WO1986002667A1 - Method for producing prestressed steel - Google Patents

Abstract

In a method for producing prestressed steel having high rupture and tensile strength, and corrosion resistance properties, there is provided a fine-grained or mixed-crystal or precipitated-phase hardening, in conjunction with a thermomechanical treatment followed by a work hardening. For the hardening process, there is used a fine-grained or mixed-crystal or precipitated-phase hardening with enhanced additional effect. The mechanical treatment is carried out by means of cylinders of fine-grained alloy steel, thereby excluding the formation of martensite.

Description

Method of manufacturing prestressing steel

The invention relates to a method for producing high-strength, weldable, corrosion-resistant and brittle fracture-proof prestressing steels.

At present, prestressing steels are generally produced from unalloyed, high-carbon, high-grade structural steels, specifically

hot-rolled, stretched and tempered rods of dimensions 15 to 40 mm round with a composition of 0.65 to 0.85 C, 0.65 to 0.85 Si, 1.10 to 1.70 Mn, 0.035 S, 0.035 P and optionally 0.10 to 0.40 V as well

- Patented or stilt-treated wire rod of dimensions 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. Cold drawn tensioning wire is produced from this.

In both cases, semi-finished billets from Approx. 120 mm 4-ct is used, which, depending on the manufacturer and existing systems, is heat-treated according to different criteria, ie brought to the rolling temperature and therefore also has different microstructures and properties, but in the end product it has the usual mechanical properties for approval certificates got to.

These prestressing steels have the considerable disadvantage that they cannot be welded. Conventional processes are used for their production, such as, for example, the known Siemens Martin, electric furnace or oxygen blowing process, 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. When casting find ^'still block and continuous casting application.

In addition to the lack of weldability, these known prestressing steels have deficiencies regarding the mechanical properties of the susceptibility to corrosion and, in particular, the insensitivity to brittle fracture, in spite of hardly any significantly changed concepts with regard to their chemical composition, structure and manufacturing conditions. A fact that has so far been overlooked in the assessment of prestressing steels is that the susceptibility of prestressing steels 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 Tu of significantly above + 20 ° C !. As with most prestressed steel structures Temperatures of up to - 40 ° C and more can occur regularly over months, particularly in the case of bridge substructures, this fact must be taken into account when designing and developing prestressing steel. The susceptibility to brittle fracture lies largely in the internal degree of purity, in oxidic and sulfidic inclusions and inclusion forms, which today can largely be improved by targeted steel aftertreatment. Then there is the likelihood of brittle fracture and, above all, its temperature dependence in the closest connection with the pearlite (cementite) part in the steel, that is to say with the carbon content which has the greatest negative influence. To date, there are no pearlitic, ie low-carbon prestressing steels.

Corrosion occurs in prestressing steel in a variety of forms, whether as trough, hole, gap, intercrystalline and transcrystalline corrosion. Particular attention should be paid to stress corrosion cracking. The corrosion-inhibiting properties of copper are known, however

Copper as an alloying element has so far not been used in prestressing steels.

So far, it has not been possible to produce high-strength, at the same time corrosion-resistant and, at the same time, brittle fracture-proof prestressing steels which are also suitable for welding. The inventor set himself the task of developing such a prestressing steel and at the same time developing a method for its production.

A method leads to the solution of this task a steel consisting of

0.05 to 0.20 mass% carbon

1.20 to 1.70% by mass of manganese

0.30 to 0.50 mass% silicon

0.04 to 0.06% by mass of niobium

0.035 to 0.05 mass% vanadium

0.30 to 0.50 mass% molybdenum

0.30 to 2.00 mass% copper

0.04 to 0.06 mass% aluminum

0.015 to 0.02 mass% nitrogen

<£ 0.030 mass% phosphorus 0.020 mass% sulfur

is subjected to a thermomechanical treatment, which takes place after solidification from the melt and reheating from a second heat, wherein

before the thermomechanical treatment, the steel is kept at the lowest possible reheating temperature (= second heat below 1150 ° C.) and subsequently a controlled rolling of the steel with a small number of passes and a high degree of deformation ( 10-45%) and a high ^Q mforrπungsgeschwindigkeit to a low Umforrr.ungsteinperatur closely above 850 ° C leads is durchge.

The reason for keeping the steel at the lowest possible reheating temperature is that vanadium and niobium go into solution at 850 ° C. and 950 ° C., but are dissolved again above 1150 ° C. The latter should be avoided. One should

Particle size of 100-200 A and a particle quantity

6 2 of 20 x 10 per mm can be achieved for the intended purpose.

This is followed by controlled rolling of the steel with a low number of passes at a high degree of deformation (10-45%) and a high deformation speed up to the lowest possible forming temperature, which is close above 850 ° C. The- temperature limit uss due to the copper present in the steel, since effective solidifying deposition of copper can only be achieved by accelerated cooling from approx. -850 ° C to around 650/550 ° C without rolling. that at a temperature below 850 ° C there is no longer any precipitation of copper during rolling.

This first stage of the thermomechanical treatment produces wire rod grades for the production of cold-drawn wire, three-wire strands, seven-wire strands and tension rods, which correspond in their properties to the Euro standard 138, but the additional usage properties (more corrosion-resistant , resistant to brittle fracture and weldable). Costly cold forming (stretching) and subsequent tempering are omitted for tensioning rods, which already means a considerable advantage of the invention.

The actual hardening processes of the hardening mechanisms used here take place especially during the range between 850 ° C. and a residence time which should be close to the Ar ^ limit. Here, in a further process step according to the invention, there is an accelerated cooling without rolling to about 650/550 ° C., as a result of which the _< T- ~ conversion is reduced with a simultaneous recrystallization delay.

When using stages 1 and 2 of the method according to the invention, strength classes of tempered wire in accordance with Euro standard 138 are achieved, and indeed without the costly tempering and tempering, a further advantage of the method according to the invention based on the chemical composition according to the invention. The tensile strength of the tension rods and wire rod is increased by at least 20% compared to conventional grades, from which cold-drawn wires and strands can be produced from them with correspondingly increased strength properties.

According to the invention, a third stage of the treatment can also be provided, in which rolling is again carried out in a controlled manner from about 650/550 ° C. 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, for example in still air, are considered below. As a result, an increased precipitation hardening process results in an increase in strength of over (ind.) 40% compared to conventional prestressing steels. The accompanying diagram serves to illustrate this process sequence.

After the thermomechanical treatment and the hardening mechanisms that run with it, the steel can additionally be strain hardened, provided that higher strength classes are aimed for or are necessary.

During the course of the thermomechanical treatment according to the present invention, the mechanisms of increasing strength due to the chemical composition and the targeted metering of the microalloying elements work together additively. 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. This means that the thermo-mechanical treatment, along with the chemical composition for fine-grain melting and hardening, is the most important step in realizing the desired The goal, namely for the production of high-strength, corrosion-resistant, brittle fracture-proof and weldable prestressing steel, is. 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 dosing 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.

Precipitation hardening in particular due to the accelerated cooling and a low final rolling temperature with a high degree of deformation and a high rate of deformation followed by a delay after the final deformation and delayed. Cooling achieves the highest effect of the increase in strength, this phase of thermomechanical treatment is also of the greatest importance, because through this phase 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 the fine-grain hardening, a fine-grain depletion being required for its optimal realization, which simultaneously increases the 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.

For this purpose, an austenite grain that is as fine as possible should already be aimed for, since this also determines the size of the ferrite grain. For this purpose it is necessary. that the microalloying elements provided in the directional analysis, in particular aluminum, nitrogen, niobium and vanadium, are incorporated into the austenite structure to inhibit grain growth and to form strength-increasing obstacles to dislocations due to fine precipitations. A particle size of 100 to 200 X is most effective for this, the

2 6

Particle quantity per mm approx. 20. 10 should be.

The fine grain melting should include the following stages according to the invention:

1. a steel pretreatment, with extensive desulfurization being aimed for. this happens

by calcium treatment CAB, for example by the TN method.

2. A steel aftertreatment, in particular an inert gas purging, a vacuum treatment, a deoxidizing, full sedation, and, if possible and relevant, an inclusion modification and / or a pan treatment with metallic calcium or calcium halide slags.

Continuous casting is a good choice for casting. Continuous casting is the most economical and at the same time the best quality way of casting and solidifying the molten steel into the primary material used for the manufacture of prestressing steel: billets. To ensure a high level of quality required for prestressing steels, special measures must be taken to prevent such defects, such as, to avoid core defects such as mitigating and solidification bridges and surface defects. B. Reoxidation protection, concealed casting, electromagnetic stirring. The low carbon content of 0.1% provided in the directional analysis largely prevents the occurrence of the abovementioned defects and at the same time favors the economy of continuous casting for the production of prestressing steel grades, since the costly measures are more extensive than for conventional ones , high-carbon prestressing steel grades are required, while at the same time ensuring a high degree of purity, homogeneity and quality.

In order to take account of the first sub-task mentioned, namely the increase in the strength of prestressing steels, the most important factors influencing the strength must be taken into account. These include in particular obstacles to dislocation movements. The structure has a particularly high influence on the strength properties of prestressing steels, since the presence or formation of obstacles to the displacement movement must be present in order to achieve any kind of strength increases. These obstacles can be divided into according to their dimensions

- zero dimensional. These are punctiform obstacles like foreign atoms in the mixed crystal. Increased strength through mixed crystal hardening.

- one-dimensional. These are line-shaped obstacles as dislocations. Strain hardening by cold working.

- two diagonal. These are surface obstacles as grain boundaries. Solidification through grain refinement.

- three-dimensional. These are spatial obstacles as excretions. Solidification by particle hardening or dispersion hardening.

The mixed crystal hardening acts through the type of chemical composition, the influence of the foreign atoms in substitution mixed crystals and the interstitially dissolved foreign atoms being of particular importance. There are There are numerous diagrams and tables from which the individual alloy elements and their effect on the increase in yield strength can be read. The influence of the different alloy elements can be explained by the distortion that these elements cause in the lattice. The greater the distortion, the higher the increase in strength.

Of all four types of hardening, 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 in a simultaneous increase in toughness. Furthermore, the ^{'two-dimensional} obstacles are just for migratory Versetzun¬ gen so strong obstacles that they can not be overcome über¬ of these. 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 strength. However, the average grain size influences the lower yield strength.

In the case of particle hardening by precipitation, it must be emphasized * that the highest degree of solidification is given when the particle size and the particle spacing are just large enough that no cutting occurs. The precipitation processes for particle hardening are strongly influenced by the degree of oversaturation, the deformation, the conversion and ultimately the recrystallization, which must be given special attention below in the thermomechanical treatment to increase the strength. When developing high-strength prestressing steels, the precipitation of carbides, nitrides or carbo- nitrides can be taken into account by particle hardening. It must also be noted that the precipitation of special carbides or carbonitrides of the micro-alloying elements niobium and vanadium has a specifically higher hardening effect than, for example, copper deposits. According to the invention, these individual hardening mechanisms can be combined with one another and with targeted cold hardening, their effect being additive, but their respective proportions can change considerably depending on the specified conditions. However, according to the invention it was found that the basic mechanisms of the individual hardenings only become optimal through a further, the most important treatment step, namely through the so-called thermomechanical treatment.

The term thermomechanical treatment is intended to subsume a number of specially controlled shaping processes in which the influencing factors

- deformation temperature,

- degree of deformation,

- deformation speed,

- time of deformation,

- final deformation temperature,

- cooling rate,

- conversion Ϊ ^** - oi. . - Dwell time after deformation as well

- subsequent cooling

each plays an important role in terms of the optimal improvement of the steel properties. Virtually all characteristic values of the mechanical properties can be influenced by a thermomechanical treatment, but in particular strength and toughness properties as well as the transition temperature and thus the insensitivity to brittle fracture.

The thermomechanical treatment in the context 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 setting a low final rolling temperature, rapid cooling before the last rolling pass and a high degree of final deformation, so that the Recrystallization leads to the finest possible austenite grain before the ferrite-pearlite transformation. In the controlled rolling of the steel according to the invention after directional analysis in the sequence according to the invention, in this microalloyed steel, the rolling process additionally causes hardening of mixed crystals as well as fine grains and particles by the precipitation of carbides, nitrides or carbonitrides. In addition, the temperature control is alloying and by rolling so ge controls that Diev - <A- conversion shortly before and / or after the lowest possible final rolling temperature, which comes to lie shortly before A _3, takes place. In any case, the formation of martensite should be excluded.

It is important for low-pearlite, micro-alloyed steels that the carbides and nitrides of the microalloying elements niobium and vanadium have face-centered cubic lattices and are isomorphic and therefore completely miscible. The highest strength-increasing effect through the aforementioned hardening mechanisms, however, is effective in the body-centered grid. The shape and size of the carbonitride precipitate must also be taken into account. The particle size and quantity, or the particle spacing, as well as the shape and arrangement of the precipitates and their strength are themselves decisive for influencing the mechanical properties.

These sizes are influenced by the chemical composition and above all by the temperature-time conditions under which the precipitates form. Depending on the temperature, the carbonitrides can precipitate in austenite, during the ^{~ *} - <transformation or in ferrite. Elimination in the ferrite is most effective for increasing the strength. The kinetics, the extent and the temperature of the precipitations depend not only on the thermodynamic conditions, but also on the diffusibility of the alloy atoms, the degree of hypothermia and the germ conditions of the precipitates.

The practical measures that have to be used to optimize the proportion of fine grain hardening in connection with the thermomechanical treatment required for this development are for the alloy according to the invention:

- Low, fine grain preserving furnace temperature, in particular to prevent or limit how- the dissolution of carbide, nitride and / or carbonitride precipitates,

- High degree of deformation with few sequences,

- low forming temperature,

- Lowering of the Y ^~ - _< {* -, conversion temperature by accelerated cooling and / or by alloy and / or

- recrystallization delay.

In the thermomechanically treated state, the optimally cheapest (finest) grain size with all the advantages with regard to the increase in strength and at the same time the most favorable effect on toughness properties and transition temperature is established.

The proportions of fine grain hardening and hardening and thus a significant proportion of the possible increase in strength in the present invention are very much determined by the production conditions, i.e. the thermomechanical treatment. This requires for this purpose

- a high forming speed and degree,

- A quick and controlled cooling before and / or after the last roll pass and

- a subsequent delayed cooling,

- tailored to the production of thermomechanically acted, cold-stretched tension rods or wire rod for the production of cold-drawn tension wires and Lit¬ zen from it.

The final rolling temperature and the degree of deformation, in particular in the last pass, are decisive for the mechanical properties that can be achieved. With a decrease in the final rolling temperature, the pearlite content decreases, which means that low-carbon, micro-alloyed structure structures in the controlled final rolled state have only a low, often no pearlite content in the structure. The mechanical properties are thereby additionally favorably influenced.

With a higher stitch decrease and a smaller number of stitches, smaller austenite grain sizes are achieved, which result in more favorable mechanical properties through a correspondingly smaller ferrite grain. Increasing stitch decreases of 10 to 45% have a particularly favorable effect on a finer ferrite grain size and then on a noticeable improvement in the transition temperature or the sensitivity to brittle fracture. Stitch acceptance and final rolling temperature as well as possible holding times must be matched to the desired properties and dimensions of the end products tension rods (thermo-mechanically treated, cold-drawn) and wire rod (for the production of cold-drawn wires) in order to achieve the desired metallurgical effect on the one hand and a rolling technology economical process on the other to ensure. The decisive factors influencing the mechanical properties that can be achieved are rapid rolling and cooling after finish rolling. A low temperature affects the ferrite grain size as a result of the accelerated cooling Y ^{~ "} -« conversion shifted lower temperatures, on the other hand, the excretion processes taking place during the subsequent slow cooling are considerably supported.

For the structure formation that results in prestressing steel, recrystallization is the ^ - <. -Conversion and excretion of microalloying elements crucial. These processes can run side by side in a very short time of a few minutes and also influence one another. For these reasons, it is necessary for the development of high-strength, corrosion-resistant and brittle fracture-proof prestressing steels, based on mechanical properties and dimensional ranges, to carry out a precise recording of the processes taking place and their assignment to the structural structures which arise and the properties caused by them optimize.

The last stage of the process following the thermomechanical treatment is work hardening, which in particular consists of stretching or drawing. 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, once again results in a considerable increase in strength compared to today's prestressing steel grades by means of the degree of deformation to be used.

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 of possible microalloying elements the fine grain hardening and hardening by the thermomechanical treatment, ie the increase in strength, is followed by vanadium. The same applies to the improvement of the transition temperature.

By microalloying with niobium and vanadium and at the same time low in pearlite, the strengthening proportion of the manganese and silicon content also increases with increasing contents.

An increase in the nitrogen content, with the simultaneous presence of vanadium, brings about an additional increase in the yield strength. This also increases the tensile strength, so that an increase in the yield strength ratio from around 70% to 90%, which is particularly important for prestressing steels, is brought about.

With steel, niobium alloy results 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 through the addition of niobium is therefore a major reason why niobium must be used here, since niobium also causes the greatest reduction in the transition temperature.

With regard to the improvement of the transition temperature or the insensitivity to brittle fracture, it must be stated that by alloying with niobium and vanadium, a connection between the increase in the yield point and the improvement in the transition temperature is independent of the microscope. Control elements exist. With the same yield point but different niobium or vanadium contents, almost the same insensitivity to brittle fracture or transition temperature is achieved.

Manganese and nickel as well as silicon with contents below about 0.5% also shift the transition temperature to lower temperatures.

In addition to solidification, grain refinement also leads to a significant improvement in toughness, which is manifested in a sharp reduction in the transition temperature. In addition, the desired influence is reinforced by a decreasing pearlite content. Low-pearlite steels are therefore generally particularly insensitive to brittle fracture in the case of fine ferrite grains.

Particular attention must also be paid to their chemical composition with regard to the cold-forming properties of prestressing steels. The sulfur content plays the decisive role in the anisotropy of toughness, the most important factor influencing the cold formability. A desired lower sulfur content, ie a reduced number of sulfide inclusions, considerably improves the toughness with regard to constriction of the fracture, a property which is particularly important for prestressing steels. In addition, the reduction in the sulfide length is particularly effective for a more favorable constriction of the fracture. Strong desulphurization can be achieved by adding calcium, which is common in ladle metallurgy, with special attention being paid to the high vapor pressure of calcium, which is 1.86 bar at a melt temperature of 1600 ° C, and its high affinity for oxygen. must be given, ie measures must be taken to prevent the evaporation of the calcium. Even with sulfur contents of 0.008%, no more manganese sulfides are found in aluminum-calmed steels, but spherical inclusions made of calcium and aluminum oxides, which contain small amounts of sulfur in solution on their surface. Due to the favorable conditions of the calcium aluminates with regard to separation from the melt, an additional improvement in the degree of oxidic purity is achieved. The mechanical properties that can be achieved with calcium treatment have a significantly reduced spatial anisotropy of the toughness properties. The constriction of the fracture, which is so important for ensuring the quality values for prestressing steels, is improved considerably by the calcium treatment and with a decreasing sulfur content. Desulphurization should be carried out to below 0.020 mass% if possible.

With regard to the most suitable combined use of microalloying elements, molybdenum-niobium-alloyed structural structures result in 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.

In addition to the consequences of the thermal treatment, the speed of rolling and cooling and the cooling in bed are also effective for the production of prestressing steels. Both strength and toughness improvements are found down to 750 ° C. The effective The mechanisms which are responsible for the increase in strength are increased considerably by alloying with molybdenum and by regulating the rolling speed with the purpose of reducing the f - <- conversion as far as possible in the range between 650 and 550 ° C, just ^'a region in which the strength-enhancing mechanisms, in particular by precipitation hardening are most effective.

The most effective means of achieving optimal mechanical properties is, however, the production of extensive fine grain. The refinement of the grain size causes an increase in the yield strength with a simultaneous improvement in the transition temperature. In practice, the aim is to achieve the finest possible austenite grain, since this also determines the size of the ferrite grain. A general empirical value is that a reduction in the size of the austenite grain has a factor of around 0.3 on the reduction in the size of the ferrite grain. The essential process in the growth of the austenite grain is not the dissolution of the precipitates, 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. In addition to aluminum, which produces this effect via aluminum nitride, it is primarily the microalloying elements niobium, vanadium and titanium in particle sizes from 100 to 200 Å which are comparable in their carbides, nitrides or carbonitrides Come into effect. Show the most favorable conditions for preventing the sharp increase in grain growth when reheating in the pusher furnace for rolling 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.

A further measure to prevent or limit the re-dissolution of such precipitates when heating before rolling is the lowest possible furnace temperature. Furthermore, the austenite grain can also be refined by means of higher degrees of deformation. The grain refinement effect is most pronounced at low final deformation temperatures.

If the transformation y- ¹ - - is shifted towards lower temperatures by an accelerated cooling, 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 .

In addition to grain refinement, there is the possibility of delaying the recrystallization of the austenite. Portions of non-recrystallized austenite are then deformed during the final rolling temperature, which results in elongated grains and thus greatly enlarged austenite grain surfaces. The transformation of this structure in the ferrite-pearlite stage results in a strong grain refinement due to the increased seed density and the inhibited growth of the grains formed from these seeds.

The delay in austenite recrystallization can also be controlled by controlling the cooling rate Alloying of small amounts of molybdenum to the microalloyed, low-pearlite structure is favored, whereby the i-pC conversion is shifted to lower temperatures. This possibility is used in the thermomechanical treatment, whereby an even more fine-grained structure is achieved with an additional improvement in the transition temperature.

The fact that the favorable transition temperature in niobium and vanadium or niobium plus vanadium-alloyed structural structures remains unchanged or can even be improved can be explained by a larger proportion of the grain refinement. The grain refinement thus also causes, in addition to solidification, that in the present invention desired significant improvement in toughness, which at the same time manifests itself in a sharp reduction in the transition temperature. In addition, this desired influence in this development is reinforced by a decreasing percentage of pearlite. Low-pearlite structures are therefore generally particularly insensitive to brittle fracture in the case of fine ferrite grains.

In the connection between the microalloy components and the fine grain hardening it has to be taken into account that incoherent niobium vanadium and titanium carbonitrides have different effects on the ferrite grain size in their effective particle size and quantity. In the thermomechanically treated state, vanadium only causes a weak 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> ¥ ^* - o conversion. It should also be noted that the optimal Grain refinement due to niobium contents between 0.04 and 0.10% is uniformly effective, but those of titanium and vanadium are also increasingly effective with increasing contents.

The carbon and nitrogen content of the steel influences the ferrite grain size considerably less in steels with niobium than in those with vanadium. With decreasing carbon contents, the influence of the nucleation by excreted particles on the grain size decreases in favor of a very pronounced and, in the present case, desirable inhibition of recrystallization by 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, niobium or titanium cause a further fine grain effect by delaying the austenite conversion desired here. Rising manganese contents also lower the transition temperature, ensure optimum particle separation and thus the optimal effect of particle hardening.

In addition to the delayed austenite transformation, there is usually a delay in recrystallization, i.e. the recrystallization takes place later at lower temperatures, which is a requirement

- lowering the ζ - cC conversion,

- recrystallization delay and thus

- the setting of the lowest possible final rolling temperature accommodates and at the same time enables the optimal separation, for example of copper, interacting to achieve a maximum possible increase in strength. Extensive microstructure refinement occurs as a result of increased germ density and growth hindrance of the newly formed ferrite grains.

With regard to precipitation hardening, it must be taken into account in connection with the thermomechanical treatment that the hardening maxima occur in the temperature range between 550 and 650 ° C. This can be explained by the effect of the chemically non-detectable coherent precipitates (clusters) of niobium, carbon and nitrogen atoms (including titanium), which precede the incoherent precipitate. After the curing maximum has been reached, the drop in the yield point is of importance. This drop is caused by increasing temperatures or exceeding the holding time and is due to the reduction of the coherent stresses when the coherent particles change into incoherent and the subsequent waxing of the particle diameter and quantity.

According to the invention, the starting material (crude steel) for carrying out the method according to the invention is to use a steel which has the following alloying elements in its directional analysis:

Carbon 0.05 to 0.20 mass%

Manganese 1.20 to 1.70 mass%

Silicon 0.30 to 0.50 mass%

Niobium 0.04 to 0.06 mass%

Vanadium 0.035 to 0.05 mass%

Molybdenum 0.30 to 0.50 mass% Copper 0.30 to 2.00 mass% aluminum 0.04 to 0.06 mass%

Nitrogen 0.015 to 0.02 mass% phosphorus ^' - 0.030 mass%

Sulfur ≤. 0.020 mass%

About the individual elements:

Via the cementite (pearlite), the carbon content causes a substantial solidification and plays an important role in this context. However, since the carbon content via the pearlite component exerts the most significant negative influence on the brittle fracture safety (transition temperature) also given in this development, as well as on the weldability, and this with increasing pearlite content, this is To limit the carbon content to proportions which both permit an increase in strength and an improvement in the corrosion resistance, but also enable the safety against brittle fracture to be improved down to around -40 ° C. and also make it weldable. With regard to the desired fine grain formation to be aimed at, it should also be taken into account that the carbon content has a considerable influence on this. With decreasing carbon content, 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 inhibition of recrystallization due to dissolved niobium. In the thermomechanically treated state, low-pearlite microstructures have smaller ferrite grain sizes than microstructures with a higher carbon content.

Manganese has a particularly fine grain refinement and at the same time through solidification of solid solution and increased hardening, so that the manganese content should preferably be placed at the upper limit, because the increase in strength due to manganese depends very much on the pearlite content and, due to a suitably low pearlite content, also a favorable transition Temperature and thus also guaranteed brittle fracture safety. Increasing manganese levels provide a significant contribution to the delay of the ^'desired austenite transformation here and thus cause an optimum 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, ie on the achievable increase in strength, followed by titanium and 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 thus not only a higher yield strength than structure structures alloyed with titanium or vanadium, but above all also a very favorable, low transition temperature. Niobium reduces the ferrite grain size to a particularly large extent. The high ratio of fine grain hardening to hardening in the structure with the addition of niobium is therefore an essential reason for the preference for niobium. Niobium causes the additional hardening effect of increasing manganese contents even when pearlite is low.

Vanadium, like niobium, 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. Vanadium, like niobium, thus helps to control the austenite grain size by incorporating fine precipitates "in the austenite structure, which inhibits grain growth. Vanadium also contributes to solid solution strengthening like niobium, but both are insoluble in ferrite. Their excretion in ferrite The carbides and nitrides of vanadium and niobium have face-centered cubic lattice, are isomorphic and therefore completely miscible and, unlike titanium, do not contribute to the formation of sulfide. If the nitrogen content is increased, vanadium has the greatest influence on the formation of a fine ferrite grain size and has an additional increase in the yield point, just like niobium, as does dissolved vanadium, by delaying the austenite conversion, this fine grain effect and hardening.

The delay in the austenite recrystallization is very much favored by alloying small amounts of molybdenum into the microalloyed, low-pearlite structure, whereby the

- Conversion is shifted to lower temperatures. In thermomechanical treatment, this possibility is enhanced by an even deeper one Final rolling temperature used, whereby an even more fine-grained structure is achieved while improving the transition temperature. In addition, by alloying molybdenum and the resulting possibility of the i - - r conversion shift to lower temperatures, it is also additionally possible to make full use of the considerable strengthening properties of copper. In the case of microalloyed microstructures of the type described here and at the same time low pearlite fractions and a high copper content, both hardening mechanisms act both through the precipitation of mixed crystals and through the formation of carbonitrides, particularly at temperatures between 650 and 550 ° C.

With additionally high manganese and molybdenum contents, such as here for high-strength prestressing steels, in addition to the particle hardening, an additional increase in strength can be achieved in the case of copper-alloyed structural structures through a high dislocation density and fine grain hardening.

Copper is used for the purpose envisaged here because of its two advantages. Firstly because of its strong hardening effect through hardening. Secondly, because of its strong corrosion-inhibiting effect. The corrosion-inhibiting effect of copper can be used particularly well in high-strength structural structures produced with thermomechanical treatment, because at the low final rolling temperatures, which at the same time also lead to the highest strength increases, the element copper simultaneously with the The precipitation-hardening elements used here between 650 and 550 ° C., in addition to its corrosion-inhibiting effect, also act as a precipitation-hardening element. By rapid cooling from the Jf ^ area At approximately 840 ° C., about 2% copper can be dissolved in low-pearlite structures and the thermomechanical treatment already provided here. A copper-rich, face-centered mixed crystal in the form of incoherent, spherical particles then separates out, which leads, from a certain particle size, to a considerable precipitation hardness effect due to the bypass mechanism. In the presence of niobium, micro-alloyed structures and at the same time a low proportion of pearlite and a high copper content both hardening mechanisms take effect through the precipitation of mixed crystals and carbonitrides. In the case of high copper contents, however, a nickel content of up to 1% must be added to the copper-alloyed structure in order to prevent the solder fragility caused by copper. With additional high manganese and molybdenum contents, as also provided here, in copper-alloyed structure structures, in addition to particle hardening, an additional increase in strength can be achieved through a high density of dislocation 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 so that, on the one hand, the corrosion-inhibiting effect and the strengthening mechanisms can be optimally used, on the other hand, the solder brittleness, which would not be acceptable for prestressing steels, should not be effective and, if possible, a nickel addition to avoid this prevention.

Due to the aluminum content, the sudden grain growth when the primary material is heated to approximately 1150 ° C. is raised, the holding time also being important. Next Aluminum, which produces this effect via aluminum nitride, is primarily the microalloying elements niobium and vanadium, which have a comparable effect via their carbides, nitrides or carbonitrides. The lowest possible pusher furnace temperature is essential for preventing or restricting the re-dissolution of such precipitates when heating before rolling. The most favorable conditions for preventing the sharp increase in grain growth when reheating for rolling show higher aluminum contents. Aluminum also contributes to solid solution strengthening.

In addition to aluminum, 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 an important contribution to increasing the strength by increasing the nitride content. Particularly when 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.

In the present case, the phosphorus content must remain limited, although a higher content would increase the yield strength, but at the same time the steel becomes very brittle. Combined oxygen blowing / inert gas flushing makes it possible to lower the phosphorus content and largely prevent its embrittling effect. A corresponding reduction in the phosphorus content is also possible with ladle metallurgy.

According to the invention, the lowest possible phosphorus of particular importance and therefore to be striven 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 considerably with regard to constriction of fracture, a property that is particularly important for prestressing steels. In addition, 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 additions that are usual in ladle metallurgy.

Regarding the titanium not involved in the invention, it should be noted that, 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 ^ C-S-. For both reasons, 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 cooperation with the other microalloying elements would be eliminated by an increased nitrogen content.

In the production of high-strength, corrosion-resistant and brittle-fracture prestressing steels according to the invention, all those difficulties which have to be taken into account in the production of conventional, high-carbon prestressing steel grades are eliminated. Above all, there are no major concerns about a Position in the continuous casting process, which results primarily from the center segregations and surface defects that impair drawing ability. The economic viability of continuous casting compared to ingot casting then comes into play, both in terms of effort and in terms of quality. The accumulation of carbon in the middle of the strand that was previously possible is largely eliminated, which leads to eutectoid deposits of cementite networks and thus to a considerable deterioration not only of the structure and the properties thereof, but also of the brittle fracture safety.

The measures which have to be taken during the entire production process due to the high affinity for carbon for oxygen are also largely eliminated, both during melting (for example during the build-up or remelting process), freshening and the subsequent steel aftertreatment, but in particular also an elaborate reoxidation protection. The achievement of a high microscopic degree of purity, largely avoiding oxidic and sulfidic inclusions, is favored. In the case of continuous casting, the high outlay which is necessary for the electromagnetic stirring in the production of high-carbon wire grades is largely eliminated, as a result of which the very disadvantageous central increases, solidification bridges, directional solidification structures, internal and surface defects can largely be prevented.

The prestressing steels produced according to the present invention have

- significantly higher strength values, - significantly reduced internal stresses,

- significantly increased brittle fracture collateral,

- significantly increased wear resistance,

- Significantly improved applications due to their suitability for welding and

- Significantly improved corrosion resistance.

With regard to the latter two advantages, it should also be pointed out that, with a view to improving the corrosion resistance according to the present invention, elements which can be used economically are considered, which act in a manner similar to that of stainless steels and chromium. In addition, such corrosion-inhibiting elements can be used particularly well in high-strength steels produced with thermomechanical treatment because, in addition to the corrosion-inhibiting effect, they are used at the low final rolling temperatures, which at the same time also bring about the highest strength increases contribute to increased strength through precipitation hardening. If, in addition to increasing the strength of high-strength prestressing steels, it is also possible to achieve welding suitability, this results in considerable and significant possibilities for constructively simplifying and improving the clamping systems currently in use. As is known, for example in bridge construction, the coupling links are the most sensitive weak points for the occurrence of damage caused by the penetration of corrosion-promoting 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 at the same time. When using the high-strength, corrosion-resistant and brittle fracture-proof prestressing steels according to the invention, it becomes possible to produce longer prestressing strands by means of which the number of coupling members and thus the weak points is reduced. In addition, if the prestressing systems are also structurally simplified and improved due to the suitability of these prestressing steels for welding, this additionally results in a significant reduction in the susceptibility to damage.

Other advantages are

smaller and therefore easier to control diameters of tensioning wires _/ rods or strands,

- Due to the higher strength properties, the construction of smaller concrete thicknesses is possible, v / or itself

- A saving of overall construction weight, on the one hand, and

- On the other hand, significantly increased possibilities in the construction design result, that is

- Designs of constructions that cannot be realized with conventional prestressing steel with lower strength for technical or economic reasons, and

- a reduction in the total load of moving structures (bridge construction, element construction, for example) and - Reduction of transport costs for moving constructions and prestressing steel.

Despite the use of microalloying elements and improved steel aftertreatment to increase strength, increase corrosion resistance and brittle fracture safety, the current price level of prestressing steels can be kept or even improved thanks to the considerable advantages in their manufacture and use. The additional design options, which result from the welding suitability of clamping systems, greatly increase cost-effectiveness. Taken as a whole, however, the economic advantages would outweigh the disadvantage of an increase in prices.

Claims

Claims

1. A process for the manufacture of high-strength, weldable, corrosion-resistant and brittle fracture-proof prestressing steels, characterized in that a steel consists of

0.05 to 0.20 mass% carbon

1.20 to 1.70% by mass of manganese

0.30 to 0.50 mass% silicon

0.04 to 0.06% by mass of niobium

0.035 to 0.05 mass% vanadium

0.30 to 0.50 mass% molybdenum

0.30 to 2.00 mass% copper

0.04 to 0.06 mass% aluminum

0.015 to 0.02 mass% nitrogen

<£ 0.030 mass% phosphorus 0.020 mass% sulfur

is subjected to a thermomechanical treatment, which takes place after solidification from the melt and reheating from a second heat, the steel prior to the thermo-mechanical treatment at the lowest possible reheating temperature (= second heat below 1150 ° C.) is held and subsequently a controlled rolling of the steel with a small number of stitches is carried out at a high degree of deformation (10-45%) and a high deformation speed up to a low forming temperature close to above 850 ° C. - -

2. The method according to claim 1, characterized in that the thermomechanical treatment comprises a second stage u, in which an accelerated cooling takes place from about 850 ° C without rolling to about 650/550 ° C, thereby reducing the ft "-" c conversion and a recrystallization delay is effected.

3. The method according to claim 1 or 2, characterized in that the thermomechanical treatment comprises a third stage in which from about 650/550 ° C again controlled with one or a few stitches, that is, with a high degree of deformation at high speed, is rolled to a low final rolling temperature near above the Ar limit and then after a dwell time there is a delayed cooling.

4. The method according to at least one of claims 1 to

3, characterized in that the steel undergoes work hardening after the thermomechanical treatment.

5. The method according to at least one of claims 1 to

4, characterized in that the steel is largely desulfurized before and / or after the refreshing.

6. The method according to claim 5, characterized in that the steel melt is subjected to a calcium treatment before and / or after the refining.

7. The method according to claim 5 or 6, characterized gekenn¬ characterized in that a steel aftertreatment, for example inert gas purging, vacuum treatment, deoxidizing, inclusion modification or ladle treatment with metallic calcium or calcium halide slags takes place.

8. The method according to any one of claims 1 to 7, characterized ge indicates that as solidification mechanisms during the thermomechanical treatment, both a Mischkri¬ stal, fine grain and particle or precipitation hardening with largely additive effect is used.

9. The method according to any one of claims 1 to 8, characterized ge indicates that the thermomechanical treatment is carried out by the controlled rolling of micro-alloyed, fine-grained steels and excludes the formation of martensite.

10. The method according to claim 9, characterized in that the recrystallization of these fine-grained, micro-alloyed steels leads to the finest possible austenite grain before the ferrite-pearlite transformation.

11. The method according to claim 10, characterized in that - in the case of the microalloyed steels, the rolling process is additionally supplemented by precipitation of carbides, nitrides and / or carbonitrides, as a result of which both a mixed crystal, as well as a fine grain and particularly enhanced particle hardening becomes.

12.The method according to any one of claims 2 to 11, characterized ge indicates that the temperature control is controlled so that a V - .- conversion takes place shortly before and / or during the lowest possible final rolling temperature, which occurs shortly before A _r3 .

13.The method according to any one of claims 4 to 12, characterized ge indicates that after the thermomechanical treatment as work hardening, stretching (for tension rods) or drawing (for cold drawn wires) takes place.

14.The method according to any one of claims 1 to 13, characterized ge indicates that by alloying the micro alloys tion elements niobium and / or vanadium and / or molybdenum, an optimally possible particle hardening in the form of carbides, nitrides and / or carbonitrides is effected by precipitation during the thermomechanical treatment in addition to fine grain and mixed crystal hardening.

15. The method according to any one of claims 1 to 14, characterized in that the austenite grain is refined by incorporating fine precipitates such as aluminum nitrides and carbides, nitrides, and / or carbonitrides, in particular the noble and vanadium microalloying elements in a particle quantity of 20 x 10 per mm and in a particle diameter of 100 to 200 A, the highest possible degrees of deformation and speeds being used in addition to the lowest possible final rolling temperature.

16. The method according to any one of claims 1 to 15, characterized in that portions of non-recrystallized austenite are deformed during the low final rolling temperatures by a recrystallization delay, in advance of elongated grains and thus greatly enlarged ^' austenite grain surfaces' give, when they are converted in the ferrite-pearlite stage, a high grain density and the inhibited growth of the grains formed from these grains result in a large grain refinement and an optimal increase in strength through fine grain as well as particle hardening. - 42 -

17. The method according to any one of claims 14 to 16, characterized in that the austenite recrystallization in addition to the control of the cooling rate by alloying with molybdenum and thus the ^r - </ L conversion is shifted to lower temperatures.

18. The method according to any one of claims 14 to 1.7, characterized in that a desired grain refinement is achieved by alloying an increased manganese content within the scope of the specified directional analysis and then an optimum increase in strength is ensured simultaneously by solid solution strengthening and increased hardening.

19. The method according to claim 18, characterized in that, due to increasing manganese contents, an optimal delay of the desired austenite conversion and thereby the optimal fine grain formation is guaranteed.

20. The method according to claim 18 or 19, characterized labeled in ^¬ characterized in that by increasing manganese contents gleich¬ time the desired Rekristallisationsverzögerung is ensured by shifting the conversion to T ^f tiefereren temperatures and

Setting the lowest possible final rolling temperature and simultaneous use of the thermomechanical treatment.

21. The method according to any one of claims 18 to 20, characterized in that the optimal separation of particles and thus the optimal effect of particle hardening for the maximum possible strength increase is ensured by an increased manganese content.

22. The method according to any one of claims 18 to 21, characterized in that the strengthening and thus strength-increasing proportion of the manganese is increased by increased manganese content and simultaneous presence of niobium and vanadium as well as low pearlite content.

23. The method according to any one of claims 1 to 22, characterized in that the pearlite content is lowered.

24. The method according to any one of claims 1 to 23, characterized in that, or the like to prevent rapid grain growth during preheating in the batch furnace. increased aluminum and nitrogen contents are used in the course of the given directional analysis, with this method for this purpose a particle quantity of

6 6 oo 2 200 xx 1100 pprroo mmmm and one 'particle diameter from 100 to

0

200 A should be aimed for.

25. The method according to any one of claims 1 to 24, characterized in that the beginning of the grain growth to be prevented is shifted to elevated temperatures by an increasing niobium content within the specified directional analysis.