WO2013007729A1

WO2013007729A1 - Hot-rolled high-strength steel strip with improved haz-softening resistance and method of producing said steel

Info

Publication number: WO2013007729A1
Application number: PCT/EP2012/063515
Authority: WO
Inventors: David Crowther; Willem Maarten VAN HAAFTEN
Original assignee: Tata Steel Ijmuiden Bv
Priority date: 2011-07-10
Filing date: 2012-07-10
Publication date: 2013-01-17
Also published as: CN103649355A; EP2729590A1; CN103649355B; EP2729590B1

Abstract

Hot-rolled high-strength micro-alloyed steel strip with improved HAZ-softening resistance having a microstructure comprising martensite, tempered martensite and/or bainite.

Description

Hot-rol led hig h-strength steel strip with improved HAZ-softening resistance and method of producing said steel

The invention rel ates to a hot-rolled high-strength steel strip with improved HAZ-softening resistance and method of producing said steel

Traditionally steels with a high yield strength have been made by austenitising and q uench ing , but by th is tech n iq ue, for insta nce optimal surface quality and impact toughness may not be achieved . Manufacturing expenses have also been high.

EP1375694 describes a method for prod ucing a h ig h-strength, high- toughness steel with good workability and weldability by means of hot rolling . However, the inventors have found that these steels are susceptible to HAZ- softening. The heat-affected zone (HAZ) is the area of base material which has had its microstructure and properties altered by welding or heat intensive cutting operations. The heat from the welding process and subsequent re- cooling causes this change in the area surrounding the weld . The heat-affected zone which forms adjacent to the weld in steels is one of the most common regions of weld failure. Modern thermo-mechanically rolled (TMCP) steels, because of their lean chemistry, and in particular the low carbon content, require carefully controlled welding parameters in order to achieve adequate strength in the weld HAZ. For conventional low-carbon steels and some TMCP pipeline steels such as X70 and X80, HAZ softening is often the result of high heat-input welding procedure because of the slow heat dissipation in the HAZ. For higher grade TMCP steels such as X100, HAZ softening can happen even under moderate welding heat-input because of the base metal's ultra fine grain size and its bainite- and martensite-dominated microstructure. This reduction in hardness and strength in HAZ makes it a weak point in a welded pipeline structure. Consequently it is important to limit the degree to which this softening takes place, given the nature of the base metal and the welding conditions. The areas wh ich have received most study i n rega rds to HAZ softening are high strength plate steels e.g . for linepipe applications and, more recently, AHSS for automotive applications.

WO2007/129676 relates to a hot pressed steel member made from a hig h carbon cold-rolled steel sheet which is austenitised, hot-pressed to produce a steel member and quenched to achieve a minimum tensile strength of 1.8 GPa. The hot-rolled steel that is to be cold rolled contains at least 50% ferrite.

EP2028284 discloses a seamless steel pipe and EP1662014 discloses a low carbon hot rolled steel plate which is coiled between 450 to 650°C and subsequently promptly reheated to between 550 and 750°C to achieve a three-phase ferrite, bainite and island martensite structure.

The object of the present invention is to achieve a high-strength hot- rolled steel strip that is less susceptible to HAZ-softening than the currently available high strength hot-rolled steel strip.

In a first aspect this object is reached by a hot-rolled high-strength micro- alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:

0.07 - 0.30% C;

0.8 - 2.0% Mn;

0.01-0.08% A ;

0.2 - 1.5% Cr;

· 0.1-0.7% Mo;

0.0005-0.005 B;

0.01-0.07% Nb;

at most 0.5% Si;

at most 0.03% P;

· at most 0.015% S;

at most 0.05% Ti;

at most 0.1% V;

at most 0.2% Cu;

at most 0.2% Ni;

· at most 0.008% N;

optionally calcium additions for sulphide shape control, at most 0.015%; other elements in amounts of impurity level, balance iron;

the yield strength of the steel strip being at least 960 MPa. All compositional percentages are given as weight percent, unless indicated otherwise.

The inventive idea is based on the fact that by selecting the combination of chemical elements in the amounts prescribed and in particular the presence of niobium and molybdenum good hardening can be maintained . The steel structure is not critical for the segregation of manganese and carbon during the casting process owing to the low manganese and carbon content. The steel properties are not critical for local fluctuations of the coiling temperature in the strip, which facilitates the steel production and has an advantageous effect in the homogeneity of its mechanical properties, which again has a positive influence both in the flatness of the end product and in the residual stress. The steel sheet is highly suitable for welding and laser cutting, and at the same time it has a good fatigue strength irrespective of said thermal treatments. Further, the steel sheet has excellent bending properties, a good impact toughness as well as a good resistance to softening in tempering. The presence of the precipitating elements in sol ution which are subsequently available for precipitation during the cooling of the HAZ after welding ensures a significant improvement of the HAZ-softening resistance.

The steel according to the invention can be thermally cut, for instance by laser, into precisely defined shapes. It has been observed that a remarkably smooth cutting surface is achieved in a laser cut object. On the other hand, it has been found out that the strength difference between the basic material and the soft zone created in the technical cutting process, which zone is located in the immediate vicinity of the hardened zone, is small or absent also as a result of the presence of the precipitating elements in solution which are subsequently available for precipitation during the cooling of after thermal cutting . These together have an advantageous affect in the fatigue strength . In addition, the relatively low carbon content reduces the peak hardness of the hardened zone, so that the cutting surface is not sensitive to embrittlement and cracking, neither in the working of the object nor in practical use.

Carbon is an important element in controlling the strength, but needs to be limited to some extent to give a good balance in toughness, weldability and formability. The C content is kept relatively low between 0.07 and 0.30% to achieve a good toughness (all percentages in compositions are expressed in weight percent). In combination with a low temperature coiling temperature the microstructures will contain martensite and/or bainite. The exact amounts strongly depend on the composition, cooling rate on the run-out table and the coiling temperature. At relatively low C contents, the Ms temperature will be quite high, so the martensite will auto temper to some extent. Depending on the strength requirement suitable carbon-windows were found . For a steel having a yield strength of at least 960 MPa the carbon content is preferably at least 0.07 and/or at most 0.13%. For a steel having a yield strength of at least 1 100 M Pa the carbon content is preferably at least 0.13 and/or at most 0.18%, and for a steel having a yield strength of at least 1300 MPa the carbon content is preferably at least 0.19%, more preferably at least between 0.23 and/or at most 0.30%.

A suitable maximum carbon content is 0.27% or even 0.25%.

Elements like manganese, chromium, molybdenum and boron provide hardenability to promote the formation of bainite and/or martensite. The manganese content is limited to between 0.8 and 2.0%. When the manganese exceeds the upper boundary, the risk of segregation becomes significant and this may adversely affect the homogeneity of the microstructure. At levels below 0.8% the effect on hardenability is insufficient. A suitable minimum manganese level is 1.1%.

Boron is added to promote the hardenability. It is important to avoid the formation of boron nitrides as this wil l render the boron ineffective for the promotion of the hardenabil ity. The role of titanium in the composition according to the invention is to protect the boron because Ti forms titanium nitrides and as a consequence no BN is formed . The amount of alloyed boron is at least 0.0005% B (i.e. 5 ppm) but no more than 0.005% B (i.e. 50 ppm) in order to reduce grain size and to increase the hardenability. If titanium is added as an alloying element, the amount of titanium is typically at least 0.01% Ti but no more than 0.05% in order to bind the nitrogen N and to prevent the creation of boron nitrides BN. An alternative within the scope of the invention is to use aluminium to bind the nitrogen and thereby protect the boron . In cases where the formation of TiN-particles is deemed undesirable, e.g . because of their effect on Charpy-toughness, protecting the free boron by aluminium may be the preferred option. Niobium partly precipitates in the austenite during hot-rolling, thereby contributing to the grain refinement of the final transformed microstructure by the retardation of the recrystallisation of austenite. In addition, the niobium remaining in sol ution at transformation has a powerful effect on red ucing transformation temperatures, especially at faster cooling rates, so it is also beneficial for ha rdena bi l ity . At l ow coil i ng tem peratu res ( < 500°C) the contribution of niobium by precipitation strengthening is expected to be small because the temperature is too low for precipitation of fine NbC. A suitable Nb content is at least 0.02%, preferably at least 0.025%.

V has a similar but less powerful effect as Nb in this case. However, the main reason for the addition of N b and V is to improve the HAZ-softening resistance. In relevant parts of the HAZ, the thermal cycle is su ch that temperatures are reached which will allow precipitation strengthening by Nb and V, thus causing an increase in hardness as a result of the precipitation of elements which were kept in solution by the low coiling temperature. The major contribution is believed to be made by Nb and V carbides, nitrides or carbo-nitrides. To a lesser extent it is believed that MoC precipitates may form having a similar effect. If present, a suitable minimum V-content is 0.04%.

The steels according to the invention are aluminium-killed or aluminium- silicon killed steels in order to reduce the oxygen content to a minimum so that no reaction occurs between carbon and oxygen during solidification. The amounts of aluminium added to the steel during production therefore include those needed for deoxidation . The remaining amount in the end product, also called soluble aluminium (Al_SOi) is between 0.01 and 0.08% Al. In the context of this invention the aluminium content referred to is soluble aluminium.

Silicon may also serve as a deoxidant in the steel of the present invention in addition to aluminium. It also acts as a solid solution hardener starting from about at least 0. 10% Si and up to 0.50 % Si, which has an advantageous effect on the impact toug h ness and workabil ity. Above 0.5% the sil icon adversely affects the surface quality of the steel to an unacceptable extent and the removal of the hot rolling scale by pickling becomes increasingly difficult with increasing silicon content.

Phosphorus P contained as an impurity should be at most 0.03%, and sulphur S should be even lower and should be limited to at most 0.015%, which means that these contents are restricted in order to achieve good impact toughness and bendability. When necessary, further properties can be improved by treating the melt with cored wire containing Ca-Si or Ca-Fe(Ni). During calcium treatment, the alumina and silica inclusions are converted to molten calcium aluminates and silicate which are globular in shape because of the surface tension effect. The calcium aluminates inclusions retained in liquid steel suppress the formation of MnS stringers during solidification of steel . This change in the composition and mode of precipitation of sulphide incl usion during solid ification of steel is known as sulphide morphology or sul phide shape control . This results in less nozzle clogg ing d u ring casting , better mechan ica l properties as the l ong stretched M nS stringers act as crack initiating points. Typical amounts of calcium in the steel for sulphide shape control is 0.0015 to 0.015% Ca. A suitable maximum is 0.005% Ca.

Chromium should be between 0 . 2 a n d 1.5%. Molybdenum should preferably be between 0.1% Mo and 0.7% Mo. Both elements are added in o rder to i n crease h a rd en i n g and tempering resistance. This enables precipitation at higher coiling temperatures, which can be used for decreasing and even preventing the softening of the steel , as well as for al leviating strength fl uctuations caused by local temperatu re d ifferences during the cooling of the coil. A suitable minimum molybdenum content is 0.15%.

Alloying with elements like copper and nickel, often used in steels of this strength level are preferably avoided in view of the surface issues associated with copper. As copper is often alloyed in conjunction with nickel to alleviate the adverse effects of copper this is also not needed . So nickel and/or copper are preferably present at most at impurity level or more preferably completely absent.

Th e m icrostru ctu re of the stee l s a cco rd i n g to the i nve nti o n a re characterised as a microstructure that consists predominantly of tempered martensite, characterised by small carbides in a Widmanstatten pattern, and/or bainite. Ideally the microstructure is free from ferrite and pearlite constituents as these will deleteriously affect the strength level to be reached . In practice it may be unavoidable that some minor patches of ferrite are present, but the amount may not exceed the level where the strength levels is significantly affected . The abovementioned deleterious and subsequently undesirable ferrite constituents which form at h ig h tra nsfo rmatio n temperatures should be clearly distinguished from the ferritic part of bainite or Widmanstatten ferrite or acicular ferrite which form at low transformation temperatures. The former constituents are undesirable, the latter are not.

The hot-rolled steel strip according to the invention that is directly hot- rolled to the thickness 2 mm - 16 mm can be manufactured as wear-resistant and with different minimum yield strength . Typical threshold values in the marketplace are 960, 1100 and 1300 M Pa, on ly by chang ing the analysis and/or the post-rolling cooling rate of the strip, and/or temperature before the coiling, within the scope of the invention. This kind of high yield strength steel can also be used in targets where the structures require properties typically demanded of structural steel, such as good workability, weldability and impact toug hness, wh ich means that the hot-rolled steel strip according to the invention is feasible also as weldable structural steel. In the steel analysis to be ex p l a i n ed i n th e s pecifi cati o n be l ow, a l l content percentages are percentages by weight, and the rest of the steel that is otherwise not defined is naturally iron, Fe, and unavoidable impurities. The value of 960 MPa can be reached even to thickness values of the hot strip of up to 16 mm . For 1100 and 1300 MPa the values can be reached for val ues of up to 12 mm . The higher thickness results in a somewhat lower cooling rate, and therefore in a less enriched austenite prior to the phase transformation . This leads to a somewhat lower strength level, meaning that with these lean compositions the SHOO or S1300 level cannot be obtained for thicknesses of the hot strip over 12 m m . Preferably the minimum thickness is 3 mm and/or the maximum thickness is 10 mm. It should be noted that the val ues of the strength as defined in this invention are measured in the longitudinal direction (i.e. the tensile specimen is taken in the longitudinal direction of the strip (the direction of movement through the rolling mill)). Values in the transverse direction (i.e. the tensile specimen is taken in the width direction of the strip) may be different from the values in the longitudinal direction, and are usually higher than those in the longitudinal direction for strength and lower for elongation.

In an embodiment the carbon content of the steel is between 0.07 and 0.13% and the yield strength is at least 960 MPa. In an embodiment the carbon content of the steel is between 0.13 and 0.18% and the yield strength is at least 1100 MPa.

In an embod iment wherein the carbon content is at least 0.19%, preferably between 0.23 and 0.30%, and the yield strength is at least 1300 MPa.

A suitable maximum tensile strength of the hot-rolled steel according to the invention is 1700 MPa.

According to a second aspect the invention is embodied in a method for manufacturing a hot-rolled high-strength micro-alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance and a yield strength of at least 960 MPa having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:

0.07 - 0.30% C;

· 0.8 - 2.0% Mn;

0.01-0.08% A ;

0.2 - 1.5% Cr;

0.1-0.7% Mo;

0.0005-0.005 B;

· 0.01-0.07% Nb;

at most 0.5% Si;

at most 0.03% P;

at most 0.015% S;

at most 0.05% Ti;

· at most 0.1% V;

at most 0.2% Cu;

at most 0.2% Ni;

at most 0.008% N;

optionally calcium additions for sulphide shape control, at most 0.015%; · other elements in amounts of impurity level, balance iron;

the strip being finish hot-rolled above the Ar₃-temperature, wherein the method includes at least the following steps:

• finish rolling to a final thickness of from 2 to 16 mm • cooling the hot rolled strip within at most 10 seconds from the last hot rolling pass to a coiling temperature of between 20 and 500°C at a cooling rate sufficient to transform the rolled microstructure into a microstructure comprising martensite and/or bainite.

The hot-rolling process is the conventional hot-rolling process, either starting from a slab having a thickness of between 150 to 350 m m, i .e . conventionally continuously cast slabs, or below 150, i.e. thin slab casting or even strip casting . Finish rolling is preferably while the steel is still austenitic to provide the steel with a fine grain and thereby good impact toughness. The coiling temperature is preferably low to achieve the desired mechanical p ro pe rties . Prefe ra b l y th e co i l i n g te m pe ratu re is be l ow 400 °C . By manufacturing this type of steel by fast cooling or quenching directly after hot rol l i n g , a n excel lent i m pact to ug h ness is o bta i ned beca use the phase transformation into martensite and/or bainite takes place from a fine-grained, worked austenite. It also improves the surface quality because the primary scale is removed in a descaler prior to the rolling . Moreover, there is no need for a very expensive lengthy additional annealing treatment to dissolve all precipitates. In a hot strip rolling line, the slabs to be rolled are reheated in a reheating furnace to a temperature range between 1100 to 1250°C, and held for several hours. In that case the dissolution of special carbides, such as Cr and Mo carbides, and the homogenization of the structure is as complete as possible. On the other hand, the growing of the austenite grain at the high heating temperature does not make the end product more brittle, because the austenite grains are refined first by recrystallisation during high temperature rolling in the initial stages of the rolling process and by transformation of the deformed austenite grain formed as a result of the retardation of austenite transformation during thermo-mechanical rolling in the last stages of the hot rolling process. The heavily deformed austenite grains transform into a very fine transformation product during cooling on the run-out table. This results in a h i g h y i e l d s stress , combined with an excellent impact toughness. Manufacturing costs and production time can be further reduced if a thin slab casting and direct rolling facility is used where the elements like Cr and Mo carbides have not even yet precipitated before the rolling starts. Accord ing to the i nvention , steel is ma n ufactu red at a fina l rol l i ng temperature at which the steel is still austenitic, i.e. above Ar3 to a final hot rolled thickness. The cooling of the strip begins no later than 10 seconds after the last hot rol l ing pass, and it is cooled sufficiently rapidly to al low the austenite to transform into a bainitic and/or martensitic microstructure, the cooling rate preferably being at least 30°C/s, down to a coiling temperature in the range 20°C - 500°C, preferably down to a coiling temperature in the range 20°C - 450°C. The obtained result is typical ly a nearly completely bainitic and/or martensitic microstructure, so that the bainite and/or martensite content preferably is at least 90 % by volume, preferably at least 95%. The microstructure is also preferably free from ferrite formed at high temperatures and free from pearlite constituents, thus rendering the microstructure substantially fully bainitic/martensitic, where Widmannstatten fe rrite o r acicular ferrite is considered to be a bainitic structure for this purpose. In the coiling temperature range of below 100°C the martensite is not tempered, although some auto-tempering may occur at these low temperatures for the low carbon grades, whereas when the coiling temperature is at least 100°C, the martensite is tempered . At temperatures above 200°C the martensite is tempered and the carbon precipitated.

In a preferable embodiment the coiling temperature is at most 450°C, more preferably 425°C, for steels containing a carbon content of at most 0.12% or at most 275°C, more preferably 250 or 225°C, for steels containing a carbon content of at least 0.13%.

In an embodiment the cooling rate during the accelerated cooling after hot rolling and before coiling is between 5 and 100°C/s.

In an embodiment the finish rolling temperature is above Ar₃ and also below 920°C, and preferably also below 900°C.

According to a third aspect the steel according to the invention is used in the production of a part for an automobile, a lorry, ship-building, construction work, heavy haul equipment, earth-moving equipment or mobile cranes. The invention is now further explained by means of the following, non limiting examples. Table 1 shows com positions of steels for each of the categories S700, S960, SHOO and S1300.

Table 1. Composition of the steels (all in weight% x1000, except B and N (ppm).AII

steels are Ca-treated, imp=impurity level)

Type ID C Si Mn Al P S Cr Mo B N Ti V Nb

S960

High Mn 4587 100 250 1770 30 14 5 260 250 20 48 33 imp 39

High Mn 4588 91 240 1780 34 15 5 500 5 20 60 35 imp 39

Low Mn 4589 110 220 1210 31 15 4 490 510 20 50 30 imp 38

Low Mn 4590 110 240 1180 32 15 4 1030 5 20 80 32 imp 39

Low Si 4591 90 100 1600 30 14 5 500 250 20 50 30 imp 40

Low Si m 94 108 1590 38 11 2 502 263 25 52 27 47 40

Base 4699 99 230 1490 33 15 4 510 250 30 60 29 imp 40

+V 4671 94 240 1510 33 15 3 530 250 25 60 31 51 39

-Nb 4670 97 230 1500 35 15 4 540 250 25 60 30 imp imp

+Cr-Nb 4673 110 240 910 32 15 4 1130 300 30 60 29 imp imp

SHOO

B 4814 155 140 1460 23 15 5 490 240 30 60 26 imp 29

+c 4815 170 140 1460 22 15 5 490 240 30 60 29 imp 29

+ Mo 4816 165 140 1470 22 15 5 490 510 30 60 30 imp 30

-Cr 4817 155 140 1460 21 15 5 260 240 30 60 29 12 19

+ Mo-Cr 4818 155 140 1460 22 15 5 250 500 30 60 28 imp 29

-Ti+AI 4819 145 140 1460 59 15 5 450 240 30 60 6 imp 29

+V 4820 155 150 1460 23 15 5 490 240 20 60 29 50 29 -Ti+AI+V 4819* 145 140 1460 59 15 5 450 240 30 60 6 50 29

S1300

Base 250 100 1490 27 14 3 480 510 20 60 3 imp 29

+ B 230 100 1510 27 14 4 490 510 16 60 2 imp 30

+AI 250 110 1500 64 14 4 480 510 18 60 2 imp 28

+Ti 270 110 1510 27 14 4 490 510 20 60 25 imp 30

SEM microstructures show that all hot-rolled steel samples have a microstructure that consists predominantly of tempered martensite, characterised by small carbides in a Widmanstatten pattern. The tempering of the martensite will have resulted due to the high M_s temperature of these relatively low C steels, and also due to the slow cooling to simulate coil cooling . The pattern of the carbides, showing several variants of an orientation relationship between the carbide and the matrix, is characteristic of tempered martensite. The prior austenite grain boundaries were also visible, and showed some elongation along the rolling direction, as was observed using optical metallography. The strength levels depend strongly on the coiling temperature (see figure 1 for the SHOO materials and figure 2 for the S1300 material) which is to be understood in view of the desired microstructure that consists predominantly of tempered martensite.

Figure 3 shows the effect of HAZ-softening on the Y-axis for steels (a, c) according to the invention and for comparative steels (b, d). It is clearly visible that the inventive steels outperform the comparative steels in terms of a reduced softening of the HAZ. Steels b and d are S960 -Nb and S960 +Cr-Nb respectively (see table 1) and a and c are S960 Base and S960 +V respectively.

Claims

1. Hot-rolled high-strength micro-alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:

0.07 - 0.30% C;

0.8 - 2.0% Mn;

0.01-0.08% A ;

0.2 - 1.5% Cr;

0.1-0.7% Mo;

0.0005-0.005 B;

0.01-0.07% Nb;

• at most 0.5% Si;

at most 0.03% P;

at most 0.015% S;

• at most 0.05% Ti;

at most 0.1% V;

• at most 0.2% Cu;

• at most 0.2% Ni;

• at most 0.008% N;

• optionally calcium additions for sulphide shape control, at most 0.015%;

• other elements in amounts of impurity level, balance iron;

the yield strength of the steel strip being at least 960 MPa.

2. Steel according to any one of the preceding claims having a chromium content of at least 0.2%.

3. Steel according to any one of the preceding claims having a nickel content at impurity level.

4. Steel according to any one of the preceding claims having a copper content at impurity level.

5. Steel according to any one of the preceding claims wherein the carbon content is between 0.07 and 0.13% and having a yield strength of at least 960 MPa. Steel according to any one of claims 1 to 4 wherein the carbon content is between 0.13 and 0.18% and having a yield strength of at least 1100 MPa.

Steel according to any one of claims 1 to 4 wherein the carbon content is at least 0.19%, preferably between 0.23 and 0.30%, and having a yield strength of at least 1300 MPa.

A method for manufacturing a hot-rolled high-strength micro-alloyed steel strip having a thickness of between 2 and 16 mm with improved HAZ-softening resistance and a yield strength of at least 960 MPa having a microstructure comprising martensite, tempered martensite and/or bainite, and where the steel contains, in percentages by weight:

0.07 - 0.30% C;

0.8 - 2.0% Mn;

0.01-0.08% A ;

0.2 - 1.5% Cr;

0.1-0.7% Mo;

0.0005-0.005 B;

0.01-0.07% Nb;

• at most 0.5% Si;

at most 0.03% P;

at most 0.015% S;

• at most 0.05% Ti;

at most 0.1% V;

• at most 0.2% Cu;

• at most 0.2% Ni;

• at most 0.008% N;

• optionally calcium additions for sulphide shape control, at most 0.015%;

• other elements in amounts of impurity level, balance iron;

• finish rolling to a final thickness of from 2 to 16 mm

• cooling the hot rolled strip within at most 10 seconds from the last hot rolling pass to a coiling temperature of between 20 and 500°C at a cooling rate sufficient to transform the rolled microstructure into a microstructure comprising martensite and/or bainite.

Method according to claim 8 wherein the coiling temperature is:

• at most 450°C for steels containing a carbon content of at most 0.12% or

• at most 275°C for steels containing a carbon content of at least 0.13%.

Method according to any one of claims 8 to 9 wherein the cooling rate during the accelerated cooling after hot rolling and before coiling is between 5 and 100°C/s.

Method according to any one of claims 8 to 10 wherein the finish rolling temperature is above Ar₃ and also below 920°C, and preferably also below 900°C.

Use of the steel according to claim 1 to 7 in the production of a part for an automobile, a lorry, heavy haul equipment, earth-moving equipment or mobile cranes.