EP1427866B1 - Method for making rolled and welded tubes comprising a final drawing or hydroforming step and resulting rolled tube - Google Patents

Abstract

The invention concerns a method for making a welded tube, comprising a final drawing or hydroforming step, characterized in that in consists in: preparing an alloy consisting in wt. % of: C </= 2 %; Mn between 10 and 40 % with %Mn > 21.66 - 9.7 %C; Si </= 5 %; S </= 0.3 %; P </= 0.1 %; Al </= 5 %; Ni </= 5 %; Mo </= 5 %; Co </= 3 %; W </= 2 %; Cr </= 5 %; Nb </= 1 %; V </= 1 %; Cu </= 5 %; N </= 0.2 %; Sn </= 0.5 %; Ti </= 1 %; B </= 0.1 %; each of the Ca and Mg contents </= 0.1 %; each of the As and Sb contents </= 0.1 %; casting it as semi-product, a) either in the form of an ingot which is then roughed down to be transformed into a slab, or directly in slab form, said slab being then hot-rolled in the form of a strip and then coiled; b) either in the form of a thin strip; in scouring the strip if the latter is oxydized at the surface; making a welded tube by forming a metal sheet cut from said strip to bring together its edges until they abut, then by welding said edges, then by eliminating the weld bead, then by cold drawing or hydroforming.

Description

The invention relates to iron and steel industry. More specifically, it concerns the manufacture of welded tubes, generally of small dimensions, this manufacture ending with a definitive shaping step by stretching or hydroforming.

A wide variety of steel grades can be used for make welded tubes of small dimensions, that is to say a few centimeters in diameter, typically 2 to 10 cm, and a few millimeters thickness, typically of the order of 5 mm. For applications not requiring no particular properties for the final product, such as scaffold tubes, carbon and carbon steels are usually used. low-end manganese. For more demanding applications for example, for the automotive market, more complex steels are used. The exhaust lines, for example, are made of steel stainless steel, ferritic or austenitic, whose properties are adjusted in playing on the conditions of annealing, hardening and drawing, or steel aluminized carbon. The structural parts of automobiles, trucks and of railway material are conventionally made of carbon-manganese steels high strength ferrito-pearlitic structure with up to 0.2% carbon and 1.5 to 2% of manganese, these steels being stretched then a normalization annealing. It is also possible to use rolled steels hot high strength ferritic-bainitic structure or rolled steels dual-phase hot-rolled ferritic-martensitic structures, or cold dual-phase. All these steels can reach high prices, less because of the cost of their raw material than the cost of multiple operations annealing and shaping they must undergo.

The object of the invention is to provide manufacturers and users of small welded tubes, particularly in the automobile industry, a process for economic manufacture resulting in the production of products with high mechanical characteristics.

• on procède à l'élaboration d'un alliage de composition, exprimée en pourcentages pondéraux :
  • C ≤ 2%;
  • Mn compris entre 10 et 40%, avec Mn% > 21,66 - 9,7 C% ;
  • Si ≤ 5%, préférentiellement ≤ 1 %, optimalement ≤ 0,5% ;
  • S ≤ 0,3%, préférentiellement ≤ 0,05%, optimalement ≤ 0,01 % ;
  • P ≤ 0,1%, préférentiellement ≤ 0,05% ;
  • Al ≤ 5%, préférentiellement ≤ 0,1 %, optimalement ≤ 0,03% ;
  • Ni ≤ 5%, préférentiellement ≤ 2% ;
  • Mo ≤ 5%, préférentiellement ≤ 1 % ;
  • Co ≤ 3%, préférentiellement ≤ 1 % ;
  • W ≤ 2%, préférentiellement ≤ 0,5% ;
  • Cr ≤ 5%, préférentiellement ≤ 1 % ;
  • Nb ≤ 1%, préférentiellement ≤ 0,1 % ;
  • V ≤ 1 %, préférentiellement ≤ 0,1 % ;
  • Cu ≤ 5%, préférentiellement ≤ 1 % ;
  • N ≤ 0,2%, préférentiellement ≤ 0,1%, optimalement ≤ 0,05% ;
  • Sn ≤ 0,5%, préférentiellement ≤ 0,1 % ;
  • Ti ≤1 %, préférentiellement ≤ 0,1 % ;
  • B ≤ 0,1 %, préférentiellement ≤ 0,01 % ;
  • chacune des teneurs en Ca et Mg ≤ 0,1 °/a, préférentiellement ≤ 0,01 % ;
  • chacune des teneurs en As et Sb ≤ 0,1 %, préférentiellement ≤ 0,05% ;
• on procède ensuite à la coulée d'un demi-produit à partir de cet alliage,
• a) soit sous forme d'un lingot qui subit ensuite un dégrossissage par laminage à chaud pour le transformer en brame, soit directement sous forme d'une brame, ladite brame étant ensuite laminée à chaud sous forme d'une bande puis bobinée,
• b) soit sous forme d'une bande mince ;
• on procède ensuite à un décapage de la bande si celle-ci est oxydée en surface ;
• on procède enfin à la fabrication du tube soudé par formage progressif d'une tôle découpée à partir de la bande précédente pour amener ses bords jusqu'à accostage, puis par soudage desdits bords, puis par élimination du bourrelet de soudure, puis par étirage à froid ou hydroformage.

For this purpose, the subject of the invention is a method for manufacturing a welded tube, of the type comprising a final stretching or hydroforming step,
characterized in that

• an alloy of composition is produced, expressed in weight percentages:
  • C ≤ 2%;
  • Mn between 10 and 40%, with Mn%> 21.66 - 9.7 C%;
  • If ≤ 5%, preferably ≤ 1%, optimally ≤ 0.5%;
  • S ≤ 0.3%, preferably ≤ 0.05%, optimally ≤ 0.01%;
  • P ≤ 0.1%, preferentially ≤ 0.05%;
  • Al ≤ 5%, preferably ≤ 0.1%, optimally ≤ 0.03%;
  • Ni ≤ 5%, preferably ≤ 2%;
  • Mo ≤ 5%, preferably ≤ 1%;
  • Co ≤ 3%, preferably ≤ 1%;
  • W ≤ 2%, preferably ≤ 0.5%;
  • Cr ≤ 5%, preferably ≤ 1%;
  • Nb ≤ 1%, preferentially ≤ 0.1%;
  • V ≤ 1%, preferentially ≤ 0.1%;
  • Cu ≤ 5%, preferably ≤ 1%;
  • N ≤ 0.2%, preferentially ≤ 0.1%, optimally ≤ 0.05%;
  • Sn ≤ 0.5%, preferentially ≤ 0.1%;
  • Ti ≤1%, preferentially ≤ 0.1%;
  • B ≤ 0.1%, preferentially ≤ 0.01%;
  • each of Ca and Mg contents ≤ 0.1 ° / a, preferably ≤ 0.01%;
  • each of the contents of As and Sb ≤ 0.1%, preferably ≤ 0.05%;
• a semi-finished product is then cast from this alloy,
• a) in the form of an ingot which is then subjected to a hot-rolling roughing to transform it into a slab, or directly in the form of a slab, said slab then being hot-rolled in the form of a strip and then wound,
• (b) in the form of a thin strip;
• the strip is then stripped if the strip is oxidized on the surface;
• the welded tube is finally made by progressive forming of a sheet cut from the preceding strip to bring its edges until docking, then by welding said edges, then by removing the weld bead, then by stretching to cold or hydroforming.

Preferably, the carbon content of the alloy is between 0 and and 1.2% and the manganese content of the alloy is between 10 and 35%.

Even more preferentially, the carbon content of the alloy is between 0.2 and 1.2%, and the manganese content of the alloy is between 10 and 30%.

Very advantageously, the carbon content of the alloy is included between 0.2 and 0.8%, and the manganese content of the alloy is between 15 and 30%.

Optimally, the carbon content of the alloy is between 0.4 and 0.8%, and the manganese content of the alloy is between 20 and 24%.

Hot rolling may be preceded by reheating a temperature not exceeding 80 ° C below the temperature of solidus of the alloy.

Hot rolling may be preceded by reheating a temperature at which the precipitation of nitrides is not caused aluminum.

The end temperature of hot rolling is preferably greater than or equal to 900 ° C.

The winding temperature after hot rolling is preferably less than or equal to 450 ° C.

Annealing followed by over-tempering of the strip wound hot-rolled, said annealing being effected under conditions allowing the resetting of carbides and avoiding their precipitation at cooling.

After hot rolling and any annealing followed by tempering, cold rolling of the strip can be carried out with a minimum reduction rate of 25%, preceded by stripping.

The rate of reduction of the thickness of the strip during the first cold rolling of the web is preferably at least 25%.

We can proceed to a recrystallization annealing of the band at a temperature of 600 to 1200 ° C for 1 second to 1 hour.

The subject of the invention is also a welded tube produced by the previous process.

As will be understood, the invention consists first of all in using a iron-carbon-manganese alloy of determined composition, and to subject it to a series of thermomechanical treatments, before its form of tubes, which provide the desired mechanical properties.

These alloys have, in fact, a strong work hardening capacity which allows them to associate, at the end of these treatments, a very high resistance (up to 1200 MPa) at high ductility (resulting in a high elongation at break up to 90%). They present the desired characteristics for the production of small tubes such as those used by the automobile industry to build, thanks to their resistance high, reinforcing parts of the vehicle structure, such as bars anti-intrusion integrated doors. Their ductility reserve makes them also suitable for use in forming side members, which must be capable of absorbing high deformation energy.

The invention will be better understood on reading the description which follows, given with reference to the attached single figure showing the energy of lack of stacking of an iron-carbon-manganese alloy according to its composition, at a temperature of 300 K.

First, we proceed to the development of a ferrous alloy austenitic iron-carbon-manganese, whose carbon and manganese are in the following ranges (all grades are data in percentages by weight).

Carbon and manganese contents are included respectively between 0 and 2% and 10 and 40%, preferably respectively between 0 and 1.2% and 10 and 35%., very preferably respectively between 0.2 and 1.2% and 10 and 30%, very advantageously respectively between 0.2 and 0.8% and 15 and 30%, and optimally respectively between 0.4 and 0.8% and 20 and 24%;

The permissible contents for the other elements of the alloy are the following, knowing that all these elements may be present only in the state trace amounts (for this reason, we will not give minimum precise in these elements).

The silicon content must be less than or equal to 5%, preferably less than or equal to 1%, optimally less than or equal to 0.5%.

The sulfur content must be less than or equal to 0.3%, preference less than or equal to 0,05%, optimally less than or equal to 0.01%.

The phosphorus content must be less than or equal to 0,1%, preferably less than or equal to 0.05%.

The aluminum content must be less than or equal to 5%, preferably less than or equal to 0.1%, optimally less than or equal to 0.03%.

The nitrogen content is less than or equal to 0.2%, preferably less than or equal to 0.1%, optimally less than or equal to 0.05%.

The nickel content must be less than or equal to 5%, preferably less than or equal to 2%.

The molybdenum content must be less than or equal to 5%, preferably less than or equal to 1%.

The cobalt content must be less than or equal to 3%, preferably less than or equal to 1%.

The tungsten content must be less than or equal to 2%, preferably less than or equal to 0.5%.

The levels of niobium and vanadium must each be lower or equal to 1%, preferably less than or equal to 0.1%.

The chromium and copper contents must each be lower or equal to 5%, preferably less than or equal to 1%.

The tin content must be less than or equal to 0.5%, preferably less than or equal to 0,1%.

The titanium content must be less than or equal to 1%, preferably less than or equal to 0,1%.

These alloys can also tolerate maximum boron content 0.1%, preferably not more than 0.01%, a maximum calcium content or magnesium of 0.1%, preferably not more than 0.01%, arsenic or antimony maximum of 0,1%, not more than 0,05%.

The upper bounds that have been laid correspond to which, for certain elements, are beginning to be harmful to properties of the alloy. This is, for example, the case for aluminum and sulfur. For other elements, it is essentially economic criteria that have such upper bounds. Thus, there would be little metallurgical drawbacks to add more than 5% nickel to the alloy, but this would unnecessarily increase its cost price.

For the application targeted by the invention, we seek shades of Fe-C-Mn ternary system that provides high resistance (preferably at least 1000 MPa) and an equally high elongation (preferably at minus 50%). Moreover, for cost issues, it is not desirable to have a too high content of manganese.

It is known that the mode of deformation of austenitic steels and alloys depends on their chemical composition and the deformation temperature.

Compared to ferritic steels that deform predominantly by dislocation slip, austenitic ferric steels and alloys have many other modes of deformation, in addition to slip. Among them, if their stacking fault energy lends itself to it, there is twinning. This mode of deformation has the advantage of providing greater plastic deformation and, consequently, higher strength than those resulting from the simple sliding of dislocations. It is therefore necessary to seek conditions that are capable of activating twinning at the commissioning temperatures of the materials to be manufactured, in particular at ambient temperature for the case of motor vehicle parts, so as to obtain a great deal of work hardening capacity. The possibility of obtaining a mechanical twinning is governed firstly by the chemical composition of the alloy; secondly by the temperature at which the material is located, these two parameters acting on the stacking fault energy, and finally by the grain size of the material which determines the kinetics of the twinning. The excessive formation of martensite ε (more than 20% of the structure) and the formation of martensite α 'at the moment of deformation are also obstacles to obtaining satisfactory mechanical properties, in particular good ductility. It is therefore important to have, before the final step of shaping the tube, by cold drawing or by hydroforming, a material having all the desirable characteristics of these points of view. The process according to the invention gives access to such materials.
The single figure shows the theoretical evolution of stacking fault energy in the C / Mn plane at room temperature (300 K), in the form of curves along which the stacking fault energy, expressed in mJ / m ² , is constant. We have also shown in the figure a series of points of the plane C / Mn (marked by the sign ▪) for which a twinning was actually found either by various authors who published their results, or by the inventors, as well as a portion , which must be avoided, in the field of martensitic transformation γ → α 'induced by deformation. Table 1 includes the chemical characteristics and (for the samples tested by the inventors, that is to say the samples E to K) mechanical samples reported in the single figure. The mechanical properties mentioned are the tensile strength Rm, elongation at break A and their product. Samples E to K were annealed at 800 ° C. for 90 s, which gave them a grain size of 2 to 5 μm. This table also includes stacking defect energies (EDE) calculated at 300 K of the samples.

It is found that the twinning is observed at ambient temperature when the stacking fault energy varies approximately between 15 and 30 mJ / m ² , this area corresponding to a carbon content of 0 to 1.6% and to a content of in manganese from 10 to 35%. Below 10% of manganese, transformation into α-martensite is spontaneous. Beyond 10% of manganese, this transformation does not occur if, moreover, the carbon and manganese contents are linked by the relation: Mn%> 21.66 - 9.7 C%

When the temperature varies, the stacking fault energy also varies in the same direction, at a rate of ± 5 mJ / m ² for a temperature variation of ± 50 ° C. This feature is important if the shaping operation is to be performed at a temperature below ambient.

The influence of the grain size on the stacking fault energy can also be appreciated. A modification of the steel manufacturing process, for example a change in the annealing conditions after winding or after cold rolling, can lead to a significant variation in the grain size. Thus, a steel having a grain size of 50 μm has a stacking failure energy of 5 mJ / m ² less than that of a steel of similar composition having a grain size of 2 to 5 μm.

If we want to maintain the assurance of obtaining a mode of deformation by twinning, it is possible to play on the chemical composition of the steel to compensate for the effects of the variations of temperature of formatting and grain size. When a grain size of 50 μm is combined with a 50 ° C reduction in the shaping temperature, the stacking fault energy decreases by 10 mJ / m ² . According to the single figure, compensation for this reduction is achieved by increasing the carbon content by 0.4%, or by increasing the manganese content by 5%. In practice, an alloy according to the invention is thus obtained if its carbon content is between 0 and 2%, if its manganese content is between 10 and 40%, and if, in addition, these contents obey the relation (1), so as to avoid the formation of martensite α 'during deformation at ambient temperature.

In addition, it has been noted that the optimal resistance / ductility compromise of an iron-carbon-manganese alloy is obtained when the mode of activated deformation is twinning at the limit of appearance of the martensite ε. However, the comparison of stack fault energy calculations with microstructural observations reveals that the transition between martensitic transformation ε and twinning takes place when the stacking fault energy is of the order of 15 mJ. / m ² . It can be seen from the samples in Table 1 that samples G, H and I, which have a stacking failure energy of about 15 mJ / m ² , have the highest Rm.A products (greater than 60,000). , therefore the best resistance / ductility trade-offs. It should be noted, moreover, that the sample F has a product Rm.A of less than 60,000, although its stacking fault energy is also of the order of 15 mJ / m ² . But its carbon content is not sufficient to make it fully benefit from the phenomenon of dynamic hardening, which will be discussed later.

Taking into account the combined effects on the stacking fault energy that can be obtained by adjusting the temperature and the grain size, it is therefore necessary to consider, from the single figure, that the composition domain preferred, from this point of view, is delimited by the curves corresponding to stacking fault energies of 5 to 25 mJ / m ² , ie (for the temperature considered in the figure of 300 K) by a carbon content of 0 to 1.2% and a manganese content of 10 to 35%, the previous relationship (1) must also be satisfied.

• l'interaction entre le carbone et les dislocations ; le carbone favorise l'écrouissage en stimulant l'émission de nouvelles dislocations mobiles servant à relayer celles qu'il a immobilisées ; ce phénomène est appelé « vieillissement dynamique » ;
• le « pseudo-maciage » ; le maclage d'une solution solide interstitielle cubique à faces centrées altère le motif cristallin ; en effet, le cisaillement de maclage convertit les sites octaédriques en sites tétraédriques ; le carbone qui occupait, avant maclage, les sites octaédriques, plus spacieux que les sites tétraédriques, se retrouve dans les sites tétraédriques ; une distorsion du réseau se produit, analogue à celle qui accompagne la transformation martensitique.

On the other hand, the single figure shows that carbon and manganese both contribute to an increase in stacking fault energy. However, an increase of this magnitude is less expensive to obtain by carbon addition than by addition of manganese, given the cost of materials to achieve these additions. In addition, the substitution of carbon with manganese, with constant stacking failure energy in the field of deformation by twinning, at the limit of the appearance of martensite s, results in an increase of the mechanical characteristics, and this for two reasons:

• the interaction between carbon and dislocations; carbon promotes hardening by stimulating the emission of new mobile dislocations used to relay those immobilized; this phenomenon is called "dynamic aging";
• "pseudo-maciage"; the twinning of a cubic interstitial solid solution with centered faces alters the crystalline pattern; in fact, twinning shear converts octahedral sites into tetrahedral sites; the carbon which occupied, before twinning, the octahedral sites, more spacious than the tetrahedral sites, is found in the tetrahedral sites; a distortion of the network occurs, similar to that which accompanies the martensitic transformation.

The softening effect of pseudo-twinning and aging dynamic increases with the carbon content. The mechanical strength of steel is therefore increased, provided that the carbon content is at least 0.2%. Finally, we see in the single figure that the spacing of constant stacking fault energy lines increases with the content in carbon. This means that alloys with a high carbon content are less sensitive to differences in carbon content than alloys low carbon.

Partial replacement of manganese with carbon present therefore both economic and metallurgical benefits. According to single figure, an addition of 0.2% carbon makes it possible to dispense with 4 to 5% of manganese with constant stacking fault energy. An area still most preferred carbon and manganese content is therefore 0.2% ≤ C ≤ 1.2% and 10% ≤ Mn ≤ 30%, the relation (1) being otherwise satisfied.

Increasing the carbon content may, however, have disadvantages beyond a certain limit. Indeed, there is a risk that in alloys whose composition is in the preceding preferred range, a precipitation of carbides of the M ₅ C ₂ and M ₂₃ C _{6 type} occurs during a slow cooling. Such slow cooling may be that experienced by a coiled strip after being cast directly as a thin strip or hot rolled. M ₃ C carbide can also precipitate during pearlitic transformation.

These carbide precipitations are preferentially to be avoided because they deplete the carbon matrix and thus decrease the energy of lack of stacking, so tend to disadvantage twinning in favor of martensitic transformations γ → ε and / or γ → α '. As we said, the formation of martensite α 'should be avoided, and it is preferable that the proportion of martensite ε does not exceed 20% to avoid embrittlement of the material. On the other hand, these carbides, some of which are acicular, are themselves fragile and may cause cracks when unwinding of the wound tape. If the tape is to be wound at a relatively high temperature, so it is best not to impose a carbon content too high for steel, if you want to avoid having to proceed then to a solution annealing of the carbides followed by a annealing. In the majority of cases corresponding to the use of tools conventional industry, it will be preferable not to exceed a certain in carbon of 0.8%. In these circumstances, to compensate for the decrease in maximum carbon content compared to the previously preferred domain defined, it is necessary to raise the minimum manganese content up to 15%. We thus obtains an even more advantageous composition range where 0.2% ≤ C ≤ 0.8% and 15% ≤ Mn ≤ 30%.

It is the combination of deformation twinning with carbon hardening which makes it possible to combine strength and ductility, and thus to obtain high mechanical characteristics. Sample E has a stacking fault energy of 17 mJ / m ² , but contains only 0.19% carbon. Its resistance is only 750 Mpa. A carbon content of at least 0.4% is necessary to obtain a resistance greater than 950 MPa, as shown in sample F. This increase in the minimum carbon content requires the reduction of the maximum 24% manganese if it is desired to remain at a stacking fault energy value of about 15 mJ / m ² , and thus maintain the same degree of twinning by deformation.

The optimal composition range for the alloys of the invention is therefore 0.4% ≤ C ≤ 0.8% and 20% ≤ Mn ≤ 24%. We can, for example, propose a carbon content of 0.6% and a manganese content of 22%, as in samples G, H and I which show the most Rm.A products Table 1.

These values of the carbon and manganese contents are optimal in that they provide at room temperature adequate stack fault energies of the order of 5 to 25 mJ / m ² . However, if the shaping of the tubes is to be carried out at a temperature substantially lower than ambient, higher carbon and manganese maximum levels may be advisable for the stacking fault energy (which, as said it, decreases when the temperature drops) is kept at a level allowing a twinning is significantly observed. Therefore, in the spirit of the invention, the carbon content of the alloy can be up to 2% and the manganese content up to 40%.

Regarding other alloying elements used in the composition of the steel according to the invention or likely to enter, the comments can be formulated.

The maximum silicon content of 5% is justified by the need to maintain good weldability to the alloy. In practice, one less than 1%, of the order of 0.5% or less, is advisable. For the high levels of silicon, weldability problems can be reduced if welding is carried out in an inert atmosphere.

Requirements on maximum levels of sulfur, phosphorus, aluminum and nitrogen are due to the desire to obtain good forgeability hot for the material. Sulfur and phosphorus weaken the joints of grains, and too high levels of aluminum and nitrogen are likely to lead to the precipitation of aluminum nitrides that will hinder the migration of grain boundaries during hot processing. Maintain these elements in the specified ranges of contents makes it possible to maintain a good ductility of the material at hot rolling temperatures low enough not to cause surface defects of the type scale encrustations.

In addition, the inclusiveness of the alloy has an influence on its resistance and its elongation at break. The manganese sulphides constitute the main source of damage leading to a rupture premature. Improvement of the characteristics at break is therefore a reason additional to limit the sulfur content.

The need to limit the titanium, niobium and vanadium contents is due to the fact that these elements are likely to form carbonitrides which tend to slow down recrystallization by hindering the migration of joints. This is also the case for aluminum. As already mentioned, the grain size is an important parameter for setting properties mechanical properties of the material, and can be controlled by means of annealing recrystallization. For this recrystallization annealing to be fully effective, it is necessary to limit the formation of these carbonitrides.

The contents of chromium, nickel, molybdenum, copper, cobalt, tungsten, tin, boron, calcium, magnesium, arsenic and antimony must be maintained within the prescribed limits so that these elements do not have significant influence on the mechanical properties of the material.

The casting of the steel whose composition has been mentioned above can be in ingots or, preferably, continuously to obtain slabs of classical format, with a thickness of about 200 mm. he is also possible to cast this alloy in the form of thin slabs (a few cm thick) which can then undergo hot rolling in line. This process gives access to hot rolled strips of small thickness, which may possibly not subsequently undergo cold rolling. In this In this case, coarse-grained alloys are obtained (on the order of 20 μm, this value depending on the end of rolling and winding temperatures), presenting a relatively average resistance but a high ductility. It is also possible to achieve the casting of steel by a method of direct casting of thin strips, possibly being laminated hot online or offline. The application of this casting process to the casting of iron-carbon-manganese alloys (different from those of the invention) already been proposed in EP-A-1 067 203.

This casting step is widely known and does not show peculiarities compared to usual practices, it will not be more detailed here.

The hot product is then hot-rolled. casting. In the case of ingot casting, hot rolling begins by a heating followed by a roughing which brings the ingot to the format of a classic slab. In the case of a conventional continuous casting, one proceeds directly to hot rolling, after a heating step of the slab. These reheatings should not bring the slab to a temperature above the solidus temperature of the segregated zones, under pain of causing the appearance of "burns" that prohibit any in hot form. For example, the solidus temperature of a Fe-C-Mn alloy at 0.6% of carbon and 22% of manganese is of the order of 1280 ° C. Given the segregation of elements such as manganese, the phosphorus and carbon all of which have a partition coefficient less than 1, when treating this alloy, it is recommended not to exceed a reheat temperature of 1200 ° C for the front ingot roughing as for the slab before hot rolling. On the one In general, we can ensure the safety of this reheating by the performer at a temperature not exceeding the solidus temperature of the alloy minus 80 ° C.

Precipitation of aluminum nitrides during reheating is also preferably avoided. This precipitation hampers the migration of the joints during hot processing. The solubility product K _s of aluminum nitride as a function of the temperature T is expressed by: log K s = log ([Al] x [N]) = - 6770 T 1.72 where [Al] and [N] are the weight concentrations of aluminum and nitrogen in solid solution and T the temperature in Kelvin. Knowing these contents in the treated alloy, one can deduce the reheating temperature before hot rolling not to exceed to avoid the precipitation of aluminum nitrides. If the nitrogen content is 0.05% (preferred maximum), the preferred maximum solid solution aluminum content is 0.03% for reheating at 1200 ° C. For reheating temperatures lower than 1200 ° C, there is a risk of ending up with nitrogen and aluminum contents that may lead to a precipitation of nitrides, but in practice, these constraints are not very difficult to meet. However, they must be observed to obtain the most satisfactory results on the products obtained by the process according to the invention.

After reheating, it is advisable to carry out a descaling, if the atmosphere of the reheating furnace has been sufficiently oxidizing for cause a significant occurrence of calamine, without this descaling lead to excessive temperature loss before rolling at hot that follows. The presence of primary scale on the surface of the slab before its hot rolling can lead to encrustations of scale in the slab degrading the surface quality of the product. The calamine inlaid also causes deterioration of the mill rolls.

• de ne pas dépasser une teneur en soufre de 0,01 % dans l'alliage coulé ;
• et de ne pas laminer en dessous de 900°C, à moins d'avoir une teneur en soufre particulièrement faible (de l'ordre de 0,002% ou moins) et des teneurs en aluminium et azote qui garantissent absolument l'absence de précipités de nitrure d'aluminium après le réchauffage.

The slab is then hot-rolled to obtain a strip of thickness of the order of, for example, 2.5 to 3 mm. Since the alloys of the composition under consideration do not show any allotropic transformation in the temperature range under consideration, the rolling path in terms of the number of passes, the reduction ratio per pass and the time interval separating the passes, is immaterial. The only constraint is, most often, to respect a rolling end temperature of at least 900 ° C. Indeed, if the iron-carbon-manganese alloys prepared in the laboratory have a sufficient ductility to be able to be rolled up to 800 ° C without risking the appearance of cracks in banks, the forgeability of an iron-carbon-manganese alloy prepared under industrial conditions is also influenced by its contents of aluminum, nitrogen and sulfur. In the typical practice of a steel plant, we therefore advise:

• not to exceed a sulfur content of 0.01% in the cast alloy;
• and not to roll below 900 ° C, unless it has a particularly low sulfur content (of the order of 0.002% or less) and aluminum and nitrogen contents which absolutely guarantee the absence of precipitates of aluminum nitride after reheating.

Regarding the rolling conditions of iron-carbon-manganese alloys concerned by the invention, they may be comparable in terms of the rate of reduction per pass and the time interval between passes to those usually used for stainless steels austenitic type SUS 304, given the similarities of hardness to between SUS 304 and iron-carbon-manganese alloys the invention. As an indication, for a slab 160 mm thick, it is possible to set an exit temperature of the reheating furnace of 1100 ° C, a roughing cage outlet temperature of 980 ° C, a thickness in 38.5 mm roughing cage outlet, a temperature at the entrance of the finishing cage of 912 ° C, a rolling end temperature of 910 ° C, a band thickness at the end of rolling of 3 mm, an exit speed of 259 m / s band and a winding temperature of 480 ° C.

• l'appauvrissement en carbone de l'austénite provoqué par cette précipitation modifie l'énergie de défaut d'empilement qui a été réglée au moyen de la composition chimique pour que la cinétique de maclage soit optimale à la température ambiante, lors de la mise en forme du tube ; en cas de précipitation significative de carbures de fer, les propriétés mécaniques escomptées grâce à cette composition ne seront donc pas obtenues ;
• la précipitation des carbures de fer rend l'alliage cassant, donc difficilement laminable à froid.

The winding of the strip obtained after the hot rolling is then conventionally carried out. It is necessary to avoid the precipitation of iron carbides during cooling of the coils, because:

• the carbon depletion of the austenite caused by this precipitation modifies the stacking fault energy which has been adjusted by means of the chemical composition so that the kinetics of twinning is optimal at room temperature, when setting tube shape; in the event of significant precipitation of iron carbides, the mechanical properties expected from this composition will therefore not be obtained;
• the precipitation of the iron carbides makes the alloy brittle, so hard to cold roll.

Knowing that the cooling itself, of the order of 10 ° C / h, starts only one to two hours after winding, you have to wind the band at such a temperature that it can not be so prolonged to temperatures at which this precipitation of carbides from iron is possible. The winding temperature can be deduced from TTT diagrams of the alloy concerned. By way of example, for an iron-carbon-manganese alloy at 0,6% of carbon and 22% of manganese, a stay of 2 hours at a temperature of 500 ° C or more and 28 hours at 450 ° C or more leads to precipitation of iron carbides. Consequently, in industrial conditions where the coil must have completely cooled before any subsequent operation (stripping, cold rolling ...), it is it is preferable not to wind the strip at a temperature of more than 450 ° C. AT this effect, we can proceed to a cooling of the band after its hot rolling, so as to bring it to the winding temperature desired. We will take care to bring the band as late as possible to the temperature winding, for example by delaying the cooling, so as to allow a complete recrystallization of the steel before winding. So early brought to 450 ° C, the steel will no longer be able to recrystallize. Forced cooling immersion of the coil in the pool is also possible, always in the aim of avoiding the carbide precipitation domain.

However, where it would not have been possible to avoid such precipitation of carbides to the winding, it can then proceed to an annealing to dissolve these precipitates and thus return the carbon in solid solution, then to a hypertrempe to avoid the reprecipitation of the carbides with cooling after annealing. Typically, the band is worn up a temperature between 1000 and 1050 ° C at a speed such that the band remains one minute above 900 ° C, and 10 to 20 s above 1000 ° C, then it is cooled at a speed of at least 5 ° C / s. In general, the quenching is carried out to the maximum of the possibilities of the line.

In the case where the casting of the steel is carried out by casting of thin slabs or thin strips with possible hot rolling in line (not requiring, therefore, not necessarily reheat), those skilled in the art will be able to adapt the process previously described in consequence, knowing that the metallurgical imperatives discussed in the winding temperature and the possible need for annealing and hyper-tempering must also be taken into account here. account.

In cases where, as has been said, one wishes to obtain a product relatively thick and not very resistant, but with high ductility, the strip can be left as it is without proceed with its cold rolling. We can even try to make the grain by annealing followed by quenching after hot rolling, even if the winding conditions had prevented the precipitation of carbides. At this point, we have a grain size that is, in general, not less than 15-20 μm. On the other hand, if you want to get a band fine to make light tubes, and / or a band having a strong resistance can only be obtained with a particle size of 5 μm or less, cold rolling is necessary. This cold rolling allows also to reduce the roughness of the surface of the strip, so to obtain a surface appearance compatible with use for forming parts intended to remain visible. It also increases the band's ability to be coated.

Prior to its eventual cold rolling, the band must classically be stripped. By way of non-limiting example, this stripping can be performed in a solution of 20% hydrochloric acid at room temperature ambient, in the presence of hexamethylenetetramine as an inhibitor.

The cold rolling of the strip is then carried out with a rate of total reduction which is a function not only of the desired final thickness, but also of the resistance and the hardness that one wishes to obtain. As indicative, for an iron-carbon-manganese alloy at 0.6% carbon and 22% manganese, the resistance reaches almost 2000 MPa after 60% of reduction, and its hardness Hv 5 under the same conditions practically reaches 700. On the same alloy, a reduction rate of 30% leads to a resistance of about 1500 MPa. Generally, for alloys concerned by the invention, it is possible to propose carrying out cold rolling with a minimum total reduction rate of 25%. A rolling mill can be used conventional cold, or a Sendzimir rolling mill that gives access, in three passes, at reduction rates of the order of 60-70%, including for alloys with very high strength, greater than 1500 MPa. A thickness of 1 mm for the cold-rolled sheet can thus be obtained.

In general, it is advisable to proceed, during the first pass of the cold rolling at a high thickness reduction, of the order of 25% less. Indeed, the heating produced by such a large reduction as soon as The beginning of rolling slows or even inhibits the twisting of the deformation, which facilitates rolling. This pass may even be sufficient to obtain from the outset the intended final thickness.

This is followed by a recrystallization annealing, so as to to obtain an adequate grain size for the control of the compromise resistance / ductility and the ratio Re / Rm (yield strength / resistance to traction). This recrystallization annealing must be carried out by the method of continuous annealing because a basic annealing would lead to a precipitation of carbides, which we saw was undesirable. This annealing can be carried out oxidizing atmosphere, followed by stripping; it can also be of the type "Bright annealing", that is to say carried out in an inert atmosphere, which makes it possible to get rid of pickling and limit surface decarburization. We can do follow this annealing by passing through a cold-rolling mill ("skin-pass") or planing. Typically, this recrystallization annealing is performed at a temperature of 600-1200 ° C, for 1 second to 1 hour, depending the size of the grains that one wishes to obtain.

For example, an iron-carbon-manganese alloy with 0.6% carbon and 22% of manganese may preferably undergo a bright annealing 800 ° C for 90s to obtain a grain size of the order of 2.5 microns. The mechanical characteristics thus obtained are a maximum resistance of 1030 MPa and an elongation at break of 60%.

In general, usable iron-carbon-manganese alloys in the process according to the invention may have an elongation at break of 90% or more, if a relatively low tensile strength of 600 MPa (figures obtained for a 0.2% carbon alloy, 27% manganese, with a grain size of 30 μm). But in the range compositions (0.4 to 0.8% of carbon and 20 to 24% of manganese), an elongation at break of about 50 to 60% and a tensile strength of the order of 1000 MPa for a size of grains of 5 μm, or even a tensile strength of the order of 1200 MPa for a grain size of 1 μm.

In addition to their mechanical characteristics favorable to the uses considered in the invention, these alloys are distinguished by excellent weldability in particular, that they contain optimally little or very little silicon, whose oxide is difficult to reduce, and because of their structure austenitic that makes the concepts of martensitic quenchability irrelevant and / or carbon equivalent which should normally be taken into account when the use of conventional ferritic steels to form small tubes welded. In addition, these alloys can easily receive a deposit uniform and adherent zinc by electrogalvanization, especially in the case where they were cold rolled.

We then proceed to manufacture the small welded tube, using for this, the conventional methods. The success of this manufacturing is in largely conditioned by the cleanliness of the weld. As a result, a excellent preliminary stripping (especially if no annealing has been carried out gloss of the cold-rolled strip) is necessary not to occlude surface oxides in the weld seam.

After slitting the sheet, shearing its shores and progressive forming to bring its edges until docking, the tube is conventionally welded by electrical resistance, laser or high frequencies. The inner and outer scraping of the bead is then carried out welding to eliminate thickness variations. These variations thick would be unfavorable to hydroforming and would damage the tool formatting.

This shaping of the tube can take place by cold drawing. In this Indeed, the thickness of the tube is reduced by pulling through a die which calibrates the outside diameter and, most often, on a mandrel which calibrates the internal diameter. Using suitable dies and chucks, stretching can be used to shape the tubes and transform a draft of circular section into a product having another geometry.

The shaping can also take place by hydroforming. according to this process, a hollow body of more or less complex shape is manufactured in deforming a tube under the joint action of internal pressure and forces compression acting at the ends of the tube. Iron-carbon-manganese alloys used in the invention have a coefficient of hardening of the order of 0.5, which is very favorable to their good behavior during hydroforming, and makes it possible to obtain pieces of complex shape which would be inaccessible by the use of more conventional steels. Only some austenitic stainless steels would likely have comparable performance.

In general, the use of iron-carbon-manganese alloys presenting the compositions indicated gives the metal a great variety behaviors, which allow either to obtain welded tubes with better mechanical characteristics than existing products, either to obtain mechanical characteristics equivalent to those of the products existing, but for a lower cost of production and / or for a quantity of subject matter brought into play less, leading to a significant reduction in the room. Thus, with the alloy with 0.2% carbon and 27% of manganese mentioned above whose tensile elongation exceeds 90%, it is possible to eliminate the anneals intermediaries, and consider increasing stitching heights. As for the alloy at 0.6% carbon and 22% manganese whose resistance to traction is from 1000 to 1200 MPa, it allows to obtain a mass gain important on the final tube and to simplify the control of its stage of implementation. because this high strength widens the loading range by reducing the burst area during hydroforming. Finally, so general, because of their high work hardening capacity, stretching and hydroforming of the iron-carbon-manganese alloys according to the invention have also the advantage of standardizing the mechanical characteristics in every respect of the tube.

Claims (14)

1. Method for making a welded tube, of the type comprising a final drawing or hydroforming step,
   characterised in that:

   an alloy is produced, having a composition, expressed in per cent by weight, of:

   C ≤2%;

   Mn of between 10 and 40%, with Mn% > 21.66 - 9.7 C%;

   Si ≤ S 5%, preferably ≤ 1%, optimally ≤ 0.5%;

   S ≤ 0.3%, preferably ≤ 0.05%, optimally ≤ 0.01%;

   p ≤ 0.1%, preferably ≤ 0.05%;

   Al ≤ 5%, preferably ≤ 0.1%, optimally ≤ 0.03%;

   Ni ≤ 5%, preferably ≤ 2%;

   Mo ≤ 5%, preferably ≤ 1%;

   Co ≤ 3%, preferably ≤ 1%;

   W ≤ 2%, preferably ≤ 0.5%;

   Cr ≤ 5%, preferably ≤ 1%;

   Nb ≤ 1%, preferably ≤ 0.1%;

   V ≤ 1%, preferably ≤ 0. ≤1%;

   Cu ≤ 5%, preferably ≤ 1%;

   N ≤ 0.2%, preferably ≤ 0.1%, optimally ≤ 0.05%;

   Sn ≤ 0.5%, preferably ≤ 0.1%;

   Ti ≤ 1%, preferably ≤ 0.1%;

   B ≤ 0.1%, preferably ≤ 0.01%;

   each of the Ca and Mg contents ≤ 0.1%, preferably ≤ 0.01%;

   each of the As and Sb contents ≤ 0.1%, preferably ≤ 0.05%;

   then a semi-finished product is cast from this alloy,

   a) either in the form of an ingot which is then roughed down by hot rolling to transform it into a slab, or directly in the form of a slab, the said slab then being hot-rolled in the form of a strip and then coiled,

   b) or in the form of a thin strip;

   then the strip is pickled if it is surface oxidised;

   finally the welded tube is made by progressive forming of a metal sheet cut from the preceding strip so as to bring its edges together, then by welding the said edges, then by eliminating the welding bead, then by cold drawing or hydroforming.
2. Method according to Claim 1, characterised in that the carbon content of the alloy is between 0 and 1.2% and in that the manganese content of the alloy is between 10 and 35%.
3. Method according to Claim 2, characterised in that the carbon content of the alloy is between 0.2 and 1.2% and in that the manganese content of the alloy is between 10 and 30%.
4. Method according to Claim 3, characterised in that the carbon content of the alloy is between 0.2 and 0.8% and in that the manganese content of the alloy is between 15 and 30%.
5. Method according to Claim 4, characterised in that the carbon content of the alloy is between 0.4 and 0.8% and in that the manganese content of the alloy is between 20 and 24%.
6. Method according to one of Claims 1 to 5, characterised in that the hot rolling is preceded by reheating performed at a temperature not exceeding 80°C below the solidus temperature of the alloy.
7. Method according to one of Claims 1 to 6, characterised in that the hot rolling is preceded by reheating performed at a temperature at which the precipitation of aluminium nitrides is not brought about.
8. Method according to one of Claims 1 to 7, characterised in that the temperature at the end of hot rolling is greater than or equal to 900°C.
9. Method according to one of Claims 1 to 8, characterised in that the coiling temperature after hot rolling is less than or equal to 450°C.
10. Method according to one of Claims 1 to 9, characterised in that annealing followed by hyperquenching of the coiled hot-rolled strip are carried out, the said annealing being performed in conditions permitting the redissolving of the carbides and avoiding their precipitation on cooling.
11. Method according to one of Claims 1 to 10, characterised in that, after the hot rolling and the possible annealing followed by a possible hyperquenching, cold rolling of the strip is carried out, with a minimum reduction ratio of 25%, preceded by pickling.
12. Method according to Claim 11, characterised in that the reduction ratio of the thickness of the strip during the first cold-rolling pass of the strip is at least 25%.
13. Method according to one of Claims 11 and 12, characterised in that recrystallisation annealing of the strip is carried out at a temperature of 600 to 1200°C for 1 second to 1 hour.
14. Welded tube, characterised in that it has been made by the method according to one of Claims 1 to 13.

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