The invention relates to a method of producing an austenitic steel sheet excellent in resistance to delayed cracking.

In view of fuel economy and safety in the case of collisions, high strength steels are used more and more in the automobile industry. This requires the use of structural materials that combine a high tensile strength with high ductility. Austenitic alloys comprising as main elements iron, carbon and high levels of manganese, which can be hot-rolled or cold-rolled and have a strength that may exceed 1000 MPa. The mode of deformation of these steels depends on the stacking fault energy: for a sufficiently high stacking fault energy, the observed mode of mechanical deformation is twinning, which results in a high work hardenability. By acting as an obstacle to the propagation of dislocations, the twins increase the flow stress. However, when the stacking fault energy exceeds a certain limit, slip of perfect dislocations becomes the main deformation mechanism and the work hardening is reduced. It is known that the sensitivity to delayed cracking increases with the mechanical strength, in particular after certain cold-forming operations since high residual tensile stresses are liable to remain after deformation. In combination with atomic hydrogen possibly present in the metal, these stresses are liable to result in delayed cracking, that is to say cracking that occurs a certain time after the deformation itself. Hydrogen may progressively build up by diffusion to crystal lattice defects, such as matrix/inclusion interfaces, twin boundaries and grain boundaries. It is in the latter areas that hydrogen may become harmful when it reaches a critical concentration after a certain time. For a constant grain size, the time required to attain a critical level depends on the initial concentration of mobile hydrogen, the intensity of the residual stress concentration field and the kinetics of hydrogen diffusion.

In particular circumstances, small amounts of hydrogen may be introduced at some stages of steel fabrication such as chemical or electrochemical pickling, annealing under special atmospheres, electroplating or hot dip galvanizing. Subsequent machining operations using lubricating oils and greases may also cause hydrogen production after decomposition of these substances at high temperatures.

It is an object of this invention to provide a method of producing an austenitic steel sheet excellent in resistance to delayed cracking.

It is a further object of this invention to provide a method of producing an austenitic steel sheet which is energy efficient and simple in comparison to the conventional route for this type of steel.

According to the invention one or more of these objects is reached by providing a method of producing an austenitic steel sheet excellent in resistance to delayed cracking comprising

• casting an ingot, or a continuously cast slab, or a continuously cast thin slab or a strip-cast strip, the composition thereof comprising, in weight:
  • 0.50% - 0.80 %C
  • 10 - 17% Mn
  • 1.0 - 5% Al
  • at most 0.5% Si
  • at most 0.020% S
  • at most 0.050% P
  • 50 - 200 ppm N
  • 0.050 - 0.25 % V
  the remainder being iron and unavoidable impurities inherent to fabrication;
• providing a hot-rolled strip by hot rolling the ingot, the continuously cast slab, the continuously cast thin slab or the strip-cast strip to the desired hot rolled thickness.
• cold-rolling the hot-rolled strip to the desired final thickness,
• continuous annealing of the cold-rolled strip in a process comprising heating up the strip at a heating rate V_h of between 3 and 60°C/s to an annealing temperature T_a for an annealing time t_a of between 15 and 300 seconds followed by cooling at a cooling rate V_c of between 10 and 100°C/s and wherein T_a is from 750 to 850°C.

By using a high aluminium content the SFE of the steel increases. Any adverse effects of the elements that lower the SFE, such as silicon, are counteracted by the addition of aluminium. Additionally, aluminium lowers the activity and the diffusivity of carbon in austenite, which reduces the driving force for forming carbides. The vanadium, which is added as an essential alloying addition, forms carbides. These vanadium-carbides act as hydrogen sinks if and when the size and distribution of the vanadium-carbides is correct. The increased aluminium content is therefore essential to control the vanadium-carbide precipitation because it prevents the vanadium-carbides from coarsening due to the reduced carbon activity and diffusivity as a result of the presence of aluminium. The inventors found that at least 1.0 % Al and from 0.050 to 0.25 % V is required to achieve this. Lower aluminium contents led to too coarse vanadium-carbides, thereby rendering them ineffective as hydrogen sinks, and the amount of Vanadium needs to be controlled between the mentioned values to achieve a sufficient number of small precipitates. Higher V-values lead to early nucleation of the precipitates and thereby inevitably to coarse and fewer precipitates, whereas values below 0.050% V simply result in to few precipitates, even if they are fine enough. The annealing treatment is crucial in that it controls the precipitation of the vanadium-carbides and causes the recrystallisation of the cold-deformed microstructure caused by the cold-rolling to result in a fine grain structure. In a preferable embodiment the silicon content is very low, i.e. at impurity level. In principle the aluminium content is limited only by the fact that the steel according to the invention is an austenitic steel.

EP2090668-A1 discloses a method of producing high strength in a Cr-bearing TWIP steel by subjecting the steel to a cold reduction. US20080035249-A1 disclose TWIP steels with Mn contents of over 21%.

The maximum aluminium content is 5%. Preferably the aluminium content is at least 1.25 and/or at most 3.5%, more preferably at least 1.5 and/or at most 2.5%.

In an embodiment the maximum annealing temperature T_a is 825°C or even 800°C. Preferably the cooling rate V_c is between 20 and 80 °C/s. The heating rate is according to the invention is between 3 and 60°C/s. The annealing time t_a according to the invention is between 15 and 300 seconds.

In a preferable embodiment, the maximum annealing temperature T_a is from 775 to 795 °C (i.e. 785 ± 10°C).

Preferably, the steel strip material has been pickled before cold rolling. Pickling is (often) necessary before cold rolling to remove oxides, to prevent rolling in of oxides. Preferably, the cold rolled strip material is produced from a hot rolled strip material or a belt cast strip material.

In a preferred embodiment of the invention during the cooling at a cooling rate V_c after the continuous annealing the strip is led through a hot dip bath for providing a metallic coating by hot dipping the strip into a molten bath of the metal making up the metallic coating. This process leads to a very economical and quick process to produce a metallically coated steel strip. The metallic coating can be any known common coating such as zinc or zinc alloy, wherein the zinc may be alloyed with elements such as aluminium and/or magnesium.

In another embodiment of the invention the strip is pickled after continuous annealing and wherein the strip is provided with a metallic coating by pickling after annealing followed by heating to a temperature below the continuous annealing temperature before the strip is led through a hot dip bath for providing a metallic coating by hot dipping the strip into a molten bath of the metal making up the metallic coating. This alternative process is available if the economical process as described hereinabove is not preferred. There may be issues with adhesion with certain specific metallic coatings for which a pickling treatment may be necessary. After pickling, it is not necessary nor desirable to heat the strip above the T_a. It is preferable that the heating temperature remain below T_a.

With this method, the strip material is heated only to a temperature high enough to form a closed inhibition layer. This temperature is lower than the normal continuous annealing temperature necessary for metallurgical reasons (such as recrystallisation to influence mechanical properties). The forming of oxides on the surface of the steel strip material is thereby reduced.

Preferably, the temperature below the continuous annealing temperature is between 400 and 600 °C. In this temperature range the forming of oxides is considerably reduced and the strip material is heated sufficiently for the subsequent hot dip galvanizing.

According to a preferred embodiment, the Fe in the strip material is reduced during or after the heating to a temperature below the continuous annealing temperature and before the hot dip galvanising. By reducing the strip material, the Fe-oxides that are formed are reduced, and in this way the amount of oxides present on the surface of the strip material before hot dip galvanizing is decreased considerably.

Preferably, the reduction is performed using H₂N₂, more preferably using 5 - 30% H₂N₂ in the reducing atmosphere. It has been found that with the use of this atmosphere most oxides can be removed.

According to a preferred embodiment, an excess amount of O₂ is provided in the atmosphere during or after the heating of the strip material and before the reduction of the strip material. The providing of an excess amount of oxygen improves the quality of the surface of the steel strip material before the hot dip galvanizing, and thus the quality of the zinc layer coated on the AHSS strip material. It is supposed that the oxygen binds the alloying elements in the AHSS strip material both at the surface of the strip material and internally, and that in this way the oxides formed cannot migrate to the surface of the strip material.

The reducing atmosphere that follows after the oxidation will then reduce the oxides at the surface of the strip material, and in this way the amount of oxides at the surface of the strip material is considerably reduced or even almost absent, as experiments have shown. Preferably, the excess amount of O₂ is provided in an amount of 0,05 - 5 % O₂. This amount of oxygen has been found to suffice.

In a preferred embodiment of the invention V-alloyed TWIP steel strip material according to the invention was hot rolled, pickled and cold rolled, and continuously annealed to a temperature according to the invention and pickled again. Then the strip material is heated to a temperature of 527 °C in an annealing line, and thereafter hot dip galvanised in a galvanising bath at approximately 450°C.

During the heating of the strip material to the temperature of 527 °C, an excess amount of 1% 02 is provided. The oxygen is provided at such a high temperature so that not only forms oxides at the surface of the strip material, but also at some depth under the surface binds the alloying elements. After the providing of the oxygen, the strip material is reduced using approximately 5% H₂N₂. The reduction of the strip material removes the oxides from the surface, but the oxides formed under the surface remain where they are and cannot migrate to the surface.

Thus, by reducing the surface the oxides are effectively removed and no new oxides can be formed at the surface. These oxides, when not removed, cause a bad adhesion of the zinc layer to the substrate, resulting in bare spots, flaking and the forming of cracks in the zinc layer when the material is bent. It is presumed that by normal reduction, the alloying elements migrate to the surface very fast at the alloying temperature and thus form oxides at the surface again before the hot dip galvanising takes place. Whatever the exact mechanism may be, it has been found that the use of this method diminishes or almost eliminates the amount of oxides found in a hot dip galvanised zinc layer on a V-alloyed TWIP steel.

In an embodiment of the invention the cold-rolling reduction is between 10 to 90%, more preferably between 30 and 85, even more preferably between 45 and 80%.

In an embodiment of the invention the annealed strip is temper rolled with a reduction of from 0.5 to 10% prior to or after the metallic coating has been provided to the strip.

In an embodiment of the invention the Vanadium content is between 0.06 and 0.22%.

In a second aspect of the invention a strip or sheet is provided produced by a method according to any one of claims 1 to 6, wherein the steel preferably is provided with a metallic coating. In a preferred embodiment of the invention the strip or sheet is used for the production of automotive inner or outer parts or wheels or for hydroforming applications.

The invention will now be further explained by means of the following, non-limitative examples:
The chemical compositions of the materials used in this study are shown in Table 1. Code C Si Mn Al V S N P 8V-1 0.74 0.36 13.8 2.2 0.080 0.011 0.011 Low* 8V-2 0.74 0.36 13.8 2.3 0.086 0.011 0.009 Low 8V-3 0.74 0.31 13.7 2.3 0.084 0.011 0.007 Low 10V 0.71 0.22 13.4 2.4 0.106 0.010 0.011 Low 16V 0.72 0.25 14.5 2.3 0.160 0.001 0.006 0.008 21V 0.69 0.21 14.9 2.6 0.213 0.010 - 0.005 Si-free 0.70 <0.20 14.5 2.5 impurity 0.004 0.007 0.030 * Low = impurity level

The finish rolling temperature (FRT) was chosen to ensure recrystallisation of the deformed microstructure and the coiling temperature was kept below 500°C to avoid carbide precipitation. Recrystallisation does not solely depend on the FRT but also depends on the time, on the rolling strain accumulated since the last recrystallisation event during hot rolling and on the strain rate.

All hot rolled materials were 50% cold rolled and subsequently recrystallisation annealed. Different annealing cycles were applied to determine the optimal annealing parameters. Note that the elongations were between 45% and 50% for all samples, except for those that were not recrystallised (36-45%) and the material annealed 920°C (65%). Since the strength is considered more important, the following discussion will focus on that.

For annealing temperatures up to 750°C, the material softens due to an increased fraction of recrystallised material and probably some grain growth. At these temperatures, the effect of precipitation is limited. The difference between the (fully recrystallised) materials annealed at 775°C and 800°C is small because the precipitation is considered optimal in this temperature region for minimising grain growth. Based on these observations, the recommended annealing temperature is 785°C ± 10°C. Grade Si-free Heating rate (°C/s) T_a (°C) t_a (s) Rp (MPa) Rm (MPa) Ag (%) Atot (%) Si-free 5 700 270 580 920 43 Si-free 5 750 270 546 900 46 Si-free 5 800 270 490 875 47 Si-free 5 850 270 450 850 48 Si-free 5 920 270 410 870 65 8V-1 5 785 270 584 962 44 46 8V-2 5 785 270 595 970 45 49 8V-3 5 785 270 585 964 46 50 10V 5 785 270 575 939 45 49 16V 5 700 270 764 1072 32 36 16V 5 725 270 687 1012 37 41 16V 5 750 270 636 977 41 45 16V 5 775 270 614 964 42 46 16V 5 800 270 590 946 43 48 16V 5 750 60 670 1004 39 43 16V 5 775 60 626 973 41 45 16V 5 800 60 597 955 44 48 21V 30 700 60 655 985 45 21V 30 750 60 600 950 43 21V 30 920 60 425 835 53

The delayed cracking and stress corrosion cracking results on the V-alloyed grades show a lower susceptibility to crack formation as the material is annealed at a higher temperature. For the stress corrosion cracking sensitivity, V addition is clearly beneficial at an annealing temperature of 750°C, but also at higher annealing temperatures.

The V-alloys were subjected to a resistance spot welding tests. Hot cracking in the weld was largely reduced compared to the Si-free non V-alloyed material.