EP0198024B1 - Procede de fabrication d'un acier pour precontrainte - Google Patents

Procede de fabrication d'un acier pour precontrainte Download PDF

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EP0198024B1
EP0198024B1 EP85905199A EP85905199A EP0198024B1 EP 0198024 B1 EP0198024 B1 EP 0198024B1 EP 85905199 A EP85905199 A EP 85905199A EP 85905199 A EP85905199 A EP 85905199A EP 0198024 B1 EP0198024 B1 EP 0198024B1
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process according
hardening
strength
grain
steel
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EP0198024A1 (fr
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Max Willy Tischhauser
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Priority claimed from CH5210/84A external-priority patent/CH667104A5/de
Priority claimed from DE19853535886 external-priority patent/DE3535886A1/de
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/06—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of rods or wires
    • C21D8/08—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of rods or wires for concrete reinforcement

Definitions

  • the invention relates to a method for producing high-strength, weldable, corrosion-resistant and brittle fracture-proof prestressing steels.
  • prestressing steels are usually made from unalloyed, high-carbon, high-grade structural steels
  • patented or style-treated wire rod with dimensions of 5.5 to 14.5 mm round with a composition of 0.60 to 0.90 C, 0.10 to 0.30 Si, 0.50 to 0.80 Mn, 0.035 S and 0.035 P. This is used to make cold drawn tension wire.
  • semi-finished billets of approx. 120 mm 4-kt are used as primary material, which is heat-treated according to different criteria depending on the manufacturer and existing systems, i.e. is brought to the rolling temperature and therefore also has different microstructures and properties, but in the end product it must have the mechanical properties customary for registration certificates.
  • prestressing steels have the considerable disadvantage that they cannot be welded.
  • Conventional processes are used to manufacture them, such as the well-known Siemens Martin, electric furnace or oxygen inflation process, with the steel being treated neither before nor after. If at all, steel pretreatment by desulfurization and steel post-treatment by vacuum treatment take place, if at all. Block and continuous casting are still used as the casting process.
  • prestressing steels In addition to the lack of weldability, these known prestressing steels have deficiencies with regard to their mechanical properties, susceptibility to corrosion and, in particular, insensitivity to brittle fracture, in spite of hardly any changed conceptions with regard to their chemical composition, structure and manufacturing conditions.
  • a fact that has so far been overlooked in the assessment of prestressing steel is that the susceptibility of prestressing steel to brittle fracture can begin considerably above 0 ° C and increases rapidly at lower temperatures. The safety against brittle fracture is expressed with the so-called transition temperature for possible brittle fracture.
  • Conventional prestressing steels usually have a door significantly above + 20 ° C !.
  • Corrosion occurs in prestressing steel in a variety of forms, be it trough, hole, gap, intercrystalline and transcrystalline corrosion. Especially. Attention should be paid to stress corrosion cracking.
  • the corrosion-inhibiting properties of copper are known, but copper has so far not been used as an alloying element in prestressing steels.
  • the reason for keeping the steel at the lowest possible reheating temperature is that vanadium and niobium dissolve at 850 ° C and 950 ° C, but are dissolved again above 1150 ° C. The latter should be avoided.
  • a particle size of 100-200 ⁇ and a particle quantity of 20 x 10 6 per mm 2 should be achieved for the intended purpose.
  • thermomechanical treatment produces wire rod grades for the production of cold drawn wire, three-wire strands, seven-wire strands and tension rods, which correspond to the Euro standard 138 in their properties, however, the additional usage properties (corrosion-resistant, brittle-break proof and weldable). There is no need for costly cold forming (stretching) and subsequent tempering for tensioning rods, which already means a considerable advantage of the invention.
  • a third stage of the treatment can also be provided, in which, starting at about 650/550 ° C., rolling is again carried out in a controlled manner with one or a few passes, that is to say with a high degree of deformation at high speed.
  • a dwell time and a delayed cooling down for example in still air, are considered.
  • an increased precipitation hardening process results in an increase in strength of over (at least) 40% compared to conventional prestressing steels.
  • the accompanying diagram serves to illustrate this process sequence.
  • the steel can also be strain hardened, provided that higher strength classes are aimed for or are necessary.
  • thermomechanical treatment in accordance with the present invention, the mechanisms of increasing the strength due to the chemical composition and the targeted metering of the microalloying elements interact in an additive manner. These mechanisms are in particular fine grain hardening, mixed crystal hardening and very particularly precipitation hardening, in which the alloying element copper is particularly effective.
  • thermomechanical treatment along with the chemical composition for melting and hardening fine grains, is the most important step in achieving the desired goal, namely for the production of high-strength, corrosion-resistant, brittle-breakable and weldable prestressing steels.
  • the dosage of the alloy elements is designed so that not only the strength is increased considerably, but also the toughness is increased at the same time, in particular through the fine grain hardening.
  • the targeted metering of the alloying elements also ensures that the highest degree of solidification takes place via precipitation hardening. Elimination in ferrite is the most effective for increasing strength.
  • this phase of the thermomechanical treatment is also the highest Importance, because the targeted dosing of the alloying elements also achieves the highest safety against brittle fracture, in particular through the interaction of the elements manganese and molybdenum.
  • Precondition for an effective increase in strength in the sense of the invention is furthermore fine-grain hardening, fine-grain melting being necessary to achieve it optimally, which simultaneously increases toughness.
  • the grain size to be achieved according to ASTM 112 should be at least 9, but if possible at least 12, to which an increased manganese content of 1.45% on average contributes.
  • an austenite grain that is as fine as possible should already be sought, since this also determines the size of the ferrite grain.
  • the microalloying elements provided in the directional analysis in particular aluminum, nitrogen, niobium and vanadium, are used to inhibit grain growth and to form strengthening obstacles to the dislocations through fine precipitations in the austenite structure.
  • a particle size of 100 to 200 ⁇ is most effective for this, the particle quantity per mm 2 approx. 20. 10 6 should be.
  • Continuous casting should be a suitable type of casting. Continuous casting is the most economical and, at the same time, the best quality of casting and solidifying the molten steel into the primary material used to manufacture prestressing steel: billets. To ensure a high * level of quality required for prestressing steel, special measures must be taken to prevent such defects, such as e.g. Protection against reoxidation, concealed casting, electromagnetic stirring.
  • the low carbon content of 0.05-0.20% provided in the directional analysis largely prevents the occurrence of the abovementioned faults and at the same time favors the economy of continuous casting for the production of prestressing steel grades by the costly measures to a greater extent, such as for the conventional high-carbon prestressing steel grades are not required, while at the same time ensuring a high degree of purity, homogeneity and quality.
  • fine grain hardening has to be given the most consideration because the hardening mechanism resulting from it is characterized not only by an increase in strength but also a simultaneous increase in toughness. Furthermore, the two-dimensional obstacles to moving dislocations are such strong obstacles that they cannot be overcome by them. The dislocation has then become impossible and numerous dislocations form a build-up at the grain boundary, which results in a significant stress concentration and therefore an influence on the strength. However, the average grain size influences the lower yield strength.
  • these individual hardening mechanisms can be combined with one another and with a specific work hardening, their effect being additive, but their respective proportions can change considerably depending on the specified conditions. According to the invention, however, it was found that the basic mechanisms of the individual hardenings can only be achieved by a further, the most important treatment - step, become optimal, namely through the so-called thermomechanical treatment.
  • thermomechanical treatment within the scope of the invention is carried out by a very specific sequence of controlled rolling of the microalloyed and fine-grained steel specifically developed for this purpose, in particular a low final rolling temperature, rapid cooling before the last rolling pass and a high degree of final deformation, so that recrystallization occurs leads to the finest possible austenite grain before the ferrite-pearlite transformation.
  • the rolling process additionally results in precipitation of carbides, nitrides or carbonitrides, as well as solid-crystal and fine-grain and particle hardening.
  • the temperature control is controlled by alloying and rolling technology in such a way that the ⁇ -a conversion takes place shortly before and / or after the lowest possible final rolling temperature which comes to be just before A, 3 . In any case, the formation of martensite should be excluded.
  • the final rolling temperature and the degree of deformation, particularly in the last pass, are decisive for the mechanical properties that can be achieved.
  • the pearlite content decreases, which means that low-carbon, micro-alloyed microstructures have only a small, often no pearlite content in the structure in a controlled final-rolled state.
  • the mechanical properties experience an additional favorable influence.
  • thermomechanical treatment The last stage of the process following the thermomechanical treatment is work hardening, which consists in particular of stretching or drawing. Through this subsequent cold working, which is used to manufacture all prestressing steels and for which the steels of the new design are particularly well suited, a considerable increase in strength is achieved compared to today's prestressing steel grades by means of the degree of deformation to be used.
  • microalloying elements In order to further improve the properties of the prestressing steels according to the invention in connection with the thermomechanical treatment, the additions of microalloying elements can be ascribed.
  • niobium has the most effective influence on fine grain hardening and hardening through thermomechanical treatment, i.e. the strength increase, followed by vanadium. The same applies to the improvement of the transition temperature.
  • Microalloying with niobium and vanadium increases the strengthening proportion of the manganese and silicon content with increasing levels while at the same time reducing pearlite.
  • niobium alloys result in a much larger proportion of fine-grain hardening than hardening and therefore not only a higher yield strength than with a titanium or vanadium alloy, but above all, as already mentioned, a very favorable low transition temperature.
  • the high ratio of fine grain hardening to hardening due to the addition of niobium is therefore a major reason why niobium must be used here, since niobium also brings about the greatest reduction in the transition temperature.
  • Manganese and silicon at levels below about 0.5% also shift the transition temperature to lower temperatures.
  • the sulfur content plays a decisive role in the anisotropy of toughness, the most important factor influencing cold formability.
  • a desired lower sulfur content ie a reduced number of sulfide inclusions, significantly improves the toughness with regard to constriction of the fracture, a property that is particularly important for prestressing steels.
  • the reduction in the sulfide length is particularly effective for a more favorable constriction of the fracture.
  • molybdenum-niobium-alloyed structural structures give the best properties.
  • An additional improvement in the properties is achieved by the combination of niobium-vanadium-molybdenum-copper with simultaneous thermomechanical treatment according to the invention, the best results being achieved by using a low final rolling temperature and the highest possible degree of final deformation.
  • the most effective way to achieve optimal mechanical properties is to produce a fine grain.
  • the refinement of the grain size leads to an increase in the yield strength with a simultaneous improvement in the transition temperature.
  • an austenite grain that is as fine as possible is sought, since this also determines the size of the ferrite grain.
  • a reduction in the austenite grain size has a factor of around 0.3 that affects the reduction in the ferrite grain size.
  • the essential process in the growth of the austenite grain is not the dissolution of the excretions, but their aggregation into large and thus effective particles.
  • One measure for controlling the austenite grain size is the incorporation of fine precipitates in the austenite structure, which inhibits grain growth.
  • aluminum which produces this effect via aluminum nitride
  • it is primarily the microalloying elements niobium and vanadium in particle sizes from 100 to 200 ⁇ that have a comparable effect via their carbides, nitrides or carbonitrides.
  • the most favorable conditions for preventing the sharp increase in grain growth when reheating in the pusher furnace for rolling show higher aluminum contents (up to 0.050%) and nitrogen contents (up to 0.020%). With increasing niobium content, the start of the sudden grain growth is also shifted to higher temperatures.
  • austenite grain can also be refined by higher degrees of deformation. The grain refinement effect is most pronounced at low final deformation temperatures.
  • the lower transformation temperature causes a higher nucleation frequency and a lower mobility of the grain boundaries, which results in a reduction in the ferrite grain size.
  • alloying small amounts of molybdenum to the micro-alloyed, low-pearlite structure can also favor the delay in austenite recrystallization, which shifts the--a conversion to lower temperatures. It is precisely this possibility that is used in thermomechanical treatment, as a result of which an even more fine-grained structure is achieved with a simultaneous additional improvement in the transition temperature.
  • incoherent niobium and vanadium carbonitrides act in different particle sizes and quantities on the ferrite grain size.
  • vanadium In the thermomechanically treated state, vanadium only causes a slight grain refinement.
  • the basic composition plays a role in that higher carbon and nitrogen contents result in a finer secondary structure via a stronger or faster excretion before or during the ⁇ -a conversion. It should also be noted that the optimal grain refinement due to niobium contents between 0.04 and 0.10% is equally effective, but that of vanadium is also increasingly effective with increasing contents.
  • the carbon and nitrogen content of the steel influences the ferrite grain size significantly less in steels with niobium than in those with vanadium. With decreasing carbon contents, the influence of nucleation due to separated particles on the grain size decreases in favor of a very pronounced and, in the present case, desirable recrystallization inhibition due to dissolved niobium. Low-pearlite steels therefore have smaller ferrite grain sizes in the thermomechanically treated state than steels with a higher carbon content.
  • Dissolved vanadium or niobium or titanium cause a further fine grain effect by delaying the austenite transformation desired here. Rising manganese contents also lower the transition temperature, ensure optimal particle separation and thus the optimal effect of particle hardening.
  • the carbon content causes a substantial solidification via the cementite (pearlite) and plays an important role in this connection.
  • the carbon content via the pearlite component has the most significant negative influence on the brittle fracture safety (transition temperature) also specified in this development and on the weldability, and increasingly with increasing pearlite component, the carbon content on components increases limit, which allow both an increase in strength and improvement of corrosion resistance, but also improve the brittle fracture safety down to around -40 ° C and the weldability.
  • the optimal fine grain formation to be aimed at it should also be taken into account that the carbon content has a considerable influence on this.
  • Manganese has a particularly fine grain refinement and, at the same time, due to solidification of the solid solution and increased hardening, so that the manganese--
  • the content should preferably be arranged at the upper limit because the increase in strength due to manganese is very strongly dependent on the pearlite content and, thanks to a suitably low pearlite content, also ensures a favorable transition temperature and thus also safety against brittle fracture. Rising manganese contents make a considerable contribution to delaying the austenite transformation desired here and thus cause optimal fine grain formation. With the simultaneous presence of niobium and vanadium as microalloying elements, the increasing solidifying proportion of manganese becomes effective in the case of low-pearlite structures with increasing manganese content.
  • the latter for manganese also applies to the silicon content. If the silicon content is below about 0.5%, the transition temperature is also shifted to lower temperatures. However, silicon also has a strengthening effect above 0.5%, but at the same time is becoming increasingly brittle, which is to be avoided here for prestressing steels.
  • Niobium has the most effective influence on fine grain hardening and hardening through thermomechanical treatment, i.e. on the achievable strength increase, followed by vanadium. It causes the greatest reduction in the transition temperature.
  • the niobium-containing structure results in a much larger proportion of fine-grain hardening than in hardening and therefore not only a higher yield strength than is achieved by vanadium alloy structure structures, but above all a very favorable, low transition temperature.
  • Niobium reduces the ferrite grain size to a particularly large extent.
  • Niobium has the additional strengthening effect of increasing manganese contents even when pearlite is low.
  • vanadium forms precipitates of special carbides, which on the one hand contribute to fine grain formation and hardening and on the other hand to precipitation hardening and thus significantly increase strength. Like niobium, vanadium therefore helps to control the austenite grain size by embedding fine precipitates in the austenite structure, which inhibits grain growth. Like niobium, vanadium contributes to solid-solution strengthening, but both are insoluble in ferrite. Their elimination in ferrite is therefore the most effective for increasing strength.
  • the carbides and nitrides of vanadium and niobium have face-centered cubic lattices, are isomorphic and therefore completely miscible.
  • vanadium In contrast to titanium, they do not contribute to the formation of sulfide. With an increased nitrogen content, vanadium has the greatest influence on the formation of a fine ferrite grain size and causes an additional increase in the yield strength. Like niobium, dissolved vanadium affects this fine grain effect and hardening by delaying the austenite transformation.
  • the delay in austenite recrystallization is greatly promoted by adding small amounts of molybdenum to the microalloyed, low-pearlite structure, which shifts the o-a conversion to lower temperatures.
  • This possibility is used in the thermomechanical treatment by an even lower final forest temperature, whereby an even more fine-grained structure is achieved with a simultaneous improvement in the transition temperature.
  • alloying with molybdenum and the resulting possibility of shifting the ole-a conversion to lower temperatures it is also additionally possible to take full advantage of the considerable strengthening properties of copper.
  • both hardening mechanisms work both through the precipitation of mixed crystals and through the formation of carbonitrides, especially at temperatures between 650 and 550 ° C.
  • Copper is used for the purpose envisaged here because of its two advantages. First, because of its strong hardening effect through hardening. Second, because of its strong anti-corrosive effects.
  • the corrosion-inhibiting effect of copper can be used particularly well in high-strength microstructures created with thermomechanical treatment, because at the low final rolling temperatures, which also lead to the highest strength increases, the element copper simultaneously with the precipitation-hardening elements used here between 650 and 550 ° C, in addition to its corrosion-inhibiting effect, also acts as a precipitation-hardening element. By rapid cooling from the 6 region at approx. 840 ° C, around 2% copper can be dissolved in low-pearlite microstructures and the thermomechanical treatment already provided here.
  • a copper-rich, cubic surface-centered mixed crystal is then separated out in the form of incoherent, spherical particles, which, from a certain particle size, leads to a considerable precipitation hardness effect by the bypass mechanism.
  • niobium is present, micro-alloyed structures with a low pearlite content and a high copper content both cause hardening mechanisms due to the precipitation of mixed crystals and carbonitrides.
  • a nickel content of up to 1% must be added to the copper-alloyed microstructures in order to increase the solder fragility caused by copper prevent.
  • an additional increase in strength can be achieved with a copper alloy structure apart from particle hardening through a high dislocation density and fine grain hardening.
  • the corrosion-inhibiting effect of copper is already very effective with a very low copper content (0.25 to 0.40%). It is therefore necessary to adjust the copper content in order to be able to optimally use the corrosion-inhibiting effect and the hardening mechanisms on the one hand, but on the other hand not to let the solder fragility, which would be unsustainable for prestressing steels, come into effect and, if possible, to avoid adding nickel to this prevention .
  • the sudden grain growth is increased to around 1150 ° C when the primary material is heated, whereby the holding time is also important.
  • aluminum which produces this effect via aluminum nitride, it is primarily the microalloying elements niobium and vanadium that have a comparable effect via their carbides, nitrides or carbonitrides.
  • the lowest possible pusher furnace temperature is essential. The most favorable conditions for preventing the sharp increase in grain growth when reheating for rolling show higher aluminum contents.
  • Aluminum also contributes to solidification of the solid solution.
  • the sudden grain growth before heating for rolling is also increased by nitrogen to higher temperatures of around 1150 ° C.
  • An increased nitrogen content also makes a significant contribution to increasing the strength by increasing the nitride content.
  • vanadium is present, the yield strength increases significantly. This also increases the tensile strength, so that an increase in the yield strength ratio from 70% to 90%, which is particularly important for prestressing steels, is brought about.
  • the lowest possible phosphorus content is of particular importance and should therefore be aimed for.
  • the sulfur content plays a decisive role in the anisotropy of toughness, the most important factor influencing their cold formability for prestressing steels.
  • a lower sulfur content i.e. a reduced number of sulfide inclusions improves the toughness significantly with regard to fracture constriction, a property that is particularly important for prestressing steels.
  • the reduction in the sulfide length is particularly effective for a more favorable constriction of the fracture.
  • a strong desulfurization can be achieved by the calcium additions that are common in ladle metallurgy.
  • titanium in contrast to niobium and vanadium, it participates in the formation of sulfide. On the other hand, it first binds all nitrogen to nitrides, TiN, and then sulfur to a titanium carbosulfide, Ti 4 C 2 S 2 .
  • titanium is not taken into account here, since, among other things, the effect of one of the austenite grain growth and that of an increase in strength in interaction with the other microalloying elements would be offset by an increased nitrogen content.
  • the coupling links are the most sensitive weak points for the occurrence of damage caused by the penetration of corrosive media up to the steel. According to today's technical possibilities, such coupling elements are usually arranged at too short intervals from one another. The resulting high number of coupling joints results in a high number of weak points.

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Abstract

Dans un procédé de fabrication d'un acier de précontrainte présentant des qualités élevées de résistance à la traction, à la corrosion et à la rupture, on produit une trempe à grain fin ou à cristal mixte, ou par précipitations, conjointement avec un traitement thermomécanique suivi d'un durcissement à froid. Pour le traitement de durcissement, on utilise une trempe à grain fin ou à cristal mixte, ou à précipitations avec effet additionnel prononcé. Le traitement mécanique est effectué au moyen de cylindres en alliage d'acier, à grain fin, ce qui exclut la formation de martensite.

Claims (25)

1. Procédé de fabrication d'aciers de précontrainte hautement résistants, soudables, plus résistants à la corrosion et moins fragiles à la rupture, caractérisé en ce qu'on soumet un acier constitué de
0,05 à 0,20% en masse de carbone
1,20 à 1,70% en masse de manganèse
0,30 à 0,50 en masse de silicium
0,04 à 0,06% en masse de niobium
0,035 à 0,05% en masse de vanadium
0,30 à 0,50% en masse de molybdène
0,30 à 2% en masse de cuivre
0,04 à 0,06% en masse d'aluminium
0,015 à 0,02% en masse d'azote
50,030% en masse de phosphore
:50,20% en masse de soufre
à un traitement thermo-mécanique qui intervient après solidification du métal en fusion et après nouveau chauffage par deuxième passage au four, l'acier étant maintenu, au cours d'une première étape, avant le traitement thermo-mécanique, à une température de réchauffage aussi basse que possibe (= deuxième passage au four à moins de 1 150°C), puis un laminage contrôlé de l'acier étant effectué, avec un nombre réduit de passes, à un degré de déformation élevé (10 à 45%) et une grande vitesse de déformation, jusqu'à une basse température de déformation, à peine supérieure à 850°C.
2. Procédé selon la revendication 1, caractérisé en ce que le traitement thermo-mécanique comporte une deuxième étape dans laquelle, à partir de 850°C environ, on procède à un refroidissement accéléré sand laminage à environ 650/ 550°C, ce qui entraîne une reduction de la transformation õ­α, et un retard de recristallisation.
3. Procédé selon la revendication 2, caractérisé en ce que le traitement thermo-mécanique comporte une troisième étape au cours de laquelle, à partir de 650/550°C environ, on contrôle une nouvelle fois par une ou quelques passes, c'est-à-dire qu'on lamine à un degré de déformation élevé à une grande vitesse, à une basse température de laminage de finition, à peine au-dessus de la limite Ar3, puis on procède à un refroidissement retardé, après un temps d'arrêt.
4. Procédé selon au moins une des revendications 1 à 3, caractérisé en ce que pendant le traitement thermo-mécanique, il se produit un écrouissage de l'acier par étirage (pour les aciers de précontrainte) ou par traction (pour les fils étirés à froid).
5. Procédé selon au moins une des revendications 1 à 4, caractérisé en ce qu'on désulfure largement l'acier avant et/ou après affinage.
6. Procédé selon la revendication 5, caractérisé en ce que l'acier en fusion est soumis à un traitement au calcium avant et/ou après affinage.
7. Procédé selon la revendication 5 ou 6, caractérisé en ce qu'on procédé en outre à un traitement ultérieur de l'acier, par exemple à un lavage au gaz inerte, à un traitement sous vide, à une désoxydation, à une modification des inclusions ou à un traitement en poche avec du calcium métallique ou des scories d'halogénure de calcium.
8. Procédé selon une des revendications 1 à 7, caractérisé en ce qu'on utilise comme mécanisme de durcissement pendant le traitement thermo-mécanique, un durcissement de la phase homogène solide, un durcissement des grains fins et des particules ou un durcissement par précipitation, avec effets s'additionnant largement.
9. Procédé selon une des revendications 1 à 8, caractérisé en ce que le traitement thermo-mécanique s'effectue par laminage contrôlé d'aciers microalliés, à grains fins, et exclut une formation de martensite.
10. Procédé selon la revendication 9, caractérisé en ce que la recristallisation de ces aciers micro- allies à grains fins conduit à un grain d'austénite aussi fin que possible avant la transformation ferrite-perlite.
11. Procédé selon la revendication 10; caractérisé en ce que les aciers micro-alliés, on complète le laminage par séparation de carbures, de nitrures et/ou de carbonitrures, ce qui provoque un durcissement de la phase homogène solide, mais aussi un durcissement des grains fins et en particulier à un durcissement renforcé des particules.
12. Procédé selon l'une des revendications 2 à 11, caractérisé en ce que la température est contrôlée de manière qu'une transformation δ-a se produise juste avant et/ou pendant la température de laminage de finition la plus basse possible, qui se situe juste avant A,3.
13. Procédé selon la revendication 1, caractérisé en ce que la première étape du traitement thermo-mécanique, on obtient des produits en fil laminé pour la fabrication de torons à trois et sept fils ainsi que de barres de précontrainte dont les caractéristiques sont conformes à l'Euronorme 138, mais que présente en plus à l'usage des propriétés supplémentaires de résistance à la corrosion, de non fragilité à la rupture et de soudabilité.
14. Procédé selon l'une des revendications 1 à 13, caractérisé en ce que l'ajout des éléments de micro-alliage niobium et/ou vanadium et/ou molybdène, provoque un durcissement optimal possible des particules sous forme de carbures, de nitrures, et/ou de carbonitrures, par séparation pendant le traitement thermo-mécanique, en plus du durcissement de grains fins et de phase homogène solide.
15. Procédé selon l'une des revendications 1 à 14, caractérisé en ce qu'on procède à un affinage du grain d'austénite par inclusion de fins dépôts tels que des nitrures d'aluminium ou des carbures, des nitrures et/au des carbonitrures, en particulier des éléments de micro-alliage niobium et vanadium, et ce à raison de 20 x 106 particules par mm2, de diamètre compris entre 100 et 200 Ä, avec des degrés de déformation et des vitesses aussi élevés que possible et une température de laminage de finition aussi basse que possible.
16. Procédé selon l'une des revendications 1 à 15, caractérisé en ce que par un retard de recristillisation, on déforme des parties d'austénite non recristallisées pendant les basses températures de laminage de finition, de sorte qu'il en résulte des grains allongés et par conséquent des surfaces de grains d'austénite agrandies, leur transformation dans la phase ferrite-perlite donnant lieu à un fort affinage des grains, par une densité accrue des germes et par la croissance inhibée des grains formés à partir de ces germes, et par conséquent à une augmentation optimale de la résistance, par durcissement des grains fins ainsi que des particules.
17. Procédé selon une des revendications 14 à 16, caractérisé en ce que la recristallisation de l'austénite, outre le contrôle de la vitesse de refroidissement par addition de molybdène, et par conséquent la transformation δ-a - est déplacée vers des témperatures plus basses.
18. Procédé selon une des revendications 14 à 17, caractérisé en ce que par addition d'une proportion accrue de manganèse, dans les limites de l'analyse donnée à titre indicatif, on garantit un affinage des grains, comme recherché, et en même temps une augmentation optimale de la résistance, par durcissement de la phase homogène solide et durcissement renforcé par précipitation.
19. Procédé selon la revendication 18, caractérisé en ce qu'en augmentant la proportion de manganèse, on garantit en même temps une retard optimal de la transformation d'austénite recherchée, et de ce fait la formation optimale de grains fins.
20. Procédé selon la revendication 18 ou 19, caractérisé en ce qu'en augmentant la proportion de manganèse, on garantit en même temps le retard de recristillisation recherché, et ce par déplacement de la transformation 5-a vers des témperatures plus basses et réglage de la témpe- rature de laminage de finition la plus basse possible, et utilisation simultanée du traitement thermo-mécanique.
21. Procédé selon une des revendications 18 à 20, caractérisé en ce qu'en augmentant la proportion de manganèse, on garantit la séparation optimale de particules et donc l'effet optimal du durcissement des particles envue d'une augmentation maximale possible de la résistance.
22. Procédé selon une des revendications 18 à 21, caractérisé en ce qu'en augmentant la proportion de manganèse et en présence en même temps de niobium et de vanadium et avec peu der perlite, on accroît aussi la proportion de manga- nése durcissant et par conséquent on augmente la résistance.
23. Procédé selon l'une des revendications 1 à 22, caractérisé en ce qu'on abaisse la proportion de perlite.
24. Procédé selon l'une des revendications 1 à 23, caractérisé en ce que pour empêcher la croissance rapide des grains lors du préchauffage dans le four poussant ou similaire, on utilise des proportions accrues d'aluminium et d'azote dans les limites de l'analyse donnée à titre indicatif, tandis que dans ce procédé et à cet effet, on recherche une quantité de particules de 20 x 106 /MM 2 et un diamètre de particules compris entre 100 et 200 A.
25. Procédé selon l'une des revendications 1 à 24, caractérisé en ce qu'on empêche le début de la croissance des grains lors des températures plus élévées, en augmentant la proportion de niobium à l'intérieur des limites de l'analyse donnée à titre indicatif.
EP85905199A 1984-10-30 1985-10-30 Procede de fabrication d'un acier pour precontrainte Expired - Lifetime EP0198024B1 (fr)

Priority Applications (1)

Application Number Priority Date Filing Date Title
AT85905199T ATE51897T1 (de) 1984-10-30 1985-10-30 Verfahren zum herstellen von spannstaehlen.

Applications Claiming Priority (4)

Application Number Priority Date Filing Date Title
CH5210/84 1984-10-30
CH5210/84A CH667104A5 (de) 1984-10-30 1984-10-30 Verfahren zum herstellen von spannstaehlen.
DE3535886 1985-10-08
DE19853535886 DE3535886A1 (de) 1985-10-08 1985-10-08 Verfahren zum herstellen von spannstaehlen

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EP0198024A1 EP0198024A1 (fr) 1986-10-22
EP0198024B1 true EP0198024B1 (fr) 1990-04-11

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EP (1) EP0198024B1 (fr)
KR (1) KR930009973B1 (fr)
AU (1) AU4966585A (fr)
BR (1) BR8507018A (fr)
DE (1) DE3577109D1 (fr)
FI (1) FI862784A0 (fr)
NO (1) NO862605L (fr)
WO (1) WO1986002667A1 (fr)

Families Citing this family (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US5187250A (en) * 1989-06-05 1993-02-16 Mitsui Toatsu Chemicals, Incorporated Poly-α-olefins
DE4224222A1 (de) * 1992-07-22 1994-01-27 Inst Stahlbeton Bewehrung Ev Baustahl, insbesondere Betonstahl und Verfahren zu seiner Herstellung
CH687879A5 (de) * 1993-12-01 1997-03-14 Met Cnam Paris Max Willy Tisch Armierungs-, Maschinen-, Apparate- und Metallbaustaehle in Feinkornguete mit stabiler Korrosionsschutzschicht.
BR9509829A (pt) * 1994-11-28 1997-09-30 Tischhauser Max Willy Processo para a preparação de aços de construção de elevada qualidade de armação de máquinas de aparelhos e de metal em qualidades de grãos finos e com camada de proteção contra corrosão estável
GR1005389B (el) 2005-11-22 2006-12-15 Powerwave Technologies Inc. Εξυπνος στυλος

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* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
GB1001233A (en) * 1961-01-23 1965-08-11 Bernhard Matuschka Improvements in or relating to steel
US3328211A (en) * 1963-12-05 1967-06-27 Ishikawajima Harima Heavy Ind Method of manufacturing weldable, tough and high strength steel for structure members usable in the ashot-state and steel so made
FR1386441A (fr) * 1963-12-11 1965-01-22 Creusot Forges Ateliers Barres en acier pour béton précontraint
AT271532B (de) * 1964-02-20 1969-06-10 Krupp Ag Huettenwerke Stahl für langzeitig unter Zugbeanspruchung stehende Bauelemente, insbesondere Spannbetonstahl
LU51509A1 (fr) * 1966-07-07 1968-03-12
GB1321304A (en) * 1970-05-09 1973-06-27 Exors Of James Mills Ltd Thermo-mechanical treatment of steel
US4299621A (en) * 1979-07-03 1981-11-10 Henrik Giflo High mechanical strength reinforcement steel

Also Published As

Publication number Publication date
BR8507018A (pt) 1987-01-06
NO862605L (no) 1986-08-27
EP0198024A1 (fr) 1986-10-22
FI862784A (fi) 1986-06-30
KR930009973B1 (ko) 1993-10-13
AU4966585A (en) 1986-05-15
NO862605D0 (no) 1986-06-27
WO1986002667A1 (fr) 1986-05-09
DE3577109D1 (de) 1990-05-17
FI862784A0 (fi) 1986-06-30
KR887000089A (ko) 1988-02-15

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