, phosphorus decreases the hardenability of steels by forming complex phosphides that affect the amount of available substitutional elements in solid solution. Also, phosphorous segregates to grain boundaries during welding in the HAZ and leads to embrittlement.

Other inevitable impurities may be hydrogen H, oxygen O and rare earth metals (REM) or the like. Their contents are limited in order to ensure excellent mechanical properties, such as impact toughness.

The steel sheet product according to the present invention has a product of ultimate tensile strength (R _m ) and total elongation (A), i.e. R _m x A £ 50 GPa %, preferably in the range of 10 GPa % to 50 GPa %, The steel sheet product according to the present invention has a tensile strength up to 1350 MPa and a yield strength up to 90 % of the tensile strength. Preferably, the yield strength is lower than 90 % of the tensile strength.

The steel sheet product according to the present invention is manufactured in a process that determines a specific microstructure which in turn dictates the above-mentioned mechanical properties of the steel sheet product.

The steel sheet product according to the present invention has a microstructure comprising a matrix consisting of, in terms of volume percentages (vol. %): martensite 45 - 95 retained austeinite 4 - 55 bainite £ 10

The 45 - 95 vol. % martensite consists of tempered martensite such as high-temperature tempered martensite, and untempered martensite. Preferably, the matrix comprises £ 40 vol. % untempered martensite.

The steel sheet product can be provided in a wide thickness range. Preferably, the steel sheet product has a thickness of 50 mm or less, more preferably 10 mm or less, and even more preferably 6 mm or less.

In some embodiments, the steel sheet product is coated by hot-dip galvanizing or galvannealing.

The method according to the present invention for manufacturing the steel sheet product comprises the following steps of

- providing a hot-rolled or cold-rolled steel sheet with a composition consisting of, in terms of weight percentages (wt. %):

C 0.05 - 0.30

Mn 0.9 - 6.0, preferably 2.5 - 6.0

Si < 5, preferably £ 2, more preferably 0.1 - 1.5

Al £ 3, preferably 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1

Nb £ 0.1 , preferably £ 0.05

Ti £ 0.2, preferably £ 0.05 B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S £ 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities, and, which hot-rolled or cold-rolled steel sheet has an essentially martensitic microstructure;

- heating to an intercritical annealing temperature in the range of Ai to A ₃ ;

- annealing at the intercritical annealing temperature for a time period in the range of 30 seconds to 24 hours; and

- cooling to ambient temperature.

The step of intercritical annealing is effectively an ART heat treatment resulting in the final microstructure with a matrix comprising mainly ultra-fine intercritical ferrite (IF), a.k.a. high- temperature tempered martensite.

Previous studies suggested that blocky austenite nucleates at prior austenite grain (PAG) boundaries whereas acicular austenite forms between martensite laths during the intercritical annealing. The acicular austenite is more easily enriched with the austenite stabilizing carbon than the blocky austenite. A martensitic starting microstructure is considered ideal for the step of intercritical annealing in the aforementioned method according to the present invention, because the martensitic starting microstructure prior to the step of intercritical annealing enhances the formation of acicular austenite. The acicular austenite is easily enriched with austenite stabilizing carbon and a high volume fraction of metastable retained austenite thereby ensues in the final microstructure.

The martensitic starting microstructure before the intercritical annealing is further beneficial for producing the steel sheet product with a higher value of R _m x A80.

The starting material for the intercritical annealing may be either a hot rolled steel sheet or a cold rolled steel sheet, for example.

In an embodiment, the starting material for the intercritical annealing is a hot rolled steel sheet having an essentially martensitic microstructure. This embodiment is also further illustrated in Fig. 6, which incorporates an example of the manufacturing process for the hot rolled steel sheet and its intercritical annealing. The hot rolled steel sheet may be produced by the following steps of: - (A) providing a steel slab with a composition consisting of, in terms of weight percentages (wt. %):

C 0.05 - 0.30

Mn 0.9 - 6.0, preferably 2.5 - 6.0

Si < 5, preferably £ 2, more preferably 0.1 - 1.5

Al £ 3, preferably 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1

Nb £ 0.1 , preferably £ 0.05

Ti £ 0.2, preferably £ 0.05

B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S < 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities, and;

- (B) heating to a temperature in the range of 1100 °C to 1250 °C, and soaking for a time period of at least 1 hour;

- (C) hot rolling, wherein the finish rolling temperature is above Ar3 and below 1000 °C, preferably in the range of 800 °C to 1000 °C; and

- (D) cooling to ambient temperature.

After step D, the hot rolled steel sheet has an essentially martensitic microstructure. In order to manufacture the final product, the hot rolled steel sheet is then subjected to a step of heating to an intercritical annealing temperature in the range of A1 to A3 (E) and annealing at the intercritical annealing temperature for a time period in the range of 30 seconds to 24 hours (F). Finally, the hot rolled and intercritically annealed steel sheet is cooled to ambient temperature (G) to obtain the final product.

The step of soaking (B) is performed at a temperature in the range of 1100 °C to 1250 °C to enable homogenization while excessive grain growth is prevented.

For this embodiment, the cooling to ambient temperature after hot rolling may be performed either as air cooling or rapid cooling such as direct quenching. The cooling rate depends for example on the thickness of the steel sheet and chemical composition of the steel. In some embodiments, the step of cooling to ambient temperature is performed by direct quenching to a temperature of 150 °C or below at a cooling rate of at least 10 °C/s, and preferably 25-450 °C/s.

Due to the rather high Mn alloying, the hot-rolled steel sheet may be air hardenable, which means that almost fully martensitic microstructure can be achieved for the intercritical annealing, even without rapid cooling.

In another embodiment, the starting material for the intercritical annealing is a cold rolled steel with an essentially martensitic microstructure, which is manufactured by the following steps of

- providing a hot-rolled steel sheet;

- heating to a sub-critical annealing temperature below Aci and above AcHOO °C;

- annealing at the sub-critical annealing temperature for a time period of at least 4 hours;

- pickling and cold rolling;

- reheating to an austenitization temperature of A3 or above, and annealing for at least 30 seconds; and

- cooling to a temperature of 150 °C or below.

The hot rolled steel sheet in the first step of this embodiment may be manufactured by the process of providing the hot-rolled steel sheet having an essentially martensitic microstructure comprises the steps of

- providing a steel slab with a composition having a composition as the hot-rolled or cold- rolled steel sheet of disclosed herein;heating to a temperature in the range of 1100 °C to 1250 °C, and soaking for a time period of at least 1 hour;

- hot rolling, wherein the finish rolling temperature is above Ar3 and below 1000 °C, preferably in the range of 800 °C to 1000 °C; and

- cooling to ambient temperature, for example.

The hot-rolled steel sheet with high Mn content usually has hardness above 400HV1, which inevitably lowers its cold reliability. In order to increase cold reliability and acquire an average hardness of 400HV1 or less, the hot-rolled steel sheet may be subjected to a further step of sub-critical annealing before cold rolling.

An intermediate hot-rolled steel sheet having an average hardness of 400HV1 or less can be achieved after the step of sub-critical annealing and before the step of pickling and cold rolling. The step of sub-critical annealing is effectively a recrystallization heat treatment enabling nucleation and growth of new grains without phase change. The step of sub- critical annealing significantly decreases the hot-rolled steel sheet’s hardness to below 400HV1 , preferably 200HV1 to 350HV1.

Preferably, the sub-critical annealing temperature is in the range of 650 °C to 670 °C.

Accordingly, a hot-rolled steel sheet having an average hardness of 400HV1 or less, preferably 200HV1 to 350 HV1, can be manufactured by a method comprising the following steps of

- providing a hot-rolled steel sheet which comprises a composition consisting of, in terms of weight percentages (wt. %):

C 0.05 - 0.30

Mn 0.9 - 6.0, preferably 2.5 - 6.0

Si < 5, preferably £ 2, more preferably 0.1 - 1.5

Al £ 3, preferably 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1

Nb £ 0.1 , preferably £ 0.05

Ti £ 0.2, preferably £ 0.05

B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S < 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities, and;

- heating to a sub-critical annealing temperature below Aci and above AcHOO °C;

- annealing at the sub-critical annealing temperature for a time period of at least 4 hours; and

- cooling to ambient temperature.

The hot-rolled steel sheet having an average hardness of 400HV1 or less can be used for manufacturing the cold-rolled steel sheet having an essentially martensitic microstructure.

Preferably, the step of intercritical annealing of the hot rolled or cold rolled steel sheet with an essentially martensitic microstructure, is performed by batch annealing or continuous annealing. In some embodiments, the step of cooling to ambient temperature is continuous cooling.

Optionally, when continuous annealing is applied the step of cooling to ambient temperature may be followed by a heat treatment at a temperature in the range of 200 °C to 400 °C for a time period in the range of 30 seconds to 200 seconds, and preferably around 100 seconds.

Bainitic and/or martensitic transformations may occur to some degree during the step of cooling to ambient temperature, which result in volume fractions of bainite and fresh untempered martensite in the final microstructure. The heat treatment at a temperature in the range of 200 °C to 400 °C facilitates formation of the preferred tempered martensite.

Optionally, the manufacturing method comprises a further separate step of hot-dip galvanizing or galvannealing.

In some embodiments, the step of intercritical annealing is performed by continuous annealing in combination with hot-dip galvanizing or galvannealing thereby obtaining a steel sheet product with a zinc coating.

A continuous annealing line integrates several process steps including surface cleaning, annealing, hot-dip galvanizing or galvannealing, cooling, temper rolling, and surface protection, thus saving processing time, space, and costs. The continuous annealing line facilitates the production of high-strength automotive steel sheets with very uniform product properties.

In some embodiments, the step of cooling to ambient temperature after intercritcal annealing step in a continuous annealing line comprises the following stages of

- a first cooling to an annealing temperature in the range of 430 °C to 530 °C;

- annealing at the temperature in the range of 430 °C to 530 °C for a time period in the range of 30 seconds to 200 seconds; and

- a second cooling to ambient temperature.

The intercritical annealing step may be performed in a continuous annealing line.

Preferably, the time period of annealing at the temperature in the range of 430 °C to 530 °C is around 100 seconds.

The annealing at the temperature in the range of 430 °C to 530 °C is effectively a partitioning annealing which enables stabilization of reverted austenite to room temperature by enriching carbon within the reverted austenite. The present invention may provide a method for manufacturing a steel sheet product as disclosed herein. The steel sheet product, comprising a composition consisting of, in terms of weight percentages (wt. %):

C 0.05 - 0.30

Mn 2.5 - 6.0

Si 0.1 - 1.5

Al 0.01 - 2.50

Ni £ 2, preferably £ 1.6

Cr £ 0.8, preferably £ 0.65

V £ 0.3, preferably £ 0.1

Nb £ 0.1 , preferably £ 0.05

Ti £ 0.2, preferably £ 0.05

B £ 0.005, preferably £ 0.002

Ca £ 0.006

N £ 0.01

P £ 0.05, preferably £ 0.02

S < 0.01 , preferably £ 0.006 remainder Fe and inevitable impurities, and the steel sheet has a microstructure comprising a matrix consisting of, in terms of volume percentages (vol. %): martensite 45 - 95 retained austenite 4 - 55 bainite £ 10; and wherein the matrix comprises £ 40 vol. % untempered martensite; and the steel sheet product has an average Mn content in the range of 5 wt. % to 10 wt. %, within the retained austenite microstructure; and the steel sheet has a product of ultimate tensile strength (R _m ) and total elongation (A80), i.e. R _m x A80 £ 50 GPa %, and preferably in the range of 10 GPa % to 50 GPa %.

The method for manufacturing the steel sheet product comprises the method as described and disclosed herein.

Further, in the method for providing the steel sheet product, the steel sheet product comprises the definitions and disclosures of the steel sheet product as disclosed herein. The following Examples further describe and demonstrate embodiments within the scope of the present invention. The Examples are given solely for the purpose of illustration and are not to be construed as limitations of the present invention, as many variations thereof are possible without departing from the scope of the invention.

Table 1 shows the chemical compositions used for producing the tested steel Examples. Table 2 shows the results from microstructure analysis of the tested steel Examples.

Table 3 shows the mechanical properties of the tested steel Examples.

Table 4 shows the Vickers hardness values of the steel Examples 22 to 27

EXAMPLES 1 to 20

The steel Examples 1-20 are prepared by a process comprising the steps of:

- providing a hot-rolled (Examples 14 to 20), or cold-rolled steel sheet (Examples 1 to 13) with the Composition A, C, D, E, G or H (Table 1);

- austenitization by continuous annealing at a temperature above A ₃ ;

- quenching to obtain a steel sheet with mainly martensitic microstructure;

- heat treatment cycle (Fig. 1) as follows

(1) heating to a intercritical temperature in the range of Ai to A ₃ ;

(2) continuous annealing or batch annealing at the intercritical temperature for 30 seconds to 24 hours;

(3) cooling to ambient temperature.

The heat treatment cycle (Fig. 1) is conducted in a laboratory scale with Thermomechanical Simulator Gleeble model 3800-GTC. In these examples, the hot rolled steel was also austenitized and quenched before intercritical annealing to ensure essentially martensitic microstructure before the intercritical annealing.

Microstructures

Microstructure can be characterized from SEM micrographs and the volume fraction can be determined using point counting or image analysis method.

X-ray powder diffraction (XRD) is applied to analyze the microstructure of the steel Examples 2 to 17 which comprise retained austenite in an amount of 4 vol. % to 44 vol. % (Table 2). Electron backscatter diffraction (EBSD) is applied to analyze the microstructure of the steel Example 1 which comprises a matrix consisting of retained austenite in an amount of 16 vol. % and high-temperature tempered martensite, a.k.a. intercritical ferrite (IF), in an amount of 84 vol. % (Table 2). The amount of fresh untempered martensite and bainite is negligible.

Energy Dispersive X-ray Spectrometry (EDS) is applied to analyze the content of Mn within the retained austenite microstructure and the matrix. The steel Examples 16 and 17 comprise up to 7.5 wt. % Mn within the retained austenite microstructure, which is far above the mean Mn content of 4.5 wt. % in the matrix.

Microstructure for Example 1

Microstructure can be characterized from SEM micrographs and the volume fraction can be determined using point counting or image analysis method.

Electron backscatter diffraction (EBSD) is applied to analyze the microstructure of the steel Example 1 which comprises a matrix consisting of retained austenite in an amount of 16 vol. %; high-temperature tempered martensite, a.k.a. intercritical ferrite (IF), in an amount of 84 vol. %; (Table 2 and Fig. 5). The amount of bainite and fresh martensite is negligible. In Fig. 5, “M/A” denotes martensite/austenite islands, “IF” denotes intercritical ferrite a.k.a high-temperature tempered martensite, and “C” denotes carbide.

Yield strength

Yield strength (R 0.2) was determined according EN ISO 6892-1 standard using sub-sized tensile specimens. The steel Examples 1 to 13 and 18 to 20 each has an yield strength in the range of 383 MPa to 826 MPa (Table 3).

Ultimate tensile strength

Ultimate tensile strength (R _m ) was determined according to EN ISO 6892-1 standard using sub-sized tensile specimens. The steel Examples 1 to 13 and 18 to 20 each has an ultimate tensile strength in the range of 693 MPa to 1289 MPa (Table 3).

Total elongation

Total elongation (A80) was determined according to EN ISO 6892-1 standard using sub sized tensile specimens. The steel Examples 1 to 13 and 18 to 20 each has a total elongation value in the range of 11 % to 35 % (Table 3). Rm X A80

The product of ultimate tensile strength (R _m ) and total elongation (A80) for the steel Examples 1 to 13 and 18 to 20 is in the range of 12 GPa % to 38 GPa % (Table 3).

EXAMPLE 21

The steel Example 21 is prepared by a process comprising the sequential steps of subjecting a steel slab with the Composition B to

- hot rolling;

- pickling and cold rolling;

- austenitization by continuous annealing at a temperature above A3;

- quenching to obtain a steel sheet with mainly martensitic microstructure;

- heat treatment cycle (Fig. 2) as follows

(1) heating to an intercritical temperature of 770 °C;

(2) continuous annealing at 770 °C;

(3) first cooling to 470 °C;

(4) annealing and hot-dip galvanizing at 470 °C for 100 seconds;

(5) second cooling to ambient temperature.

In this example, the hot rolled steel sheet was not annealed at the sub-critical annealing temperature before pickling and cold rolling. The size of the steel Example 21 is 10 mm ^c 60 mm. The heat treatment cycle (Fig. 2) was conducted in a laboratory scale with Thermomechanical Simulator Gleeble model 3800-GTC. The Ai, A3 and M _s temperatures are estimated with dilatometry tests to 700 °C, 890 °C and 320 °C, respectively.

Microstructures

Microstructure can be characterized from SEM micrographs and the volume fraction can be determined using point counting or image analysis method.

Electron backscatter diffraction (EBSD) is applied to analyze the microstructure of the steel Example 21 which comprises a matrix consisting of retained austenite in an amount of 4 vol. %; high-temperature tempered martensite, a.k.a. intercritical ferrite (IF), in an amount of 56 vol. %; and fresh untempered martensite in an amount of up to 40 vol. % (Table 2 and Fig. 3). The amount of bainite is negligible. In Fig. 3, “M/A” denotes martensite/austenite islands, “IF” denotes intercritical ferrite, and “C” denotes carbide.

Yield strength

Yield strength (R _p 0.2) was determined according EN ISO 6892-1 standard using sub-sized tensile specimens. The steel Example 21 has an yield strength of 460 MPa (Table 3). Ultimate tensile strength

Ultimate tensile strength (R _m ) was determined according to EN ISO 6892-1 standard using sub-sized tensile specimens. The steel Example 21 has an ultimate tensile strength of 1100 MPa (Table 3).

Total elongation

Total elongation (A80) was determined according to EN ISO 6892-1 standard using sub sized tensile specimens. The steel Example 21 has a total elongation value of 16 % (Table 3).

Rm x A80

The product of ultimate tensile strength (R _m ) and total elongation (A80) for the steel Example 21 is in 18 GPa % (Table 3).

EXAMPLES 22 to 27

The steel Examples 22 to 27 is prepared by a process comprising the steps of

- providing a hot-rolled steel sheet with the Composition F;

- austenization at 1000 °C for 5 minutes and cooling to room temperature;

- reheating to a sub-critical temperature in the range of around 670 °C to 690 °C (Examples 22 to 26), or to a temperature 750 °C that is above Aci (Example 27);

- annealing at the sub-critical temperature for 4 hours, 8 hours, 24 hours or 1 week (Examples 22 to 26), or annealing at 750 °C (Example 27); and

- cooling to ambient temperature.

Vickers hardness (HV1)

The Vickers hardness test is performed by indenting the test material with a diamond indenter, in the form of a right pyramid with a square base and an angle of 136° between opposite faces subjected to a load of 1 kilograms-force (kgf) for 10 to 15 seconds. The two diagonals of the indentation left in the surface of the material after removal of the load are measured using a microscope and their average calculated. The area of the sloping surface of the indentation is calculated. The Vickers hardness is the quotient obtained by dividing the 1 kgf load by the mm ² area of indentation. The Vickers hardness value is expressed as xxxHVyy, e.g. 400HV1, where 400 is the Vickers hardness number, HV is the Vickers hardness scale, 1 indicates the load used in kgf.

Vickers hardness of the tested steel Examples 22 to 27 is dependent on the annealing temperature as well as the annealing time. In the case of Example 27 when the annealing temperature is above Aci, e.g. 750 °C, the Vickers hardness remains at an unacceptably high level after annealing (Table 4). Fig. 4 shows that when annealing is performed at the same sub-critical temperature of 670 °C the Vickers hardness decreases with increasing annealing time. The inventive Examples 23, 24 and 25 subjected to the annealing temperature of 670 °C for 8 hours, 24 hours and 168 hours respectively has an average Vickers hardness in the range of 233HV1 to 293 HV1 (Table 4) which is significantly less than the comparative Example 22, 26 or 27.

The comparative Example 22 subjected to the annealing temperature of 670 °C for 4 hours has an average Vickers hardness of 410HV1 (Table 4).

The comparative Example 26 subjected to the annealing temperature of 690 °C for 24 hours has an average Vickers hardness of 411 HV1 (Table 4).

The comparative Example 27 subjected to the annealing temperature of 750 °C has an average Vickers hardness of 545HV1 (Table 4).

Table 1. Chemical compositions (wt. %)

Table 2. Analysis of microstructure of the steel Examples 1 to 17 and 21

Table 3. Mechanical properties of the steel Examples 1 to 13 and 18 to 21

Table 4. Vickers hardness of the steel Examples 22 to 27