Field of the invention

This invention relates to a cold-rolled martensite steel with high strength and high bendability useful for components for vehicles and automobiles, and method of production thereof.

Background of the invention

In recent years, (advanced) high strength steel sheets, AHSS, are increasingly used in car components to reduce weight and fuel consumption. A series of (advanced) high strength steels, such as HSLA, Dual phase (DP), Ferritic-bainitic (FB) including stretch- flangeable (SF), Complex phase (CP), Transformation-induced plasticity (TRIP), Twinning- induced plasticity (TWIP) steels has been developed to meet the growing requirements.

However, AHSS sheet steels cannot be applied easily to a wide variety of car components because their formability is relatively poor. As steels became increasingly stronger, they simultaneously became increasingly difficult to form into automotive parts. Actually, the real application of AHSS steels (DP, CP and TRIP) to car components is still limited by their formability and poor crash performance. Therefore, improving formability, manufacturability and crashworthiness becomes an important issue for AHSS application.

High strength cold-rolled martensitic steels with tensile strength ranging from 900 to 1500 MPa have been developed over the last two decades. Steel sheets used for making reinforcing parts for automotive bumpers and door beams require good bending ability and good weldability in addition to high strength. Martensitic steels are increasingly being used in applications that require high energy absorption capacity for side impact and roll over vehicle protection, and have long been used for applications such as bumpers that can readily be rolled formed . One particular example of these steels is the steel referred to as 22MnB5, i.e. 0.22% of C, 1.2% of Mn, maximum 50 ppm of B, specified in EN10083. Hot press forming of 22MnB5 steel can produce complex parts such as bumpers and pillars with ultrahigh strength, minimum springback, and reduced sheet thickness. The tensile strength of boron steels is up to 1600 MPa, which is far above that of the highest-strength conventional cold stamping steels. However the bendability is limited which reduces its applicability.

DE102016013466-A1 discloses an alloy with, apart from unavoidable impurities due to manufacturing, the following composition (in weight %) comprising : C < 0.12%, Si < 0.6%, Mn < 1.8%, Cr < 0.5%, Nb < 0.09%, Ti < 0.05%, P < 0.03%, S < 0.01%, Al 0.015 - 0.1%, B < 0.005%, remainder iron. The properties of the finished base sheet after press hardening are preferably: yield strength between 800 M Pa and 1000 MPa, ultimate tensile strength between 1000 MPa and 1200 MPa, total elongation (gauge length 30 mm) > 6%, the bending angle for 1 mm strip > 90° .

EP1865086 discloses a steel composition comprising (in wt.%) 0.1-0.2% C, 0.05- 0.3% Si, 0.8-1.8% Mn, 0.5-1.8% Ni, <0.015% P, <0.003% S, 0.0002-0.008% B, optionally 0.01-0.1% Ti, optionally 0.01-0.05% Al, optionally 0.002-0.005% N. This composition makes it possible to manufacture a press hardened part with a tensile strength higher than 1000 MPa and with elongation higher than 10%. However, due to its high nickel content, this steel is costly to manufacture.

US6296805 discloses a steel composition comprising 0.20%-0.5% C, 0.8%-1.5% Mn, 0.1%-0.50% Si, 0.01%-1% Cr, <0.1% Ti, <0.1% Al, <0.05% P, <0.03% S, 0.0005%- 0.01% B, the remainder being iron and impurities inherent in processing . For galvanizing purposes, steels containing high silicon can greatly affect the surface quality of the steel and consequently, the thickness and appearance of the galvanized coating .

The energy absorption capacity of car components depends on the balance between strength and crash ductility, which is in general characterized by bendability. However, as the strength is increased, the ductility and the bendability decreases. A lower strength steel has good tensile elongation and bendability, but it cannot absorb higher energy due to the lower strength level . A higher strength steel has a decreased bending angle and the energy absorption capability also decreases. The 22MnB5 steel has a high tensile strength of 1600 MPa but a low bending angle of about 60° so the crash performance is not optimal .

Alternative high strength DP steel microstructures consist of ferrite, martensite and some bainite. The presence of ferrite in a hard matrix causes a poor bendability and also lower the yield strength. Both (lower bendability and lower yield strength) reduce the crash performance. Furthermore, in many cases the currently used DP steel is a cold formable steel . Controlling the phase fractions within the temperature accuracy of a reheating furnace and a hot stamping device followed by the accelerated cooling is difficult. This is because a DP microstructure by hot stamping may be created by reheating the blanks in the intercritical temperature range, where small temperature variations can give high scatter in final phase fractions causing unstable mechanical properties of the manufactured article.

Objectives of the invention

It is an object of the invention to provide a steel grade which combines high yield and ultimate tensile strength with high bendability and high crashworthiness.

It is also an object of the invention to provide a steel grade with a yield strength (YS) of at least 880 MPa, an ultimate tensile strength (UTS) of at least 1180 MPa, and an average bending angle (BA-L+BA-T)/2 of at least 90°. It is also an object of the invention to provide a steel grade with a yield strength (YS) of at least 880 MPa, an ultimate tensile strength (UTS) of at least 1180 MPa with a good coating quality and welding quality.

Description of the invention

One or more of the objects is reached with a cold-rolled and heat treated martensitic steel having high strength and high bendability having a microstructure consisting of 90 to 100 volume% (vol .%) of martensite and 0 to 10 vol.% in total of å(bainite + retained austenite) having the following composition:

• 0.13 - 0.25 wt.% C;

• 0.40 - 1.20 wt.% Mn;

• 0.04 - 1.00 wt.% Cr;

• 5 - 50 ppm by weight B;

• 0.020 - 0.20 wt.% AI_tot;

• 5 - 80 ppm by weight N;

• 0.001 to 0.040 wt.% Ti;

• 0 - 0.003 wt.% Ca ;

• 0 - 0.030 wt.% P;

• 0 - 0.010 wt.% S;

• 0 - 0.080 wt.% Si;

• 0 - 0.10 wt.% Nb;

• 0 - 0.10 wt.% Cu;

• 0 - 0.20 wt.% Mo;

• 0 - 0.060 wt.% Ni;

• 0 - 0.10 wt.% V;

• å(Cu+V+Ni+Mo) < 0.50 wt.%;

• CE = C + Mn/6 + (Cr + Mo + V)/5 + (Ni + Cu + Si)/15 < 0.46;

• remainder iron and residual elements,

the steel having a yield strength of at least 880 MPa and an ultimate tensile strength of at least 1180 MPa as determined on JIS5 tensile specimens according to NEN- EN 10002 : 2001 and an average bending angle (BA-L+BA-T)/2 of at least 90°, wherein BA-L and BA-T are the bending angles with the bending axis parallel (L) to or perpendicular (T) to the rolling direction of the cold-rolled steel as determined in a three-point bending test according to the VDA 238-100 standard.

Preferably the steel according to the invention has a yield strength of at least 900 MPa and an ultimate tensile strength of at least 1250 M Pa. A suitable maximum ultimate tensile strength is 1450 MPa . A suitable maximum yield strength is 1050 MPa.

Preferred embodiments are provided in the dependent claims 2 to 8. Residual elements (aka inevitable impurities, such as Cu, Ni, As, Pb, Sn, Sb, Mo, Cr, etc.) are defined as elements which are not added on purpose to steel and which cannot be removed by simple metallurgical processes. Some of the elements can be present as residual element, like nickel, but nickel can also be added to the steel, in which case it would no longer be a residual element or inevitable impurity, but an alloying element. In many cases it cannot be determined in the finished steel product whether (e.g .) nickel was added as an alloying element, or whether it was present as a residual element. Residual elements enter steel from impurities in ore, coke, flux and scrap; from these, scrap is considered to be the main source of residuals. Consequently, the level of residual elements in the Electric Arc Furnace process route (100% scrap based) is significantly higher than in the Basic Oxygen Steelmaking process route. The most commonly found residuals are Cu, Ni, Cr, Mo, and Sn. The acceptance limits of these residuals depend mainly on product requirements. Preferably the steel according to the invention is produced in a BOS-process route, in which case the allowable levels of residual elements for Cu, Ni, Cr, Mo and Sn are 0.04, 0.04, 0.04, 0.02 and 0.02 wt.% respectively. For Sn a preferred maximum residual level is 0.002 wt.%.

The chemistry is designed to achieve a fine-grained martensitic single phase microstructure. The steel has high strength (high yield strength and ultimate tensile strength) and high bendability and therefore has high crash crashworthiness. The high bendability resists fracture in the bend-folds during crash, in combination with the high strength allows high crash energy absorption.

The elements in this chemistry have the following roles:

Carbon is added as the main strengthening element and also to promote hardenability. Carbon is controlled to a value between 0.13 wt.% to 0.25 wt.% to ensure the formation of lath martensite instead of plate martensite and to attain the desired target strength. Flowever, when the carbon content exceeds 0.25 wt.%, the bendability and the weld toughness (weldability) are reduced . Preferably the carbon content is at least 0.15 wt.%, or even at least 0.17 wt.%, and a suitable maximum value is 0.22 wt.%.

Manganese is a hardening element. Manganese must not be less than 0.40 wt.% to ensure sufficient hardenability. Flowever, manganese content higher than 1.2 wt.% increases the risk of the formation of segregations with band-type microstructures associated with a ductility decrease. Preferably the manganese content is at least 0.60 wt.%, and a suitable maximum value is 1.10 wt.%, preferably the manganese content is at most 1.00 wt.% .

Chromium is effective for increasing hardenability. Flowever, in combination with the other elements of the composition that also increase hardenability, a Cr addition higher than 1.0 wt.% is not necessary and is excessively cost increasing. The minimum chromium addition is 0.04 wt.%, preferably 0.08 wt.%. A suitable minimum value is 0.10 wt.% and a suitable maximum value is 0.80 wt.%. Boron is added to increase the hardenability of the alloy. Small amounts suffice in the range of 5 to 50 ppm by weight provided the boron is not bound to nitrogen (i .e. soluble B is needed). At a content of at least 0.0005 wt.% (5 ppm), boron prevents the formation of ferrite on cooling and increases hardenability of the steel. Its content is limited to 0.005 wt.% because above this level, its effect is saturated and further addition is not effective. A preferable range of the B content is from 10 ppm to 35 ppm.

Aluminium is added as a deoxidizer in the liquid state and a total aluminium content is needed of at least 0.020 wt.%. Aluminium also combines with N to protect the effectiveness of the boron. The formed AIN refines austenite grains size during annealing. Aluminium may also segregate or partition to grain boundaries to affect the grain size. However, when Al exceeds 0.2 wt.%, there is a risk of formation of coarse aluminates in the liquid state, which could reduce the ductility of the steel. Consequently the aluminium ranges from 0.020 wt.% to 0.2 wt.%. A preferable range of the Al content is from 0.03 to 0.10 wt. %, and more preferably from 0.04 to 0.08 wt.%.

Nitrogen is introduced during steelmaking. The hardenability effects of B can be undermined by the presence of N in the steel since N can form boron nitrides and reduce the soluble B in the matrix. Therefore, N content in the steel was limited to a maximum of 80 ppm by weight, preferably to 70 ppm, more preferably to 60 ppm.

Titanium is an element effective for the purpose of improving the strength and for forming Ti-based sulphides with relatively little effect on the local formability and reducing the harmful MnS. Titanium also combines with N form TiN in the liquid state of steel during steelmaking so as to avoid formation of BN at lower temperatures during cooling or reheating to fully develop the effect of the element boron to enhance hardenability. At least 0.001 wt.% Ti is needed for this purpose. If an excess of titanium is added, coarse TiC particles can be formed during steelmaking, or during reheating and hot rolling or during coiling and subsequent cooling on the coil. These TiC particles cannot be completely disscolved in the following austenization process. The coarse TiC particles can deteriorate the bendability of the product and reduce the hardenability of the steel. Therefore, Ti content needs to be kept low to avoid the formation of TiC particles. Therefore, the Ti content in the steel of the present invention should be in the range of 0.001-0.040 wt.%. Preferably Ti is at most 0.030 wt.% and more preferably at most 0.020 wt.%, even more preferably at most 0.010 wt.% or even at most 0.005 wt.%.

Niobium forms fine Nb(CN) in combination with carbon and/or nitrogen, which refine austenite grain size during annealing and is an optional element in the steels according to the invention. This finer austenite grain results in finer lath structure and increased ductility and toughness. Nb-precipitates also assist in hydrogen capture to improve the resistance to the delayed fracture of the steel. However, the Nb content higher than 0.1 wt.% causes an excessive increase in rolling forces during hot-rolling. Consequently niobium is maximised to 0.1 wt.%. If added, the minimum niobium content is 0.01 wt.%. Preferably Nb is at most 0.05 wt.% and more preferably at most 0.038 wt.%.

Sulphur causes the formation of sulphides which lower bendability and ductility of the steel or the press hardened part. The sulphur content must not be higher than 0.010%, and preferably not higher than 0.005%. Lowering the sulphur content to a very low content (e.g. 0.0002 or 0.0005 wt.%) requires a costly desulfurization treatment, without significant additional benefit for these steels.

When present in quantity higher than 0.030 wt.%, phosphorus can segregate at the austenite grain boundaries and reduce toughness of the steel. However, P content lower than 0.001 wt.% needs costly treatment at the liquid stage, without significant benefit on the mechanical properties of the steel or the press hardened part. Preferably the phosphorus content is at most 0.020 wt.%. More preferably the maximum phosphorus content is 0.012 wt.%.

Silicon is usually added to improve the strength through solution hardening and transformation hardening. However, with increasing amounts of silicon, the HAZ toughness, weldability and coatability are impaired. Meanwhile, due to better affinity of silicon with oxygen than that with iron, silicates having low melting point tend to be generated easily during welding, which increases slag and the mobility of molten metals, and thus impacts the quality of the weld. Therefore, silicon content shall be controlled strictly. The content of silicon in the invention is controlled to be less than 0.080 wt.%, preferably less than 0.07 wt.%, more preferably less than 0.06 wt.%, even more preferably less than 0.05 wt.%.

Calcium is optionally added to the steel composition to modify the shape of sulphides. Calcium combines with sulphur and oxygen, thus creating oxysulphides that do not exert a detrimental effect on ductility, as in the case of elongated manganese sulphides. Furthermore, these oxysulphides act as nucleants for a fine precipitation of (Ti,Nb)(C,N). This effect is saturated when the calcium content is higher than 0.003 wt.%. If added, the minimum calcium amount is 0.0001 wt.%.

The composition is designed to have a carbon equivalent CE of less than 0.46 wherein CE = C + Mn/6 + (Cr+Mo+V)/5 + (Ni+Cu+Si)/15, where CE is the carbon equivalent, and C, Mn, Cr, Mo, V, Ni, and Cu are in wt.% of the elements in the alloy. The amount of the elements that increase the electric resistance of steel such as Si, Mn and P is reduced to expand the suitable welding current range. So the weldability of the invented steel is good. Preferably the carbon equivalent is 0.45 or less, more preferably 0.44 or less.

The alloying elements Copper, Nickel, Molybdenum, Vanadium can be optionally present to further increase the hardenability of the steel. As the composition of the invented steel is designed from the point of view of weldability, the total amount of these elements should be less than 0.5 wt.%, and the amount of each element should be less than 0.10, 0.20, 0.20 and 0.10 wt.% for copper, nickel, molybdenum and vanadium respectively, if added as alloying elements. In a preferable embodiment the steel contains no nickel, copper, molybdenum and vanadium as alloying elements, i.e. any nickel, copper, vanadium or molybdenum present is present as an inevitable impurity (aka residual element). For practical purposes, a maximum amount of 0.060 wt.% Ni is acceptable as residual element, and preferably a maximum amount of 0.030 wt.% Ni as a residual element is acceptable. For vanadium a preferred maximum residual level is 0.001 wt.%. At these amounts the effect of the nickel and vanadium is negligible.

The cold-rolled and heat treated steel according to the invention has a preferably fully martensitic microstructure. The total amount of bainite and retained austenite should be as low as possible, and preferably limited to at most 10%, preferably at most 5%, more preferably at most 2% . Small amounts of bainite and/or retained austenite are permissible. But a higher amount will lead to a lower bendability. Other microstructural constituents like ferrite or pearlite must be avoided. Presence of ferrite and pearlite are detrimental to achieve high strength, high ductility and high bendability. This martensitic microstructure is instrumental in achieving the high values for ultimate tensile strength (aka Rm or UTS) and the yield strength (aka Rp or YS) in combination with the bendability and the weldability. A martensitic microstructure of the invention provides a high yield strength in combination with excellent bendability, and tensile elongation higher or at least in the range of ultrahigh strength 22MnB5. As a result, the steel can absorb high crash energy. The steel chemistry and the thermal cycle are the most critical steps to achieve the desired microstructures, properties and performance. The alloying elements (C, Mn, Cr, B) are deliberately controlled in the steel of invention to increase the hardenability so that ferrite and pearlite can be avoided during cooling after re crystallization and austenization in a continuous production line or during transfer of blanks from the reheating furnace to the hot forming press.

The high YS and UTS are derived from martensite resulting from the good hardenability and fine grain size of the alloys. The high bending angles emanate from lowering the weak phase boundaries in the microstructures due to single phase martensite. Usually the bending cracks initiate from the weak phase boundaries. The high fracture toughness results from also a well-tailored microstructure as defined for the bendability. High bendability and high fracture toughness resulted in good crash resistance of the material.

The cold-rolled steel according to the invention is preferably provided with a thickness between 0.5 and 3 mm in addition to a yield stress YS of at least 880 MPa, an ultimate tensile strength UTS of at least 1180 MPa, and a high ductility characterized by an average bending angle (BA-L+BA-T)/2 higher than 90°. The steel combines high strength with excellent bendability which makes it perfect for energy absorption applications. The steel is excellently coatable due to the low silicon content and has a good weldability due to the low CE-value, and a substantial abrasion wear resistance. The cold-rolled martensitic steel may be in the form of a cold-rolled strip, sheet or coiled strip having a thickness of between 0.5 to 3 mm .

The steel can be also applied to manufacture the automotive parts by using conventional hot press forming process. Typical applications for the invented steel include energy absorption parts such as front and rear rails and lower B-pillars.

The steel according to the invention has a yield strength of at least 880 MPa and an ultimate tensile strength of at least 1180 MPa .

In an embodiment the steel according to the invention has a bending angle BA-L as measured in accordance with the procedure described in the description of at least 100°, preferably of at least 110°, more preferably of at least 120° and even more preferably of at least 125°.

In an embodiment the steel according to the invention has a bending angle BA-T as measured in accordance with the procedure described in the description of at least 100°, preferably of at least 110°, more preferably of at least 120° and even more preferably of at least 125°.

In an embodiment the steel according to the invention has bending angles BA-T and BA-L after bake hardening at 180 °C for 20 minutes, as measured in accordance with the procedure described in the description, of at least 100°, preferably of at least 110°, more preferably of at least 120° and even more preferably of at least 125°.

In an embodiment the steel according to the invention has an average bending angle (BA-L+BA-T)/2 higher than at least 100°, preferably of at least 110°, more preferably of at least 120° and even more preferably of at least 125° or even more than 130°.

In an embodiment the excellent bendability is maintained after bake hardening at 180 °C for 20 minutes. The bending angles obtained by the steels according to the invention after bake-hardening are at least 90° .

In an embodiment the martensitic single phase microstructure has a grain size ( = martensite packet size) below 10 pm, preferably below 8 pm, more preferably below 7 pm, even more preferably below 6.75 pm to obtain high bendability and high strength. The size is preferably at least 1 pm, more preferably at least 2 pm .

According to a second aspect the invention is embodied in a process for producing a cold-rolled martensitic steel having high strength and high bendability having a microstructure consisting of 90 to 100 vol.% of martensite and 0 to 10 vol.% in total of å (bainite + retained austenite), a yield strength of at least 880 MPa and an ultimate tensile strength of at least 1180 MPa as determined on JIS5 tensile specimens according to NEN- EN10002 : 2001 and an average bending angle (BA-L+BA-T)/2 of at least 90°, wherein BA- L and BA-T are the bending angles with the bending axis parallel (L) to or perpendicular (T) to the rolling direction of the cold-rolled steel as determined in a three-point bending test according to the VDA 238-100 standard, said process comprising :

casting a melt into a slab or strip having the following composition o 0.13 - 0.25 wt.% C;

o 0.40 - 1.20 wt.% Mn;

o 0.04 - 1.00 wt.% Cr;

o 5 - 50 ppm by weight B;

o 0.020 - 0.20 wt.% AI_tot;

o 5 - 80 ppm by weight N ;

o 0.001 to 0.040 wt.% Ti;

o 0 - 0.003 wt.% Ca;

o 0 - 0.030 wt.% P;

o 0 - 0.010 wt.% S;

o 0 - 0.080 wt.% Si ;

o 0 - 0.10 wt.% Nb;

o 0 - 0.10 wt.% Cu;

o 0 - 0.20 wt.% Mo;

o 0 - 0.060 wt.% Ni ;

o 0 - 0.10 wt.% V;

o å(Cu+V+Ni + Mo) < 0.50 wt.% ;

o CE = C + Mn/6 + (Cr + Mo + V)/5 + (Ni + Cu + Si)/15 < 0.46;

o remainder iron and residual elements,

hot-rolling the slab or strip to a hot-rolled strip;

cooling the hot-rolled strip at an average cooling rate of 15 to 100 °C/s on the run-out table to a coiling temperature of between 500 °C and 700 °C and allowing the coil to cool to room temperature;

pickling and cold-rolling the hot-rolled strip, followed by:

a) heat treating the cold-rolled strip by reheating the strip to a temperature above Ac3, soaking above Ac3 to form an austenite microstructure, followed by cooling at a first cooling rate CR1 higher than the critical cooling rate for the formation of martensite to a temperature below the martensite start temperature (Ms), followed by a further cooling with a second cooling rate CR2 to a temperature below the martensite finish temperature (Mf), wherein CR2 may be the same as CR1 or CR2 may be lower than CR1, or by

b) heat treating the cold-rolled steel as a sheet or blank cut from the cold- rolled strip by reheating the sheet or blank to a temperature above Ac3 and soaking above Ac3 to form an austenite microstructure in a reheating furnace of a hot-forming unit, and wherein the steel sheet or blank having the austenitic microstructure is hot-formed in a mould into the desired shape and subsequently cooled in the mould or outside the mould at a cooling rate CR1 higher than the critical cooling rate for the formation of martensite to a temperature below the martensite start temperature (M _s ), followed by a further cooling with a second cooling rate CR2 to a temperature below the martensite finish temperature (M _f ), wherein CR2 may be the same as CR1 or CR2 may be lower than CR1.

The steelmaking process is performed according to known steelmaking processes such as refining in a converter or electric arc furnace, and then secondary refining in a vacuum degassing furnace. The casting step may involve continuous thick or thin slab casting or strip casting, or ingot casting. The slab or strip is subsequently hot-rolled to a hot-rolled strip and cooled at an average cooling rate of 15 to 100 °C/s to a coiling temperature of between 500 °C and 700 °C, preferably between 550 and 700 °C, more preferably of at least 580 °C or even of at least 600 °C and/or at most 680° and allowing the coil to cool to ambient temperature. The microstructure after hot-rolling is not particularly critical, and a fine ferrite-pearlite microstructure would be ideal, because that would provide a good cold-rollable steel . Ambient temperature is the temperature of the surroundings, and often also referred to as room temperature. After cooling to ambient temperature and optional storage, the coil is uncoiled and pickled to remove the oxides from the strip surface and cleaned before being cold-rolled . The cold-rolling reduction should be at least 40% in order to ensure complete recrystallization during the following heat treatment. The cold-rolled strip is subsequently heat treated by reheating the strip to a temperature above Ac3 followed by soaking above Ac3 to form a fully austenitic microstructure, followed by cooling at a first cooling rate (CR1) higher than the critical cooling rate for the formation of martensite to a temperature below the martensite start temperature (Ms), followed by further cooling with a second cooling rate (CR2) to a temperature below the M _f point. CR2 may be the same as CR1 or CR2 may be lower than CR1. The critical cooling rate is defined as the slowest cooling rate which produces only martensite during austenite decomposition upon cooling the steel from above An to ambient temperature. In a CCT-curve the critical cooling rate is determined by the slowest cooling rate that just avoids the noses of other microstructure constituents that may occur during austenite decomposition, such as pearlite, ferrite or bainite. In Figure 3 a CCT- diagram (Fig . 11-4.10, Chapter 11, page 428, Elements of Materials Science and Engineering, van Vlack, 4 ^th edition (1980)) is shown that demonstrates the critical cooling rate for martensite formation (CR _M ) for eutectoid steel (0.8 wt.% C). At a lower cooling rate the pearlite 'nose' is not avoided and pearlite is formed (the hatched area) . M _s is the temperature at which the martensite formation starts. M _f is the temperature at which the martensite formation finishes. A practical problem with M _f is that the martensite fraction during cooling approaches the maximum achievable amount only asymptotically, meaning that it takes very long for the last martensite to form . For practical reasons and in the context of this invention, M _f is therefore defined as the temperature at which 90% of martensite has formed . This value can be easily determined from dilatation tests. An example of this determination is provided in Figure 4.

The heat treatment may be performed in a separate furnace, in a continuous annealing line, a hot-dip coating line or in hot stamping equipment, depending on the application of the steel. So it is possible to produce a coiled strip, a strip, a sheet or a blank produced from the strip or sheet of the steel according to the invention, or a part produced from the steel according to the invention.

Depending on the products, the cold-rolled strip is heat treated either in a continuous annealing line as sheets or in the hot press forming tools as blanks.

The thermal cycles are now described in detail. The cold-rolled sheet or strip is heated up at a suitable heating rate between 2 and 100 °C/s to a temperature in the single-phase austenitic phase field, i.e. higher than Ac3, preferably to a temperature between (Ac3 + 10) and (Ac3 + 50) °C. Soaking at a temperature above Ac3 preferably takes place over a time period between 0.1 and 10 minutes. Preferably the soaking time is at most 7 minutes, preferably at most 5 minutes. Austenization in the single austenite phase field is necessary to avoid the presence of any proeutectoid ferrite in the final microstructure and to dissolve all carbides, and the sufficient time is prescribed in order to form a homogenous austenite from the initial cold-rolled (usually ferritic-pearlitic) microstructure. Using too low an austenization temperature or too short an austenising time will result in undissolved precipitates such as carbonitrides or incomplete austenization which is detrimental to the final properties. Too high a temperature or too long a time will cause austenite grain growth which is undesired since it will affect the bendability and toughness. It is also economically unattractive.

After soaking, the steel is quenched in a medium which ensures a sufficiently high cooling rate for the material and dimensions of the workpiece and results in the formation of more than 90 vol.% martensite, preferably more than 95 vol.%, more preferably more than 98 vol .% martensite, and any remainder being lower bainite and/or retained austenite (å(bainite + retained austenite)) being at most 10 vol .%, preferably at most 5 vol .% and more preferably at most 2 vol .% . Only insignificant amounts of other microstructures may be present, provided these other microstructures do not affect the favourable properties of the steels according to the invention. For example, if during hot-press forming an uncoated steel is heated in an oxidising atmosphere, there may be a small degree of decarburisation at the surface of the steel . Also when annealing a coated steel there may be a small amount of coating metal diffusing potentially leading to an intermetallic compound of coating metal and iron. Preferably however no such other microstructures are present in the steel according to the invention.

The quenching stop temperature is below the M _f point of the steel (defined herein above as the temperature at which 90% martensite has transformed). M _f is typically 300 °C or below, or even below 250 °C. The quenching rate CR1 in the temperature range from the austenization temperature to M _s (CR1 in Figure 5) should be higher than the critical cooling rate (CR _M (see figure 3)) to prevent the formation of other microstructures than martensite. For the steels according to the invention a suitable cooling rate CR1 is at least 50 °C/s, preferably higher than 70 °C/s, more preferably higher than 90 °C/s. Optionally, a pre-cooling stage can be applied between the austenization temperature and the An. An depends on the chemical composition and the cooling rate, and is between about 650 to 780 °C for the steels and cooling rates according to the invention. The cooling rate CR2 from Ms to below M _f (CR2 in Figure 5) can be the same as the CR1 (i .e. at least 50 °C/s, preferably higher than 70 °C/s, more preferably higher than 90 °C/s) or lower than the cooling rate CR1. A suitable minimum cooling rate for CR2 is 1 °C/s.

If the heat treatment is conducted in a continuous product line, after austenization at a top temperature of e.g . 900 °C for 0.1-10 minutes, the steel sheet can be cooled at a slow cooling rate of 1 to 10 °C/s to 700 to 780 °C and then cooled at a fast rate CR1 higher than 50 °C/s to a temperature below M _s and then cooled at a low rate CR2 to a temperature below M _f . Preferably the soaking time is at most 7 minutes, preferably at most 5 minutes. Optionally, the steel sheet is then reheated to a temperature between 400 and 500 °C and hold for 0.1 to 2 min for tempering, depending on the line speed . Optionally, a hot dip coating (such as Zn or Zn-alloy, Al or Al-alloy, Al-Si) can be applied in this step. Thereafter, the steel sheet is cooled down at a cooling rate 2-10 °C/s to room temperature. The temperatures and cooling rates depend on the lay-out of the continuous annealing line, and the above example is not intended to be limiting .

Optionally, the hot-dip coating can be applied to the quenched material as a separate process. If a Zn based coating (including zinc alloys) is applied, the Zn bath temperature is typically between 440 and 480 °C. If an aluminium-based coating is applied, the Al bath temperature is typically between 660 and 720 °C. The coating may also be an aluminium- based coating, for example by dipping in an aluminium alloy bath containing, in addition to aluminium, up to 11% silicon, and from 2% to 4% iron, the sheet having a high mechanical resistance after thermal treatment and a high resistance to corrosion.

Coating of the strip, sheet or coiled strip of the steel according to the invention or part produced from the steel according to the invention may also be performed by electrogalvanizing or a chemical conversion treatment or physical vapour deposition. The coating can be applied prior to, during or after the heat treatment process. Optionally the coated or uncoated strip, sheet or coiled strip of the steel according to the invention or part produced from the steel according to the invention may be painted and subjected to paint-baking.

If the heat treatment is conducted in a hot stamping tool, and not in a continuous annealing furnace on the entire strip, the cold-rolled sheet should be first cut into blanks with the proper size for the intended products.

The cold-rolled blanks are heated up at a suitable heating rate between 2 and 100 °C/s to a temperature in the single-phase austenitic phase field, i .e. higher than Ac3, preferably to a temperature between (Ac3 + 10) and (Ac3 + 50) °C. Soaking at a temperature above Ac3 preferably takes place over a time period between 0.1 and 10 minutes. Preferably the soaking time is at most 7 minutes, preferably at most 5 minutes.

After reheating, the forming is done conventionally: the blanks are transferred to the press, hot formed, quenched (CR1) and the formed article is taken out from the press below the Ms temperature of the steel, and further cooled to room temperature with CR2, preferably in air. During press forming, the cooling rate (CR1) should be higher than CRM to prevent the formation of other microstructures than martensite. For the steels according to the invention a suitable cooling rate is at least 50 °C/s, preferably higher than 70 °C/s, more preferably higher than 90 °C/s in the temperature range from 750 °C to M _s to avoid the formation of non-martensitic microstructures. The cooling rate CR2 can be the same (i.e. at least 50 °C/s, preferably higher than 70 °C/s, more preferably higher than 90 °C/s) or lower than the cooling rate CR1. A suitable minimum cooling rate for CR2 is 1 °C/s.

Before the steel blank is hot press formed, coating can also be applied through a continuous hot dip coating line. In this case, the top annealing temperature can be lower in the range of 700 to 780 °C (no annealing in the austenitic single-phase field is necessary at this stage) and then a hot-dip coating can be applied as a standard process stage in the production line. The annealing of the steel in the austenite single phase field occurs later, during the hot-press forming .

Temper rolling or tension levelling for shape correction of the cold-rolled strip may be performed .

According to a third aspect the invention is also embodied in specific parts e.g. for automotive applications, such as lower B-pillar, reinforcement parts for automotive bumpers and doors, producible by hot press forming or by a roll forming process.

The invention is now further explained by means of the following, non-limiting examples. Examples

Table 1. Chemical composition (all elements are in wt.%, except N and B in ppm) .

The alloys listed in Table 1 were cast, hot-rolled and cold-rolled to 1.5 mm thick strips. Steels Al, A6, A7 and A9 are the invented compositions, while steels R1 (22MnB5 steel) and R2 (DP1000 steel) are reference compositions. The slabs were reheated at 1225 °C, rough-rolled to 39 mm and hot-rolled to a final thickness of 4 mm, with a finish rolling temperature above 900 °C (i .e. above Ac3, see table 1) . The average cooling rate at the run out table after hot-rolling and before coiling was 25 °C/s. The coiling temperature was 650 °C. The hot-rolled microstructure consisted of ferrite and pearlite. This was then pickled and subsequently cold-rolled to 1.5 mm thickness. Calcium treatment of the steels in Table 1 would lead to a calcium content of 0.0015 to 0.0025 wt.%)

Dilatometry was performed on cold-rolled samples of 10 mm x 5 mm x 1.5 mm dimensions (length along the rolling direction) . The samples were heated to 900 °C at a rate of 15 °C/s, held for 2 minutes at 900 °C and then quenched at 100 °C/s to room temperature (CR1 =CR2) . The critical phase transformation points were determined from the dilatometry curves by observing the inflection in the dilatation curves (dilatation vs temperature, see e.g. M . Gomez et al ., Phase Transformation under Continuous Cooling Conditions in Medium Carbon Micro-alloyed Steels. Journal of Materials Science & Technology, 2014, 30(5) : 511-516) and are given in Table 1.

For the inventive steels and the reference steel Rl, blanks of 220 mm x 110 mm x 1.5 mm were prepared from cold-rolled sheets and they were subjected to the following thermal cycles in a hot dip annealing simulator (HDAS) . The blanks were reheated at 15 °C/s to 900 °C, soaked for 2 minutes in nitrogen atmosphere to minimize surface degradation. After the austenization, the blanks were cooled at 4 °C/s to 750 °C and then cooled at 80 °C/s to 400 °C and then cooled at 35 °C/s to 20 °C. Thereafter, the blanks were reheated (at 8 °C/s) to 460 °C in 53s and hold for 17s and thereafter cooled down to room temperature at 9 °C/s to simulate the hot dip Zn coating process in a continuous production line.

After heat treatment, a paint bake hardening process at 180 °C for 20 minutes was applied for some samples. For steel R2, blanks of 550 mm x 110 mm x 1.5 mm were prepared. The heat treatment was conducted in a continuous annealing simulator. The blanks were reheated at 15 °C/s to 810 °C, held for 2 minutes and then cooled to 700 °C at 3 °C/s to 700 °C and then cooled to 470 °C at 40 °C/s and hold for 60s and then cooled to room temperature at 9 °C/s.

From the heat-treated samples, JIS5 tensile specimens (gauge length = 50 mm; width = 25 mm) along the rolling direction were prepared and tested following NEN- EN10002-1 : 2001 standard to determine tensile properties (yield strength YS (MPa), ultimate tensile strength UTS (MPa), total elongation TE (%)). Bending specimens (40 mm x 30 mm x 1.5 mm) from parallel and transverse to rolling directions were prepared from each of the conditions and tested till fracture by three-point bending test according to the VDA 238-100 standard. The samples with bending axis parallel to the rolling direction are identified as longitudinal (L) bending specimens whereas those with bending axis perpendicular to the rolling direction are denoted as perpendicular (T) bending specimens. The measured bending angles at 1.5 mm thickness were also converted to the angles for 1 mm thickness (= original bending angle x square root of original thickness) . For each type of test, three samples were tested and the average values from three tests are presented for each condition.

About less than 2% of retained austenite was detected by X-ray measurement in all the heat-treated samples. The substructures of martensite were analysed using scanning electronic microscopy. The EBSD analysis was carried out on a FEI Quanta 600 FEG-SEM . An area of 100 pm x 100 pm was scanned at a step size 0.1 pm. Since martensite packet boundaries are high angle grain boundaries (FIAGB), the minimum boundary angle for martensite packet size estimation was also set to 15°. To reveal the martensite packet boundaries, grain boundaries with mis-orientation angles between 15° and 50° were reconstructed on the basis of images as shown in Figure 1 for the A6 steel . The martensite pocket size is given by diameter of equivalent area circle. The size distribution is shown in Figure 2 for A6 steel . The average pocket size is determined as å(area of individual grain x area fraction of each grain) and is given in Table 2. Table 3 shows tensile properties and bending angle. High bending angles, larger than 120° at 1 mm thickness are achieved. Bake hardening slightly increases the strength but does not significantly affect the elongation and the bending angle. The reference R1 (22MnB5) steel has a slightly higher strength but have a much lower bending angle. The reference R2 (DP1000) steel has yield strength and ultimate tensile strength in the same range as the invented steel and also higher tensile elongation, but has much lower bendability. After bake hardening at 180 °C for 20 minutes, the ultimate tensile strength, the elongation and the bending angle are not changed although the yield strength of the steels is slightly increased. Figure 1 shows EBSD Image Quality (IQ) maps from which the EBSD processing software can identify the high angle grain (FIAGB) boundaries. FIAGB with an angle difference between 15 and 50° were used to identify the martensitic grain boundaries in steel A6. Figure 2 shows the martensite pocket size distribution in steel A6. The EBSD software determines an average martensite pocket size (i .e. martensite "grain size") on the basis of the data depicted in Figure 2.

Table 2 : Cold-rolled and heat-treated microstructure and wear resistance test

Table 3 : Tensile properties and bending angle of heat treated steels

*BA-L or BA-T are the bending angles with the bending axis parallel (Longitudinal, L) to or perpendicular (transverse, T) to the rolling direction.

Additionally the abrasion wear resistant property was measured from hot-rolled 4 mm thick strips, which have a fully martensitic microstructure. The common testing standard ASTM G65 - the dry sand rubber wheel abrasion test was carried out according to the procedure B - 10 minutes testing time. The abrasive material is the rounded quartz grain sand, as specified as AFS 50/70 silica sand in the standard ASTM G65, was used for the wear testing . The wear sample weight was measured before and after wear testing with a scale to an accuracy of 10 ⁴ g to determine weight loss. The relative wear life to reference steel S355 was calculated by dividing the weight loss of S355 by the weight loss of the inventive steels. The results are shown in Table 2. The S355 steel is a hot-rolled steel grade and has a composition of C 0.096, Si 0.37, Mn 1.48, Ni 0.4, Cr 0.016, Cu 0.2, Nb 0.022, Al 0.034, and N 0.006 with a microstructure of ferrite and pearlite.