A COLD-ROLLED AND HOT-PRESS FORMED HIGH STRENGTH STEEL PRODUCT WITH EXCELLENT HOLE EXPANDABILITY AND PROCESS OF PRODUCING THE SAME

Field of the invention

This invention relates to a high strength steel sheet with excellent hole expandability and to a method of producing the same.

Background of the invention

The demand for advanced high strength steels (AHSS) with higher strengths is increasing in the automotive industry. Automotive parts are often cold formed and the manufacturing process includes shear cutting and flanging. Problems often may when the sheared edges are stretched during pressing as they are susceptible to crack formation. The flangeability (characterized by hole expansion capacity (HEC)) tends to decrease with an increase in the strength of a steel.

There are two types of formability, drawability and stretch flangeability. The drawability is governed by global ductility, characterized by elongation in tensile tests, the stretch flangeability is governed by local ductility, characterized by the hole expansion ratio in hole expansion tests or by the bending angle in bending tests, respectively. Global ductility and local ductility are incompatible. The total elongation depends mainly on the strain localization in the matrix phase while the local ductility depends on the micro-scale uniformity: the more uniform the microstructure is, the better the stretch flangeability is, although stretch flangeability is also known to decrease with increasing strength. Bendability can also be construed to be similar to stretch flangeability, as a fracture due to a large local deformation, and can be adjusted by the scatter of hardness (i.e. the distribution of hard and soft phases in the microstructure). As steels became increasingly stronger, they simultaneously became increasingly difficult to form into automotive parts. While there have been major improvements recently in the trade-off between ductility and strength, the stretch flangeability of AHSS remains a critical issue.

Dual-phase (DP) steels and low alloy transformation-induced plasticity (TRIP) steels are known to show high elongation. However, since these steels are characterized by a large difference in hardness between the main phase, which consists of soft polygonal ferrite and the hard second phase of martensite. These types of steel are inferior in stretch flangeability (local ductility) because of voids that form at the interphase between these phases during deformation processes such as punching. In contrast, Complex phase (CP) steels yield a lower global ductility as a result of the reduced amount of the soft ferrite and a higher local ductility as a result of the reduced hardness difference among the phases and together with a more uniform distribution of microstructure. In practice, the application of AHSS steels (DP, CP and TRIP) to car components is still limited by their formability. Therefore, improving formability and manufacturability is an important issue for AHSS applications. To remedy this problem, bainite single phase and bainitic ferrite single phase steel sheets have been developed in order to reduce the hardness difference between the main and second phases, but it cannot be said that the balance of strength and hole expansion ability is adequate, particularly in materials with tensile strength of 780 MPa class and higher.

Objectives of the invention

It is an objective of the present invention to provide a cold-rolled high strength steel. It is also an object of the invention to provide a cold-rolled high strength steel with an improved formability and HEC value in particular.

It is also an object of the invention to provide a cold-rolled high strength steel excellent in hole expandability, aiming at the application for automobiles, such as passenger cars and trucks, etc., for industrial machines, or the like. It is also an object of the invention to provide a process for producing these cold- rolled high strength steels.

Description of the invention

According to a first aspect the invention is embodied in a cold-rolled and hot press formed steel product comprising (in wt.%): - C: 0.050-0.180;

- Mn: 1.50-3.50;

- Si: 0.40-1.80;

- Al: 0.01-1.00;

- Al + Si: 0.50-2.00; - P: at most 0.050;

- S: at most 0.020;

- N : at most 0.010; the steel optionally also comprising one or more elements selected from

- B: at most 0.0050; - Ti: at most 0.050;

- Nb: at most 0.050;

- V: at most 0.10;

- Mo: at most 0.20;

- Cr: at most 0.30; - Cu: at most 0.50;

- Ni: at most 0.50;

- Ca: at most 0.0030; - Rare earth elements (REM): at most 0.0030; the balance being Fe and inevitable impurities, having a microstructure containing

BF and/or PM, å(HF+LF) and second phase comprising FM, RA and carbide, in the following amounts (in vol.%):

- F at most 35%;

- HF at most 5%;

- (BF+PM) at least 60% and (BF+PM) at most 90%;

- Carbides at most 3%;

- FM at most 5%;

- RA 3 to 10%;

- (carbides + FM + RA) at most 15%,

In the context of this invention the following abbreviation for microstructural components will be used:

BF: bainitic ferrite is formed during austempering in the temperature range between Ms and Bs.

PM: partitioned martensite is formed during austempering first at a temperature below the Ms point and then heated to the temperature range between Ms and Bs.

HF: ferrite formed at intercritical temperatures (high temperature ferrite)

LF: ferrite formed at cooling below Acl after annealing (low temperature ferrite)

F: å(HF + LF)

RA: retained austenite

FM: fresh martensite refers to the martensite formed during cooling after the austempering treatment.

When in the context of this invention use is made of the term "strip", then this term is deemed to encompass a sheet or any sheetlike material produced from the strip.

In order to achieve high strength with excellent hole expandability, it is necessary to form a uniform microstructure including BF and/or PM as matrix phase. The present invention solves the problems using following measures.

Firstly, the amount of the second phases (FM + RA + carbides) is controlled to below a certain level. The boundaries between FM, carbides and the ferritic matrix are the original sources of void formation during forming. The RA can produce the TRIP effect to increase the elongation. At the same time the transformation of RA to FM during the early stage of deformation can lead to low HEC values.

Secondly, LF has a higher oversaturation with carbon and a smaller grain size than HF. Therefore, LF has a higher hardness, which is beneficial to obtain high HEC property. The presence of LF instead of HF results in a smaller hardness difference between the F and matrix (BF and PM). HF should be at most 5%, preferably at most 1%, even more preferably at most 0.2% and most preferably be avoided altogether and be absent (0%).

The inventors found that it is possible to obtain a uniform microstructure comprising BF and/or PM as the matrix, in which (FM + RA + carbides) is less than 15% to have high strength and good formability using a specific chemical composition and a specific method to produce it.

Intensive studies were conducted on high-strength galvanized steel sheets having a tensile strength (TS) of 800 MPa or more, an elongation (EL) of 10% JIS5 or more and a hole expansion ratio of 40% or more and the inventors observed that these properties could be obtained by cold-rolled and continuously annealed steel sheet or strip having a chemical composition and a microstructure as claimed.

In conventional steels, bainite forms in steels at temperatures between Bs and Ms. Bainite contains ferritic plates intertwined with finely dispersed carbides, which is supersaturated with carbon and has a high density of dislocations. The martensite is obtained when the steel is quenched below Ms. Although FM is very strong, it is normally brittle as a highly supersaturated solid solution of carbon in a-iron. The mechanical properties are typically further adjusted by heat treatment called tempering in an elevated temperature. In conventional steels, during tempering, carbon is rejected from the supersaturated solid solution and forms finely divided carbide phases within the martensite. Conventional tempered martensite therefore contains a fine dispersion of carbides in an a-iron matrix which decrease the overall elongation and formability properties of the steel.

In the steel according to the invention, the BF and/or PM formed during austempering treatment are different from the conventional bainite and martensite. During the specified austempering process, bainite and/or martensite first form. The carbon is rejected from the supersaturated solid solution but the formation of carbides is suppressed or retarded, due to the presence of 0.5% or more of the sum of Si and Al in the composition and due to time control during further heat treatment. As a result, the carbon partitions from bainite or martensite to austenite which leads to carbon enriched retained austenite with higher stability. The resultant bainite or martensite is carbide free and is referred to as BF or PM. The optical microstructure of BF and PM is similar as both have a plate-like microstructure with an ultrafine grain size (typically about 5 pm long and about 200 nm thick) although the PM has a finer size of substructure as the formation temperature is lower. The definitions of the critical phase transition temperatures such as Acl and Ac3 during heating and Ar3 and Arl during cooling are well known to a person skilled in the art and are dependent on the alloy composition. Bs is defined as the temperature at which, during cooling, transformation of the austenite into bainite starts. Bn is defined as the nose temperature of the bainitic transformation in the time-temperature transformation (TTT) curve of a steel, at which transformation of the austenite into bainite has the fastest kinetics. Ms is defined as the temperature at which, during cooling, transformation of the austenite into martensite starts. All these temperatures can be accurately determined by dilatometer experiments combined with hardness measurements and microstructure evaluation. Ac3 is a function of the heating rate and Ac3 in this invention was measured at a heating rate of 2 °C/s.

Alternatively, Ac3, Ar3, Bs, Ms and Bn points can be calculated beforehand, using the following empirical formulae, and assuming a heating rate of 2 °C/s:

AC ₃ = 952-203[C]° ⁵ +44.7[Si]-30[Mn]+400[AI]-ll[Cr]-15.2[Ni]-20[Cu] + 70[P] (1)

Ar ₃ = 914-650[C]-134[Mn]-15[Cr]-20[Cu] + 179[Si]+ 150[AI]-6.85CR (2)

Bs = 839-86[Mn]-23[Si]-67[Cr] + 35[AI]°- ⁵ -270(l-exp(-1.33[C])) (3)

Ms = 539-423[C]-30.4[Mn]-7.5[Si]+30[AI]-12. l[Cr] (4) Bn = ((Bs+Ms)/2)-20 (5)

Here [M] denotes the content (wt. %) of the chemical element M. CR is the cooling rate. The calculated critical transformation points are given in °C.

Now the reasons for the limitations on the microstructure of the high-strength steel sheet according to aspects of the present invention will be described. The amount of each microstructure is given by volume percentages unless otherwise indicated. The use of the phrase "x to y" to indicate a numerical range means: from x including x to y including y and is equivalent with "x - y".

BF + PM: 60 - 90% The matrix of BF or a mixture of PM and BF ensures a good balance in the strength and ductility of the steel. It is difficult to achieve a TS of 800 MPa or more and a hole expansion ratio of 40% or more when the amount of BF and PM is less than 60%. On the other hand, when the amount is greater than 90%, the amount of the ferrite and retained austenite is not enough and the elongation of the steel is remarkably reduced. Therefore, the fraction of (BF + PM) is 60% to 90%, and preferably at least 65%.

F: O - 35%

The ferrite fraction is the sum of HF and LF. The presence of HF increases the total elongation but decreases strength and formability. The presence of F can accelerate the bainitic transformation during austempering and helps to increase the stability of the retained austenite. LF has a higher hardness, which is beneficial to obtain good HEC values. Therefore, a controlled amount of LF is allowable and HF should most preferably be avoided. It is difficult to achieve both a TS of 800 MPa or more and a hole expansion ratio of 40% or more when the area fraction of F is greater than 35%. Therefore, the area fraction of F is 0% to 35%. Preferably F is at most 32% and more preferably F is at most 30%. F is preferably at least 4% and more preferably at least 6%.

FM: 0 - 5%

Some amount of FM may be produced during final cooling if the bainitic transformation is incomplete during the austempering treatment. The presence of FM further increases the strength of the steel. However, the hole expansion ratio becomes too low when the amount of the FM is greater than 5%. Therefore, the volume percentage of FM is 0 to 5%, preferably limited to at most 3%.

RA: 3 - 10% RA is useful for increasing El by producing the TRIP effect. The volume fraction of retained austenite needs to be 3% or more to achieve such an effect. However, when the volume fraction is greater than 10%, the RA is not stable enough as the carbon content of the RA is too low. It will transform to fresh martensite at an early stage of deformation which deteriorates the stretch flangeability. Therefore RA is 3 to 10%, preferably at least 5% and/or preferably at most 8%.

It is recommended that the carbon content in the RA is 0.9% or more. This carbon content greatly affects the TRIP-effect, and controlling it to 0.9% or more is particularly effective for improving elongation while not damaging the stretch flangeability. It is preferably at least 1.0%, more preferably at least 1.2%. The higher the carbon content of the RA is, the more preferable it is. However, in practice, the upper limit of the carbon content in the RA that can be reliably achieved is 1.6%.

Carbides: 3.0% or less

Carbides may be present in pearlite or as cementite or as other metal alloying carbides (exists at grain boundaries or at finely dispersed inside lath-structured ferrite) when the cooling rate is not high enough or when the overageing temperature is too high and/or the overageing time is too long. The formation of carbides reduces the amount of the retained austenite and the carbon content in the retained austenite, which deteriorate both the strength and ductility. In order to achieve good formability, in any case, it is necessary that the amount of carbides is 3% or less, preferably carbide free. In the case where the amount of carbides is more than 3%, there is a decrease in formability. Therefore, the amount of the carbides should preferably be limited to be less than 3.0%, more preferably less than 1.0%, most preferably 0%.

FM+RA+carbides: The sum of these phases is referred to as the second phase and is controlled to be 15% or less. More preferably less than 10%.

When the above microstructure conditions are satisfied, high strength and excellent formability are achieved. The volume fraction of each of the metallographic structures, F, BF+PM, and FM can be determined using an optical microscope and scanning electron microscopy. The volume fractions of RA and carbides are determined using an X-ray diffractometer (XRD). The details of the microstructural characterization are described in the experimental section.

Secondly, the above-mentioned chemical composition will be described. All compositional percentages are in weight percent (wt.%) unless otherwise indicated.

C: 0.050% - 0.180%

Carbon (C) is a necessary element for strength and ductility by forming the required microstructure. When the content of C is less than 0.050%, it is difficult to secure 60% or more BF and/or PM in the microstructures. On the other hand, when the carbon content is greater than 0.180%, too much fresh martensite or retained austenite can form in the most of available production lines due to the limited overageing time. It is not possible to achieve a desired hole expansion capacity. Therefore, the carbon content is 0.050% to 0.180% and preferably 0.070% to 0.160%.

In a preferred embodiment the carbon content is at most 0.140% and more preferably at most 0.130%, even more preferably at most 0.120% and most preferably at most 0.110% or even 0.100%

Si: 0.40 - 1.80, Al: 0.010 -1.00, Si+AI: 0.50 - 2.00%

Silicon (Si) is an element which is effective in increasing the strength of the steel by solute strengthening and effective in preventing carbon from precipitating in the form of iron carbides and increasing the stability of residual austenite. Si has a function of improving the ductility, work hardenability and stretch flangeability through the retardation of austenite grain growth during annealing. At least 0.40% Si is needed to achieve such effects. When the Si content is higher than 1.80%, surface quality of the steel sheet and the weldability of the steel are deteriorated. Accordingly, the Si content is 0.40 to 1.80%, preferably 0.50 to 1.50%, more preferably 0.90 to 1.35%.

The primary function of Aluminium (Al) is to deoxidise the liquid steel before casting. Furthermore, Al can partially replace Si as Al has a similar function as Si to prevent the formation of carbides and to stabilize the retained austenite and therefore is effective in improving strength and ductility balance. In order to have such effects, 0.010% or more of Al is needed. If the content of Al is too high, the ferrite to austenite transformation temperature Ac3 is increased to a level that is incompatible with conventional installations and will therefore not provide a microstructure according to the invention. In addition, the risk of cracking during casting and rolling increases as the Al content is increased. Therefore, the Al content is 0.010 to 1.00%, preferably 0.020 to 0.70%. When both Si and Al are added, the total amount of Si + Al is 0.50 to

2.00%. If the amount of aluminium is to be limited, for instance for (but not limited to) specific surface quality reasons, then the aluminium content is preferably between 0.020 and 0.090%, more preferably between 0.020 and 0.050%. Mn: 1.50% - 3.50%

Manganese (Mn) is an element which is indispensable for obtaining hardenability of the steel to achieve a desired strength. Mn is also added to balance the elevated phase transformation point as a result of high amounts of Si and Al. In order to produce such effects, the Mn content needs to be 1.50% or more. However, when the Mn content is greater than 3.50%, the elongation significantly deteriorates and formability is reduced due to the formation of macro-segregation and band microstructure. Therefore, the Mn content is 1.50% to 3.50% and preferably 1.80% to 3.00%.

P: 0.050% or less

Phosphorus (P) is an impurity in the steel. Since P is also effective for increasing the strength of steel and for increasing the stability of the RA. In order to produce such an effect, it is preferable that the P content be 0.002% or more. On the other hand, in the case where the P content is more than 0.050%, the steel is embrittled due to grain boundary segregation and, therefore, deteriorates stretch flangeability. Therefore, the P content is 0.002% to 0.050%. If excellent weldability is required, it is preferable that the P content be 0.020% or less, more preferably 0.015% or less, even more preferably 0.012% or less.

S: 0.020% or less

Sulphur (S) is present in the form of an inclusion such as MnS and deteriorates formability and/or weldability. Hence, the content thereof is preferably minimized. However, the manufacturing cost increases significantly if the S content is supressed below 0.001%. The content of S is preferably 0.010% or less for an optimal balance of the properties and the manufacturing cost.

N: 0.010% or less

Nitrogen (N) is inevitably present in the steel making process. N can be present in three nitrogen-related phenomena i.e. formation of blowholes, precipitation of one or more (carbo-)nitride compounds with elements such as Ti, Nb, Al and B and the interstitial solid element in solution. N in solid solution can markedly increases hardness and yield strength and decreases the tensile elongation. However, N deteriorates the toughness and formability of steel when an excess amount of 0.010% is added due to the formation of coarse nitride compounds or blowholes. Accordingly, nitrogen content in steel is to be 0.010% or less, preferably 0.0060% or less. The practical lower limit of nitrogen content is around 0.0005% because decreasing nitrogen content in steel further significantly increases production costs. A suitable and practical N content is between 0.0010 to 0.0050%.

One or more of the following elements can be optionally added.

Cr: 0.30% or less, Ni: 0.50% or less, Cu: 0.50% or less, Mo: 0.20% or less These elements are useful as strengthening elements for steel, and are also elements effective for stabilizing RA and ensuring a predetermined amount. In order for any of these elements to effectively individually exert such an effect the following amounts must at least be added Cr: 0.05% or more, Ni: 0.05% or more, Cu: 0.05% or more and Mo: 0.05%. When the content of the each of these elements is below the low limit, it is considered to be an impurity in the BOS-process. However, when each of these elements is added in an amount exceeding the upper limit, Cr > 0.30%, Ni > 0.50%, Cu > 0.50% and Mo > 0.20%, the above effects are saturated, which is economically wasteful. Preferably one, more or all of Cr, Ni, Cu and Mo are at most at impurity level.

Ti: 0.050% or less, Nb: 0.050% or less, V: 0.10% or less

At least one of these elements is an element that has effects of strengthening precipitation and refining the structure, and is useful for increasing strength. In order to effectively exhibit such an effect, it is recommended to add Ti: 0.005% or more, Nb: 0.005% or more, V 0.01% or more, respectively. However, if any of the elements is added in excess of the upper limit, the above effect is saturated, which is economically useless. When the content of the each of these elements is below the lower limit, it is considered to be an impurity in the BOS-process. Preferably one, more or all of Ti, Nb and V are at most at impurity level Ca: 30 ppm or less and / or REM: 30 ppm or less

Ca and REM (rare earth element) are elements that control the form of sulphide in steel and are effective for improving workability. Here, examples of the rare earth element used in the present invention include Sc, Y, and lanthanides. In order to effectively exert the above effects, it is recommended to add 3 ppm or more, respectively. However, even if it is added in excess of 30 ppm, the above effect is saturated, which is economically useless. It is more preferably at most 25 ppm. The same applies to the amount of Sc, Y, and lanthanides to be added if required.

B: 0.0050% or less

Although not essential to the invention Boron (B) is nevertheless a useful element in suppressing formation and growth of polygonal ferrite from austenite grain boundaries and therefore increases the hardenability of the steel. It helps to obtain a sufficient amount of BF and/or PM. To achieve such an effect at least 0.0003% B should be added. When B exceeds 0.0050%, the formability of the steel is deteriorated. Therefore, B, when added, is set to be in the range of 0.0003 to 0.0050%, preferably 0.0005 to 0.0020%, more preferably 0.0005 to 0.0015%. For B to be able to perform this role effectively, it is essential that free nitrogen is as low as possible so that the formation of BN is minimized. A suitable amount of Ti or Al should be added to scavenge nitrogen and to maximize the effect of B. According to a second aspect the invention is also embodied in a process according to the independent method claim 9.

A steel slab is prepared to have the preferred component composition described above based on a melt from the BOS-process. The slab is preferably manufactured by a continuous casting process for the purpose of preventing macro-segregation. These processes are not particularly restricted and may be carried out according to conventional methods. Thin slab casting, strip casting or the like can also be applied which, for the purpose of this invention, will also be considered a slab. In the case of strip casting it is acceptable for the manufacturing method to skip at least a part of the hot rolling process.

The slab is subjected to hot-rolling and then cold-rolling to obtain a cold-rolled steel strip. Preferable manufacturing conditions of a cold rolled steel strip include: (re)heating the slab to a temperature in the range of 1100 °C to 1300 °C; hot-rolling the slab to a hot-rolled strip wherein the finishing rolling temperature is at temperature equal to or higher than the Ar3 transformation point and preferably between 800 and 1050 °C; cooling at an average cooling rate of 20 °C/s or more (preferably 30 °C/s or more) to a temperature below 650 °C and coiling the hot-rolled strip at a temperature in the range of from 200 to 550 °C.

The coiled hot-rolled strip is descaled by pickling or the like prior to cold-rolling. The resulting steel sheet is preferably cold-rolled at a total cold rolling reduction of between 30 and 80%. If the rolling reduction of cold rolling is less than 30% then the recrystallization during annealing is incomplete so that a non-recrystallized ferrite phase is retained in a microstructure after the continuous annealing process, which may result in a decrease in formability. In addition, in the case where the rolling reduction of cold rolling is excessively high, since there is an increase in rolling loads during cold-rolling and rolling issues such as chattering and fracturing of strips may occur. Therefore, it is preferable that the rolling reduction of cold rolling is 80% or less.

The cold-rolled steel strip can be prepared in the form of a strip, sheet or blank for the ensuing heat treatment process. The cold rolled steel thus obtained is subjected to a thermo-mechanical treatment in the form of a hot press forming operation. The heat treatment includes an austenitizing process, a fast cooling process, an austempering or overaging process and a final cooling process. The heat treatment process is described with reference to the Fig. 1. The cold rolled steel sheet is first heated to a temperature T1 in the range of 680-

750 °C at a heating rate VI. The heating rate is not limiting, but to limit the length of this heating stage the heating rate is preferably at a rate of 10-25 °C/s, and then to the temperature T2, at a slow heating rate V2 of 0.5-10 °C/s, preferably at a heating rate V2 of 1-5 °C/s. In this second heating stage, the phase transformation of ferrite to austenite and recrystallization of ferrite occur concurrently. If V2 is too low, the resulting austenite grains may become too large, which may deteriorate the formability. if V2 is too high, recrystallization of ferrite is slow and results in a banded-like structure with small austenite grains and large austenite grains, and therefore promotes the formation of ferrite during cooling.

The cold rolled steel is annealed at temperature T2 in the austenite single-phase region for a time t2 to obtain a fully austenitic microstructure. T2 should be in a range of Ac3-30 °C to Ac3+50 °C, preferably Ac3-20 to Ac3+30 °C. It should be noted that Ac3 is usually tested during heating and is higher than Ae3. Therefore when isothermally holding a steel at Ac3-30 °C for a certain time, it is fully austenitized. The soaking time t2 is in a range from 1 to 300 seconds, preferably 5 to 200 seconds. Annealing at a temperature in the austenite single-phase region is necessary to avoid the formation of polygonal ferrite, which leads to a lower flange formability due to a lower hardness. If T2 is higher than Ac3+50 °C or t2 is longer than 300 seconds, austenite grains will grow, which influences the size and distribution of the retained austenite and also slow down the bainitic transformation kinetics later in the partitioning process. An excess amount of fresh martensite may form during final cooling as a result of this incomplete bainitic transformation, which leads to a higher strength but a low ductility and formability. If T2 is lower than Ac3-30 °C or t2 is shorter than Is, reverse transformation to austenite may not proceed sufficiently and/or carbides in the steel sheet may not be dissolved sufficiently. Accordingly, the annealing temperature needs to be equal to or higher than Ac3-30 °C, but should not exceed Ac3+50 °C. The annealing time t2 is 1 second to 300 seconds, preferably 5 to 200 seconds. After austenization at T2, the steel strip is first cooled at a slow cooling rate of V3 to a temperature T3. This slow cooling section is important for the current invention. To obtain steel sheets with high formability, ferrite formed in the high temperature range above Acl is avoided. Ferrite formed in the low temperature range below the Acl point during fast cooling is allowed. During this slow cooling, the C redistribution occurs in the austenite, leading to the formation of a metastable intermediate structure (MIS) including C-lean parts and the C-rich parts, which acts as a central connection among austenite, ferrite and cementite. The MIS is not a stable bulk phase, but acts as precursors of new phases. The formation of the MIS in the austenite phase will facilitate the formation of ferrite in the fast cooling stage and accelerate the bainitic transformation kinetics during austempering. Moreover, the temperature distribution in the steel sheet becomes more uniform, reducing the variation in the microstructure during fast cooling. Thus, the slow cooling has a significant effect of the microstructure although there is no ferrite formed. To make use of this slow cooling effectively, the T3 temperature should be higher than Ar3 to prevent the formation of high temperature ferrite during the slow cooling step. The upper limited of T3 is set to be T2-40 °C. T3 is preferably in the range of 680 to 850 °C, preferably 700 to 820 °C. The V3 does not have a significant effect on the microstructure and properties as long as T3 is above Ar3, preferably V3 is in the range of 0.5-15 °C/s, more preferably 1- 5 °C/s.

The steel sheet is then cooled to a temperature T4 ranging from Ms-200 °C to Bn, preferably from Ms-150 °C to Bn-20 °C at an average cooling rate of at least by Vc (°C/s), which is given by equation (6). When V4 is less than Vc °C/s, the microstructure specified is not obtained because an excess amount of ferrite is produced during cooling. Therefore, the sheet needs to be cooled at an average rate of Vc °C/s or more. The upper limit of the cooling rate V4 is not particularly restricted unless variation in temperature occurs in the steel sheet when the cooling is stopped. The upper limit of the average cooling rate is preferably 200° C/s or less because the shape of the steel sheet is distorted or it is difficult to control the ultimate cooling temperature, T4.

V _c = exp (7.1 - 9.18C - 1.22 Mn - l.OCr + 0.32Si + 1.124/) (6)

The temperature T4 should be in between Bn and Ms-200 °C, typically at Bn-20 °C and Ms-150 °C to initiate the formation of BF and/or martensite. If T4 is between Bn and Ms, BF is formed and if T4 is in between Ms and Ms-200 °C, some martensite is first formed. The formation of PM accelerates the bainitic transformation in the following step. Therefore, the amount of PM and BF in the final microstructure can be adjusted by changing T4. In general, the lower the T4 is, the more partitioned martensite but less bainitic ferrite forms. If T4 is below Ms-200 °C, the amount of untransformed austenite will be too low, thereby minimizing the TRIP effect and associated ductility of the obtained product. The temperature T4 should at most be Bn, preferably Bn-20 °C. Above Bn high temperature BF is obtained, which leads to lower strength due to the large BF grain size.

The steel sheet is then heated within a time t4 to a temperature T5 ranging from Ms-50 °C to Bs temperature, preferably from Ms to Bn and overaged at T5 for a period t5 ranging from 15 seconds to 150 seconds for partition process. The time t4 should preferably be controlled within 1 to 10s, preferably within 1 to 5s. Although the time t4 is not critical for the microstructural properties, limitation of the typical available production lines requires a short time at t4 such that sufficient time is left for the partition step to complete bainitic transformation and to stabilize the retained austenite. At T5, C partitioning occurs between the BF or martensite and untransformed austenite. The martensite transforms to partitioned martensite and the untransformed austenite continues to transform into carbide-free bainitic ferrite. The average carbon content in the retained austenite is increased as the time t5 is increased, so that retained austenite is made stable.

If T5 exceeds Bs, carbides may precipitate in the remaining austenite and the desired microstructure of steel cannot be obtained. Preferably, the T5 is below Bn temperature to obtain low temperature bainitic ferrite and or PM. If T5 is below Ms-50 °C, the degree of C partitioning is insufficient and the carbon concentration in retained austenite is not high enough to stabilize it in a limited time, which is a known constrain in typical available production lines. Accordingly, the T5 temperature should be above Ms-50 °C, preferably above Ms. The partition time t5 at T5 has to be long enough to allow the non-transformed austenite into bainitic ferrite and to allow C enrichment in the retained austenite but short enough to avoid the formation of carbides. When t5 is less than 15s, the partitioning of martensite is insufficient, the desired microstructure may not be obtained, and thus good formability of the steel sheet may not be sufficiently ensured. When t5 is longer than 150 seconds, carbides tend to precipitate in non-transformed austenite, which decreases the carbon content in the retained austenite. Accordingly, t5 is 15 to 150s, preferably 30 s to 100s. A shorter t5 is applied for a higher T5. It is not necessary that the holding temperature T5 be constant as long as the temperatures are within the ranges described. Actually, T5 might gradually increase due to the latent heat produced by bainitic transformation.

The steel is then cooled down to below 300 °C at a cooling rate V7 of at least 1 °C/s, preferably at least 5 °C/s, after which it is further cooled down to ambient temperature at V8, by either forced cooling or uncontrolled natural cooling. A suitable cooling rate V8 for an annealing production line is 5-30 °C/s. In a further embodiment of the invention, the steel product may further be subjected to a coating process as well known to a person skilled in the art, for example hot dip galvanizing, galvannealing or electrodeposition. The coating may be applied after cooling to ambient temperature or in between the process steps as described above.

T6 should be in the range from Bn to Bs, preferably in the range of 450 °C to 500 °C. The time t6 is preferably in the range of 1 to 30s. During this step, the overall microstructure transformation continues and the untransformed austenite continues to transform to bainitic ferrite. Therefore, the total time t5+t6 should be in the range of 15 to 150s, preferably in the range of 30 to 100s to limit the precipitation of carbides in non-transformed austenite and obtain the microstructure according to the invention. In the invention the heat treatment process as described above is incorporated in a hot forming press to form parts directly. In this application, blanks suitable for producing parts are prepared from cold rolled sheet of the invented compositions. A blank is first heated in an electric furnace at T2 for t2. After austenization, the blank is removed from the furnace and placed on the bottom die. In this transfer phase, the blank temperature decreases from T2 slowly to T3 due to heat losses to the environment. The hot press forming is then applied to form a part. During hot press forming, the blank temperature decreases at a rate V4 between 60 and 250 °C/s to T4, which can be controlled by initial die temperature. Finally, the formed part is moved to a furnace with a predetermined temperature T5 and holding for t5 and the cooling to ambient temperature in air. Zn coating can be applied by a separate process.

Examples

The invention will now be described by the following non-limiting examples.

Steels having compositions as shown in Table 1 were cast into 25 kg ingots of 200 mm x 110 mm x 110 mm in dimensions using vacuum induction. The following process schedule was used to manufacture cold rolled strips of 1 mm thickness:

• Reheating of the ingots at 1225 °C for 2 hours;

• Rough rolling of the ingots from 140 mm to 35 mm;

• Reheating of the rough-rolled ingots at 1200 °C for 30 min;

• Hot rolling from 35 mm to 4 mm in 6 passes;

• Run-out-table cooling: cool from finish rolling temperature (FRT) about 850 to 900 °C to 600 °C at a rate of 40 °C/s;

• Furnace cooling: strips transferred to a preheated furnace at 600 °C and then cooled to room temperature to simulate the coiling process;

• Pickling: The hot rolled strips were then pickled in HCI at 85 °C to remove the oxide layers;

• Cold rolling: The hot rolled strips were cold rolled to 1 mm strips;

• Heat treating according to the invention.

Cold rolled sheets with suitable size were used to simulate the annealing process in a production line by using a continuous annealing simulator. The process parameters are given in Table 2.

Samples for microstructure observations, tensile tests and hole expansion tests were machined from the heat-treated strips.

Dilatometry was done on the cold rolled samples of 10 mm x 5 mm x 1 mm dimensions (length along the rolling direction). Dilatation tests were conducted on a Bahr dilatometer type DIL 805. All measurements were carried out in accordance with SEP 1680. The critical phase transformation points Ac3, Ms and Mf were determined from the quenched dilatometry curves. Bs and Bn were predicted using JmatPro 10 software. The phase fractions during annealing for different process parameters were determined from dilatation curves simulating the annealing cycles. The microstructure was determined by optical microscopy (OM) and scanning electron microscopy (SEM) using a commercially available image-processing program. The cross section in the thickness direction parallel to the rolling direction of a steel sheet is polished then etched with a 3% nital solution or a 10% aqueous sodium metabisulfite solution. The microstructure is observed at a position located at 1/4 t (position at 1/4 of the thickness from the surface). The SEM used for the EBSD measurements is a Zeiss Ultra 55 machine equipped with a Field Emission Gun (FEG- SEM) and an EDAX PEGASUS XM 4 HIKARI EBSD system. The EBSD scans were captured using the TexSEM Laboratories (TSL) software OIM (Orientation Imaging Microscopy) Data Collection. The EBSD scans were evaluated with TSL OIM Analysis software. The EBSD scan area was in all cases 100 x 100 pm, with a step size of 0.1 pm, and a scan rate of approximately 80 frames per second.

The retained austenite (RA) was determined by X-ray diffraction (XRD) according to DIN EN 13925 on a D8 Discover GADDS (Bruker AXS). The XRD measurements were conducted on the subsurface at V¾ thickness of the steel sheet. The steel sheet is mechanically and chemically polished and is then analysed by measuring the integral intensity of each of the (200) plane, (220) plane, and (311) plane of fee iron and that of the (200) plane, (211) plane, and (220) plane of bcc iron with an X-ray diffractometer using Co-Ka radiation. The amount of RA and the lattice parameter in the RA were determined using Rietveld analysis. The C-content in the RA is calculated using the formula (D. Dyson and B. Holmes, Effect of alloying additions on the lattice parameter of austenite, J. Iron Steel Inst. 208 (1970) 469-474):

C =(a[A]-3.572-0.0012 Mn+0.00157 Si-0.0056 AI)/0.033 (7)

Where a is the lattice parameter of the retained austenite in angstrom, and C, Mn and Al are the content of the elements in wt.%.

Tensile tests - JIS5 test pieces (gauge length = 50 mm; width = 25 mm) were machined from the annealed strips such that the tensile direction was parallel to the rolling direction. Room temperature tensile tests were performed in a Schenk TREBEL testing machine following NEN-EN10002-1:2001 standard to determine tensile properties (yield strength YS (MPa), ultimate tensile strength UTS (MPa), total elongation TE (%)). For each condition, three tensile tests were performed, and the average values of mechanical properties are reported.

Hole Expansion Test (Stretch Flangeability Evaluation Test) - Test pieces for testing hole expandability (size: 90x90 mm) were sampled from the obtained rolled strip. In accordance with The Japan Iron and Steel Federation Standards JFS T 1001, a 10 mm diameter punch hole was punched in the centre of the test piece and a 60° conical punch was pushed up and inserted into the hole. When a crack penetrated the strip thickness, the hole diameter d (mm) was measured. The hole expansion ratio l (%) was calculated by: A(%) = {(d-d0)/d0}xl00, with dO = 10 mm.

The process parameters are presented in Table 2 using the indications in Figure 1. The Ac3, Ar3, Bs Ms, and the critical cooling rate Vc are also given. The Ac3 is used for determining T2, Ar3 for T3, Bs and Ms for T4 and T5, Vc for V4. The resulting microstructures and properties are given in Table 3. High HEC values are obtained when the amount of ferrite is well controlled, which is adjusted by selecting T3 and V4. Examples P7, P8, P13, P14, P20, P24, P28 and P33 are out of the range of the current invention for T3 and/or V4, which lead to an excessive amount of ferrite or retained austenite, thus the required high formability cannot be ensured. Comparative experiment were performed with steel 98, which has a carbon content outside the claimed range. Table 2 shows that and processes P34 and P35 do provide high strength and elongation values but the HEC values are much too low due to the large amounts of retained austenite resulting from the high C content of the steel. Heat treatment in combination with hot press forming

Steel blanks with dimensions of 200 mm x 110 mm x 1.0 mm have been prepared from cold-rolled steel sheets having the composition 58A as given in Table 1. The hot forming was performed in a Schuler SMG company press (hereafter SMG press).

The blanks were first heated at 890 °C (T2) in a box furnace for 4 min and then transferred to the SMG press. The blanks were transferred to the SMG press in 10s (T3 = 780 °C) and hot-press formed into a U-shaped article and cooled in the die to T4. The SMG press tools were preheated to a temperature of 400 °C (T4). The cooling rate during press forming (V4) is about 65 °C/s. The article was transferred to a muffle furnace preheated at a temperature of 400 °C and held for 120s and then taken out of the furnace and cooled in air. JIS5 tensile specimens were prepared from the flat part of the hot formed articles. The results obtained from SMG samples are shown in Table 4.

Table 1. Chemical compositions of the steels

*Alloy 98 is comparative example.

Table 2. The annealing process parameters for continuous annealing simulations

Table 3. The microstructures and properties of the continuously annealed sheets

Table 4. The microstructures and properties of the hot press formed sheets