TWIP steel sheet having an austenitic matrix

The present invention relates to a cold-rolled and recovered TWIP steel sheet having an austenitic matrix and a method for the manufacture of this cold- rolled and recovered TWIP steel. The invention is particularly well suited for the manufacture of automotive vehicles.

With a view of saving the weight of vehicles, it is known to use high strength steels for the manufacture of automobile vehicle. For example for the manufacture of structural parts, mechanical properties of such steels have to be improved. However, even if the strength of the steel is improved, the elongation and therefore the formability of high steels decreased. In order to overcome these problems, twinning induced plasticity steels (TWIP steels) having good formability have appeared. Even if these products show a very good formability, mechanical properties such as Ultimate tensile strength (UTS) and yield stress (YS) may not be high enough to fulfill automotive application.

The patent application US2006278309 discloses a hot-rolled austenitic iron/carbon/manganese steel sheet, the strength of which is greater than 900 MPa, the product (strength (in MPa) ^* elongation at fracture (in %)) of which is greater than 45000 and the chemical composition of which comprises, the contents being expressed by weight: 0.5%<C<0.7%, 17%<Mn<24%, Si<3%, Al<0.050%, S<0.030%, P<0.080%, N<0.1 %, and, optionally, one or more elements such that: Cr<1 %, Mo<0.40%, Ni<1 %, Cu<5%, Ti<0.50%, Nb<0.50% and V<0.50%, the composition further comprising iron and inevitable impurities resulting from the smelting, the recrystallized fraction of the steel being greater than 75%, the surface fraction of precipitated carbides of the steel being less than 1.5% and the mean grain size of the steel being less than 18pm.

However, the strength of this austenitic steel sheet is really low. Indeed, in the examples, the strength is of 1130MPa in the range of the invention.

Thus, the object of the invention is to solve the above drawbacks by providing a TWIP steel having a high strength, an excellent formability and elongation. It aims to make available also an easy to implement method in order to obtain this TWIP steel. This object is achieved by providing a TWIP steel sheet according to claim 1. The steel sheet can also comprise characteristics of claims 2 to 12.

Another object is achieved by providing a method for producing a TWIP steel sheet according to claim 13. The method can also comprise characteristic of claims 14 to 16.

Other characteristics and advantages of the invention will become apparent from the following detailed description of the invention.

The following terms will be defined:

- all percentage "%" in the steel composition are defined by weight,

- UTS: ultimate tensile strength (MPa) and

- TE: total elongation (%).

The invention relates to a cold-rolled and recovered TWIP steel sheet having an austenitic matrix comprising by weight:

0.71 < C < 1.20%,

13.0 < Mn < 25.0%,

S < 0.030%,

P < 0.080%,

N < 0.10%,

0.1 < Si < 3.0%,

0.1 < V < 2.50%,

and on a purely optional basis, one or more elements such as

Cu < 5.0%,

Al < 4.0%,

Nb < 0.50%,

B < 0.0050%,

Cr < 1.0%,

Mo < 0.40%,

Ni < 1.0%,

Ti < 0.50%,

0.06 < Sn < 0.2%,

the remainder of the composition being made of iron and inevitable impurities resulting from the elaboration. Without willing to be bound by any theory, it seems that the TWIP steel sheet according to the invention allows for an improvement of the mechanical properties thanks to this specific composition. Indeed, it is believed that the above composition comprising the high amount of C allows for an improvement of, among others, ultimate tensile strength.

Regarding the chemical composition of the steel, C plays an important role in the formation of the microstructure and the mechanical properties. It increases the stacking fault energy and promotes stability of the austenitic phase. When combined with a Mn content ranging from 13.0 to 25.0% by weight. In case there are vanadium carbides, a high Mn content may increase the solubility of vanadium carbide (VC) in austenite. However, for a C content above 1.2%, there is a risk that the ductility decreases due to for example an excessive precipitation of (Fe,Mn) ₃ C cementite. Preferably, the carbon content is between 0.71 and 1.1 %, more preferably between 0.8 and 1.0% and advantageously between 0.9 and 1.0% by weight so as to obtain sufficient strength combined optionally with optimum carbide or carbonitride precipitation.

Mn is also an essential element for increasing the strength, for increasing the stacking fault energy and for stabilizing the austenitic phase. If its content is less than 13.0%, there is a risk of martensitic phases forming, which very appreciably reduce the deformability. Moreover, when the manganese content is greater than 25.0%, formation of twins is suppressed, and accordingly, although the strength increases, the ductility at room temperature is degraded. Preferably, the manganese content is between 15.0 and 24.0%, more preferably between 17.0 and 24.0% so as to optimize the stacking fault energy and to prevent the formation of martensite under the effect of a deformation. Moreover, when the Mn content is greater than 24.0%, the mode of deformation by twinning is less favored than the mode of deformation by perfect dislocation glide.

Al is a particularly effective element for the deoxidation of steel. Like C, it increases the stacking fault energy which reduces the risk of forming deformation martensite, thereby improving ductility and delayed fracture resistance. However, Al is a drawback if it is present in excess in steels having a high Mn content, because Mn increases the solubility of nitrogen in liquid iron. If an excessively large amount of Al is present in the steel, the N, which combines with Al, precipitates in the form of aluminum nitrides (AIN) that impede the migration of grain boundaries during hot conversion and very appreciably increases the risk of cracks appearing in continuous casting. In addition, as will be explained later, a sufficient amount of N must be available in order to form fine precipitates, essentially carbonitrides. Preferably, the Al content is below or equal to 2%. When the Al content is greater than 4.0%, there is a risk that the formation of twins is suppressed decreasing the ductility. Preferably, the amount of Al is above 0.1 %.

Correspondingly, the nitrogen content must be 0.1 % or less so as to prevent the precipitation of AIN and the formation of volume defects (blisters) during solidification. In addition, when elements capable of precipitating in the form of nitrides are present, such as vanadium, niobium, titanium, chromium, molybdenum and boron, the nitrogen content must not exceed 0.1%.

According to the present invention, the amount of V is between 0.1 and 2.5%, preferably between 0.1 and 1.0%. Preferably, V forms precipitates. Advantageously, vanadium elements have a mean size below 7 nm, preferably between 0.2 and 5nm and are intragranular in the microstructure.

Silicon is also an effective element for deoxidizing steel and for solid-phase hardening. However, above a content of 3%, it reduces the elongation and tends to form undesirable oxides during certain assembly processes, and it must therefore be kept below this limit. Preferably, the content of silicon is below or equal to 0.6%.

Sulfur and phosphorus are impurities that embrittle the grain boundaries. Their respective contents must not exceed 0.030 and 0.080% so as to maintain sufficient hot ductility.

Some Boron may be added up to 0.005%, preferably up to 0.001 %. This element segregates at the grain boundaries and increases their cohesion. Without intending to be bound to a theory, it is believed that this leads to a reduction in the residual stresses after shaping by pressing, and to better resistance to corrosion under stress of the thereby shaped parts. This element segregates at the austenitic grain boundaries and increases their cohesion. Boron precipitates for example in the form of borocarbides and boronitrides.

Nickel may be used optionally for increasing the strength of the steel by solution hardening. However, it is desirable, among others for cost reasons, to limit the nickel content to a maximum content of 1.0% or less and preferably between below 0.3%.

Likewise, optionally, an addition of copper with a content not exceeding 5% is one means of hardening the steel by precipitation of copper metal. However, above this content, copper is responsible for the appearance of surface defects in hot-rolled sheet. Preferably, the amount of copper is below 2.0%. Preferably, the amount of Cu is above 0.1 %.

Titanium and Niobium are also elements that may optionally be used to achieve hardening and strengthening by forming precipitates. However, when the Nb or Ti content is greater than 0.50%, there is a risk that an excessive precipitation may cause a reduction in toughness, which has to be avoided. Preferably, the amount of Ti is between 0.040 and 0.50% by weight or between 0.030% and 0.130% by weight. Preferably, the titanium content is between 0.060% and 0.40 and for example between 0.060% and 0.1 10% by weight. Preferably, the amount of Nb is above 0.01 % and more preferably between 0.070 and 0.50% by weight or 0.040 and 0.220%. Preferably, the niobium content is between 0.090% and 0.40% and advantageously between 0.090% and 0.200% by weight.

Chromium and Molybdenum may be used as optional element for increasing the strength of the steel by solution hardening. However, since chromium reduces the stacking fault energy, its content must not exceed 1.0% and preferably between 0.070% and 0.6%. Preferably, the chromium content is between 0.20 and 0.5%. Molybdenum may be added in an amount of 0.40% or less, preferably in an amount between 0.14 and 0.40%.

Furthermore, without willing to be bound by any theory, it seems that precipitates of vanadium, titanium, niobium, chromium and molybdenum can reduce the sensitivity to delayed cracking, and do so without degrading the ductility and toughness properties. Thus, at least one element may be chosen from titanium, niobium, chromium and molybdenum under the form of carbides, nitrides and carbonitrides.

Optionally, tin (Sn) is added in an amount between 0.06 and 0.2% by weight, without willing to be bound by any theory, it is believed that since tin is a noble element and does not form a thin oxide film at high temperatures by itself, Sn is precipitated on a surface of a matrix in an annealing prior to a hot dip galvanizing to suppress a pro-oxidant element such as Al, Si, Mn, or the like from being diffused into the surface and forming an oxide, thereby improving galvanizability. However, when the added amount of Sn is less than 0.06%, the effect is not distinct and an increase in the added amount of Sn suppresses the formation of selective oxide, whereas when the added amount of Sn exceeds 0.2%, the added Sn causes hot shortness to deteriorate the hot workability. Therefore, the upper limit of Sn is limited to 0.2% or less.

The steel can also comprise inevitable impurities resulting from the development. For example, inevitable impurities can include without any limitation: O, H, Pb, Co, As, Ge, Ga, Zn and W. For example, the content by weight of each impurity is inferior to 0.1 % by weight.

Preferably, the mean size of grain of steel is up to 5pm, preferably between 0.5 and 3pm.

In a preferred embodiment, the steel sheet is covered by a metallic coating.

The metallic coating can be an aluminum-based coating or a zinc-based coating.

Preferably, the aluminum-based coated comprises less than 15% Si, less than 5.0% Fe, optionally 0.1 to 8.0% Mg and optionally 0.1 to 30.0% Zn, the remainder being Al.

Advantageously, the zinc-based coating comprises 0.01-8.0% Al, optionally

0.2-8.0% Mg, the remainder being Zn.

For example, the coated steel is a galvannealed steel sheet obtained after an annealing step performed after the coating deposition.

In a preferred embodiment, the steel sheet has a thickness between 0.4 and 1 mm.

The method according to the present invention for producing a TWIP steel sheet comprises the following steps:

A. feeding of a slab having the above composition,

B. Reheating such slab and hot rolling it,

C. A coiling step,

D. A first cold-rolling,

E. An recrystallization annealing,

F. A second cold-rolling and G. A recovery heat treatment.

According to the present invention, the method comprises the feeding step A) of a semi product, such as slabs, thin slabs, or strip made of steel having the composition described above, such slab is cast. Preferably, the cast input stock is heated to a temperature above 1000°C, more preferably above 1050°C and advantageously between 1 100 and 1300°C or used directly at such a temperature after casting, without intermediate cooling.

The hot-rolling is then performed at a temperature preferably above 890°C, or more preferably above 1000°C to obtain for example a hot-rolled strip usually having a thickness of 2 to 5 mm, or even 1 to 5 mm. To avoid any cracking problem through lack of ductility, the end-of-rolling temperature is preferably above or equal to 850° C.

After the hot-rolling, the strip has to be coiled at a temperature such that no significant precipitation of carbides (essentially cementite (Fe,Mn) ₃ C) occurs, something which would result in a reduction in certain mechanical properties. The coiling step C) is realized at a temperature below or equal to 580°C, preferably below or equal to 400°C.

A subsequent cold-rolling operation followed by a recrystallization annealing is carried out. These additional steps result in a grain size smaller than that obtained on a hot-rolled strip and therefore results in higher strength properties. Of course, it must be carried out if it is desired to obtain products of smaller thickness, ranging for example from 0.2 mm to a few mm in thickness and preferably from 0.4 to 4mm. A hot-rolled product obtained by the process described above is cold- rolled after a possible prior pickling operation has been performed in the usual manner.

The first cold-rolling step D) is performed with a reduction rate between 30 and 70%, preferably between 40 and 60%.

After this rolling step, the grains are highly work-hardened and it is necessary to carry out a recrystallization annealing operation. This treatment has the effect of restoring the ductility and simultaneously reducing the strength. Preferably, this annealing is carried out continuously. Advantageously, the recrystallization annealing E) is realized between 700 and 900°C, preferably between 750 and 850°C, for example during 10 to 500 seconds, preferably between 60 and 180 seconds.

Then, a second cold-rolling step F) is realized with a reduction rate between 1 to 50%, preferably between 10 and 40% and more preferably between 20% and 40%. It allows for the reduction of the steel thickness. Moreover, the steel sheet manufactured according to the aforesaid method, may have increased strength through strain hardening by undergoing this re-rolling step. Additionally, this step induces a high density of twins improving thus the mechanical properties of the steel sheet.

After the second cold-rolling, a recovery step G) is realized in order to additionally secure high elongation and bendability of the re-rolled steel sheet. Recovery is characterized by the removal or rearrangement of dislocations in the steel microstructure while keeping the deformation twins. Both deformation twins and dislocations are introduced by plastic deformation of the material, such as rolling step.lt is believed that the recovery step allows for an increase of the mechanical properties such as the elongation.

Thus, in addition to the high amount of C in the TWIP steel according to the present invention, a recovery step is performed allowing an improvement of notably the elongation. And, thanks to the combination of the specific TWIP steel and the method comprising the recovery step according to the present invention, it is possible to obtain a cold-rolled and recovered TWIP steel having a high mechanical resistance and a high elongation.

In a preferred embodiment, a recovery step G) is performed by heating the steel sheet at a temperature between 390 and 700°C and preferably 410 and 700°C in a batch annealing or a continuous annealing furnace. In this embodiment, a hot-dip galvanizing step H) can then be performed.

In another preferred embodiment, the recovery step G) is performed by hot- dip galvanization. In this case, the recovery step G) and the hot-dip galvanization are realized in the same time allowing cost saving and the increase of the productivity.

Preferably, the temperature of the molten bath is between 410 and 700°C depending on the nature of the molten bath. Advantageously, the steel sheet is dipped into an aluminum-based bath or a zinc-based bath. Preferably, the dipping into a molten bath is performed during 1 to 60 seconds, more preferably between 1 and 20 seconds and advantageously, between 1 to 10 seconds.

In a preferred embodiment, the aluminum-based bath comprises less than

15% Si, less than 5.0% Fe, optionally 0.1 to 8.0% Mg and optionally 0.1 to 30.0% Zn, the remainder being Al. Preferably, the temperature of this bath is between 550 and 700°C, preferably between 600 and 680°C.

In another preferred embodiment, the zinc-based bath comprises 0.01-8.0% Al, optionally 0.2-8.0% Mg, the remainder being Zn. Preferably, the temperature of this bath is between 410 and 550°C, preferably between 410 and 460°C.

The molten bath can also comprise unavoidable impurities and residuals elements from feeding ingots or from the passage of the steel sheet in the molten bath. For example, the optionally impurities are chosen from Sr, Sb, Pb, Ti, Ca, Mn, Sn, La, Ce, Cr, Zr or Bi, the content by weight of each additional element being inferior to 0.3% by weight. The residual elements from feeding ingots or from the passage of the steel sheet in the molten bath can be iron with a content up to 5.0%, preferably 3.0%, by weight.

Advantageously, the recovery step G) is performed between 1 second and 1 hour and 10 minutes, preferably between 30 seconds and 1 hour and more preferably between 30 seconds and 30minut.es.

For example, an annealing step can be performed after the coating deposition in order to obtain a galvannealed steel sheet.

A TWIP steel sheet comprising an austenitic matrix having a high strength, an excellent formability and elongation is thus obtainable from the method according to the invention.

Example

In this example, TWIP steel sheets having the following weight composition were used: * 0.856 0.21 21.94 0.027 0.114 1.35 0.155 0.04 0.891 0.008

* 0.876 0.502 17.63 0.032 0.108 2.78 0.149 - 0.384 0.0061 - -

^* examples according to the present invention.

Firstly, the samples were heated and hot-rolled at a temperature of 200°C. The finishing temperature of hot-rolling was set to 890°C and the coiling was performed at 400°C after the hot-rolling. Then, a 1 ^st cold-rolling was realized with a cold-rolling reduction ratio of 50%. Thereafter, a recrystallization annealing was performed at 850°C during 180seconds. Afterwards, a 2 ^nd cold-rolling was realized with a cold-rolling reduction ratio of 30%.

Finally, a recovery heat step was performed during 1 hour at 400°C for Trials 1 and 2 in a batch annealing.

For Trials 3 to 5, a recovery heat treatment was performed during 60 seconds in total. The steel sheet was first prepared through heating in a furnace up to 625°C, the time spent between 460 and 625°C being 54seconds and then dipped into a zinc bath during respectively 6s. The molten bath temperature was of 460°C.The following Table shows the mechanical properties of all Trials, after the recrystallization annealing E), after the second-rolling step F) and after the recovery step G).

Results show that Trials 2, 4 and 5, having a composition according to the invention have higher mechanical properties than Trials 1 and 3 having a composition outside the range of the invention. Indeed, the specific composition of the TWIP steel in addition to the method according to the present invention allows for a high UTS and a high TE.