EP2488671A1

EP2488671A1 - Heat treatment of martensitic stainless steel after remelting under a layer of slag

Info

Publication number: EP2488671A1
Application number: EP10781971A
Authority: EP
Inventors: Laurent Ferrer; Patrick Philipson
Original assignee: SNECMA SAS
Current assignee: Safran Aircraft Engines SAS
Priority date: 2009-10-12
Filing date: 2010-10-11
Publication date: 2012-08-22
Anticipated expiration: 2030-10-11
Also published as: RU2567409C2; EP2488671B1; RU2012119551A; FR2951198B1; FR2951198A1; CA2776851A1; US20120199252A1; US8808474B2; CA2776851C; BR112012008524B1; CN102575311B; JP2013507532A; WO2011045515A1; CN102575311A; JP5778158B2; BR112012008524A2

Abstract

The invention relates to a method for producing a martensitic stainless steel that includes a step in which an ingot of the steel is remelted under a layer of slag, a subsequent step in which the ingot is cooled and at least one austenitic thermal cycle consisting in heating the ingot above the austenitic temperature thereof, followed by a cooling step. During each of the cooling steps, if the cooling step is not followed by an austenitic thermal cycle, the ingot is maintained at a holding temperature within the ferrite‑pearlite transformation nose region for a holding time greater than that required to transform the austenite as completely as possible into a ferrite‑pearlite structure in the ingot at the holding temperature, whereby the ingot is maintained at the holding temperature once the temperature of the coolest point in the ingot has reached said holding temperature. Moreover, during each of the cooling steps, if the cooling step is followed by an austenitic thermal cycle, before the minimum temperature of the ingot drops below the martensitic transformation start temperature Ms, the ingot is either: maintained at a temperature above the heating‑induced austenitic transformation finish temperature Ac3 for the entire duration between these two austenitic thermal cycles, or maintained at the holding temperature within the ferrite‑pearlite transformation nose region, as above.

Description

THERMAL TREATMENTS OF STAINLESS STEEL MARTENSIAL STEELS

ORE REFUSAL UNDER DAIRY

The present invention relates to a method for manufacturing a stainless steel martensitic comprising a slag remelting step of an ingot of this steel then a cooling step of the ingot and then at least one austenitic thermal cycle consisting of heating this ingot above its austenitic temperature.

In the present invention, the percentages of composition are percentages by weight unless otherwise specified.

A stainless martensitic steel is a steel whose chromium content is greater than 10.5%, and whose structure is essentially martensitic.

It is important that the fatigue strength of such a steel is the highest possible, so that the life of parts made from this steel is maximum.

For this, it seeks to increase the inclusion cleanliness of the steel, that is to say, to reduce the amount of undesirable inclusions (certain alloy phases, oxides, carbides, intermetallic compounds) present in the steel. Indeed, these inclusions act as crack initiation sites which lead, under cyclic stress, to premature failure of the steel. Experimentally, a considerable dispersion of the results of fatigue tests is observed on test specimens of this steel, that is to say that for each level of fatigue strain imposed strain, the service life (corresponding to the number of cycles leading to the rupture of a fatigue test piece in this steel) varies over a wide range. Inclusions are responsible for the minimum statistical values of steel fatigue life (low values in the range).

To reduce this dispersion of the fatigue strength, that is to say up those low values, and also to increase its average value in resistance to fatigue, it is necessary to increase the inclusion cleanliness of the steel. We know the technique of slag remelting, or ESR (Electro Slag Refusion). In this technique, the steel ingot is placed in a crucible in which a slag (mineral mixture, for example lime, fluoride, magnesia, alumina, spath) has been poured in such a way that that the lower end of the ingot quenched in the slag. Then an electric current is passed into the ingot, which serves as an electrode. This current is high enough to heat and liquefy the slag and to heat the lower end of the steel electrode. The lower end of this electrode being in contact with the slag, melts and passes through the slag in the form of fine droplets, to solidify below the layer of supernatant slag, into a new ingot that grows gradually. The slag acts, inter alia, as a filter which extracts the inclusions from the steel droplets, so that the steel of this new ingot located below the slag layer contains fewer inclusions than the initial ingot (electrode). . This operation is carried out at atmospheric pressure and air.

Although the ESR technique makes it possible to reduce the dispersion of the fatigue strength in the case of stainless martensitic steels by elimination of inclusions, this dispersion in terms of lifetime of the parts still remains too important.

Non-destructive ultrasonic testing, performed by the inventors, showed that these steels practically had no known hydrogen defects (flakes).

The dispersion of the fatigue strength results, specifically the low values of the range of results, is therefore due to another undesirable mechanism of premature initiation of cracks in the steel, which leads to its premature failure in fatigue.

The present invention aims to provide a manufacturing method that allows to raise these low values, and thus reduce the dispersion of the fatigue strength of stainless martensitic steels, and also to increase its average value in resistance to fatigue.

This goal is achieved by virtue of the fact that during each of the cooling steps;

- If the cooling step is not followed by an austenitic thermal cycle, said ingot is maintained at a holding temperature included in the ferrito-pearlitic transformation nose for a holding time longer than the time necessary to transform as completely as possible the austenite in ferritic-pearlitic structure in this ingot at the holding temperature, the ingot being maintained at this holding temperature as soon as the temperature of the coldest point of the ingot has reached the holding temperature,

If the cooling step is followed by an austenitic thermal cycle, the ingot is, before its minimum temperature is lower than the martensitic transformation start temperature Ms, to be maintained, for the entire duration between these two thermal cycles. austenitic, at a temperature above the end of austenitic transformation temperature in Ac3 heating, is maintained at the holding temperature included in the ferrito-pearlitic transformation nose as above.

Thanks to these provisions, it reduces the formation of microscopically sized gas phases (not detectable by industrial non-destructive testing means) and consist of light elements within the steel, and thus avoids the premature crack initiation from these microscopic phases which leads to the premature failure of the steel in fatigue.

Advantageously, the ingot is placed in an oven before the skin skin temperature is lower than the end of ferritic-pearlitic transformation Arl cooling Arl temperature which is higher than the martensitic transformation start temperature Ms.

The invention will be better understood and its advantages will appear better on reading the detailed description which follows, of an embodiment shown by way of non-limiting example. The description refers to the accompanying drawings in which:

FIG. 1 compares fatigue life curves for a steel according to the invention and a steel according to the prior art,

FIG. 2 shows a fatigue stress curve,

FIG. 3 is a diagram illustrating dendrites and interdendritic regions,

FIG. 4 is a photograph taken under an electron microscope of a fracture surface after fatigue, showing the gas phase having initiated this fracture,

FIG. 5 schematically shows cooling curves on a time-temperature diagram for a region richer in alphagenic elements and less rich in gamma-ray elements, FIG. 6 schematically shows cooling curves on a time-temperature diagram for a region less rich in alpha-gene elements and richer in gammagenic elements.

During the ESR process, the steel that has been filtered by the slag cools and gradually solidifies to form an ingot. This solidification takes place during cooling and is carried out by growth of dendrites 10, as illustrated in FIG. 3. In accordance with the phase diagram of the stainless martensitic steels, the dendrites 10, corresponding to the first solidified grains, are by definition richer in elements. alphagenes while the interdendritic regions 20 are richer in gamma-containing elements (application of the known rule of the segments on the phase diagram). An alphagene element is an element that favors a ferritic type structure (structures that are more stable at low temperature: bainite, ferrite-pearlite, martensite). A gamma element is an element that promotes an austenitic structure (stable structure at high temperature). There is therefore segregation between dendrites 10 and interdendritic regions 20.

This local segregation of chemical composition is then preserved throughout the manufacturing process, even during subsequent hot forming operations. This segregation is therefore found on both the solid ingot of solidification and on the subsequently deformed ingot.

The inventors have been able to show that the results depend on the diameter of the ingot coming directly from the ESR crucible or the ingot after hot deformation. This observation can be explained by the fact that the cooling rates decrease with increasing diameter. Figures 5 and 6 illustrate different scenarios that may occur.

FIG. 5 is a temperature (T) -time (t) diagram known for a region richer in alphagenes and less rich in gamma elements, such as dendrites 10. Curves D and F mark the beginning and the end of the transformation of austenite (region A) into a ferrito-pearlitic structure (FP region). This transformation takes place, partially or fully, when the cooling curve that follows the ingot passes respectively in the region between the curves D and F or in addition in the region FP. It does not occur when the cooling curve is entirely in region A.

FIG. 6 is an equivalent diagram for a region richer in gammagenic elements and less rich in alphagenic elements, such as the interdendritic regions 20. It will be noted that with respect to FIG. 5, the curves D and F are shifted to the right, that is, the ingot will have to be cooled more slowly to obtain a ferritic-pearlitic structure.

Each of Figures 5 and 6 shows three cooling curves from austenitic temperature, corresponding to three cooling rates: fast (curve C1), average (curve C2), slow (curve C3).

During cooling, the temperature begins to decrease from an austenitic temperature. In the air, for the diameters concerned in our case, the cooling rates of the surface and the heart of the ingot are very close. The only difference is that the surface temperature is lower than that of the core because the surface was the first to cool relative to the core.

In the case of faster cooling than rapid cooling (curve C1) (FIGS. 5 and 6), the ferrito-pearlitic transformations do not take place.

In the case of rapid cooling according to the curve C1, the transformations are only partial, only in the dendrites (Figure 5).

In the case of an average cooling according to the C2 curve, the transformations are only partial in the interdendritic spaces 20 (FIG. 6) and almost completely in the dendrites 10 (FIG. 5).

In the case of slow cooling according to the C3 curve and even slower cooling, the transformations are almost complete both in the interdendritic spaces 20 and in the dendrites 10.

In the case of rapid cooling (Cl) or medium (C2), there is more or less cohabitation between ferritic regions and austenitic regions.

Indeed, once the material has solidified, the dendrites 10 first turn into ferritic structures during the course of cooling (crossing the curves D and F of Figure 5). While the interdendritic regions 20 either do not change (in the case of rapid cooling according to the curve C1) or change later, in whole or in part (in the case of average cooling according to the curve C2 or slow according to the curve C3), to temperatures inferior (see Figure 6).

The interdendritic regions thus retain a longer austenitic structure.

During this cooling in the solid state, locally, there is a structural heterogeneity with coexistence of austenitic microstructure and ferritic type. Under these conditions, the light elements (H, N, O), which are more soluble in the austenite than in the ferritic structures, therefore tend to concentrate in the interdendritic regions 20. This concentration is increased by the higher content in gammagens in the interdendritic regions 20. At temperatures below 300 ° C, the light elements only diffuse at extremely low speeds and remain trapped in their region. After the interdentitic zones 20 have been converted into a total or partial ferritic structure, the solubility limit of these gaseous phases is reached under certain concentration conditions and these gaseous phases form pockets of gas (or a substance in a physical state allowing great malleability and incompressibility).

During the cooling phase, the more the ingot at the outlet of the ESR (or the subsequently deformed ingot) has a large diameter (or, more generally, the larger the dimension of the ingot is large) or the cooling rate of the ingot is low. , the lighter elements are able to diffuse ferritic structure dendrites towards the interdendritic regions 20 of austenitic or all-part structure and to concentrate during the period of coexistence of the ferritic and austenitic structures. The risk that the solubility of these light elements is exceeded locally in the interdendritic regions is accentuated. When the concentration in light elements exceeds this solubility, it appears then in the steel microscopic gas pockets containing these light elements.

Moreover, during the end of cooling, the austenite of the interdendritic regions tends to transform locally into martensite when the temperature of the steel falls below the martensitic transformation temperature M s, which is slightly above ambient temperature (FIGS. 5 and 6). However, martensite has a threshold of solubility in light elements even lower than other metallurgical structures and that austenite. There is therefore more microscopic gaseous phase within the steel during this martensitic transformation.

During subsequent deformations that the steel undergoes during hot forming (eg forging), these phases flatten in sheet form.

Under fatigue stress, these leaves act as stress concentration sites, which are responsible for premature crack initiation by reducing the energy required to initiate cracks. There is thus a premature failure of the steel, which corresponds to the low values of the fatigue resistance results.

These conclusions are corroborated by the observations of the inventors, as shown by the electron micrograph of FIG.

In this photograph of a fracture surface of a stainless steel martensitic, there is a substantially globular zone P from which radiates fissures F. This zone P the footprint of the gaseous phase consisting of light elements, and which is the origin of the formation of these fissures F which, by propagating and agglomerating, created a zone of macroscopic fracture.

The inventors have carried out tests on stainless martensitic steels, and found that when performing on these steels a precautionary heat treatment according to the invention, during the cooling of the ingot immediately after the exit of the crucible ESR, as well as immediately after each of the austenitic thermal cycles at a austenitic temperature (which may include hot forming) performed subsequent to the ESR remelting, the fatigue results are improved. Such precautionary heat treatment is described below, corresponding to a first embodiment of the invention.

According to the first embodiment of the invention, the ingot is, during its cooling at the outlet of the austenitic thermal cycle, or after its exit from the crucible ESR and before the skin temperature of the ingot is less than the martensitic transformation start temperature Ms, placed and maintained in an oven whose temperature, called the holding temperature, is between the start temperatures and at the end of cooling ferrit-pearlitic transformation, Ar1 and Ar3 ("ferrito-pearlitic nose", region to the right of the curve F, FIGS. 5 and 6), for at least one holding time t, as soon as the temperature of the The coldest point of the ingot has reached the holding temperature. This time is greater than (for example at least twice) the time required to transform the austenite as completely as possible into a ferrito-pearlitic structure at this holding temperature.

The mechanisms are illustrated by the diagrams of Figures 5 and 6, and in particular by the cooling curves C1, C2, and C3, already discussed above. These cooling curves show the average evolution of the ingot temperature (surface and core) for different increasing thicknesses. This temperature begins to decrease from an austenitic temperature. Before the austenitic regions turn into martensite, i.e., before the skin temperature of the ingot becomes less than Ms, this ingot is then placed and held in an oven. The cooling curve thus becomes horizontal (curve 4 in FIG. 5 which corresponds to the treatment according to the invention).

When the ferrito-pearlitic transformation is complete (curve 4 enters the FP region to the right of curve F), the ingot is allowed to cool to room temperature.

Once at room temperature, it is possible to deposit the ingot on any surface, for example on the ground. Being able to thus deposit the ingot at a time of manufacture makes it possible to considerably increase the flexibility in the manufacturing workshops to improve logistics and costs.

During cooling from the austenitic temperature, the temperature of the ingot is most often greater than 300 ° C, which promotes the diffusion of light elements within the ingot. At the moment when the surface temperature of the ingot becomes greater than that at the heart of the ingot, degassing occurs in the ingot, which advantageously reduces the content of gaseous elements therein. The inventors have experimentally found that when, during each cooling following an austenitic thermal cycle, and during the cooling after its exit from the ESR crucible, a precautionary heat treatment is carried out on the ingot as described above, the formation is reduced. of gaseous phases of light elements within the ingot.

Indeed, there is no longer any variation in the concentration of light elements (H, N, O) from one zone to another of the ingot, and therefore there is less risk of exceeding the solubility of these phases in a zone. ingot data. Consequently, no preferential concentration of light elements is created in this or that zone.

After the precautionary heat treatment according to the first embodiment of the invention, it is possible to subject the ingot one or more austenitic cycles.

Another precautionary heat treatment is described below, corresponding to a second embodiment of the invention.

According to the second embodiment of the invention, during cooling from an austenitic temperature (temperature above the austenitic transformation end temperature to Ac3 heating), the ingot is placed, before its minimum temperature (normally the temperature of skin) is less than the martensitic transformation start temperature Ms, in an oven whose temperature is higher than the temperature Ac3. It is in the case where a subsequent austenitic thermal cycle is provided at a temperature greater than Ac3 just after cooling following a previous austenitic cycle or according to the ESR method). The ingot is then maintained in this oven at least the time necessary for the coldest part of the ingot to become greater than Ac3, the ingot then being immediately subjected to the subsequent austenitic thermal cycle. Curve 5 in FIG. 5 corresponds to this treatment according to the invention.

If, after this subsequent austenitic thermal cycle, one or more other austenitic thermal cycles are carried out, the holding in the furnace of the ingot as described above between two successive austenitic thermal cycles is carried out.

Indeed, the inventors have experimentally found that when it is ensured that the minimum temperature of the ingot between two austenitic thermal cycles does not become lower than the temperature Ms beginning of martensitic transformation, the formation of gaseous phases of light elements within the ingot is reduced.

Indeed, it remains then, within the ingot, always homogeneous in austenitic structure, homogeneous in concentration in light elements, and therefore the risk of exceeding the solubility level of the gaseous phases in a given zone of the ingot is constant, and is less.

In addition, during this cooling from the austenitic temperature, the temperature of the ingot is most often greater than 300 ° C, which diffusion of light elements within the ingot. At the moment when the surface temperature of the ingot becomes greater than or equal to that at the heart of the ingot, degassing occurs in the ingot, which advantageously reduces the content of gaseous elements therein.

In addition, at austenitic temperatures, diffusion of the alloying elements from the high concentration zones to the low concentration zones allows a reduction of the intensity of the segregations in alphagenes elements in the dendrites 10, and a reduction of the intensity of segregations in gammagens in interdendritic regions 20. The reduction of the intensity of segregations in these gammagens consequently reduces the difference in solubility of dendrites and interdendritic regions 20 in light elements (H, N, O), allowing better homogeneity in terms of structure (less coexistence of austenitic and ferritic structures) and chemical composition including light elements.

The term "segregation intensity" of an element is the difference between the concentration of this element in an area where this concentration is minimal, and the concentration of this element in an area where this concentration is maximum.

After the last austenitic thermal cycle, the ingot is maintained in the ferritic-pearlitic transformation nose for a time sufficient to obtain a quasi-complete ferrito-pearlitic transformation, in accordance with the first embodiment of the invention, which allows to deposit the ingot at room temperature.

For example, in the case of a stainless martensitic steel Z12CNDV12 (AFNOR standard) used by the inventors in the tests, the ferrito-pearlitic transformation nose is in the temperature band T between 550 ° C. and 770 ° C. Temperatures T between 650 ° C and 750 ° C are optimal, and the ingot must be maintained for a time t varying between 10 hours and 100 hours. For temperatures between one hand between 550 ° C and 650 ° C, and on the other hand between 750 ° C and 770 ° C, the holding time varies between 100 and 100OOh.

For such a steel, the temperature Ms is of the order of 200 ° C - 300 ° C.

The inventors have found that one of the precautionary thermal treatments against gaseous phases, as described above, was especially necessary when;

The maximum dimension of the ingot before cooling is less than about 910 mm or the minimum dimension is greater than 1500 mm, and the H content of the ingot before slag remelting is greater than 10 ppm, and

- The maximum dimension of the ingot before cooling is greater than about 910 mm and the minimum dimension of the ingot is less than about 1500 mm, and the H content of the ingot before slag remelting is greater than 3 ppm.

The maximum dimension of the ingot is that of the measurements in its most massive part, and the minimum dimension of the ingot is that of the measures in its least massive part:

at. immediately after slag remelting when the ingot does not undergo hot forming before its subsequent cooling.

b. When the ingot undergoes hot forming after slag remelting, just before its subsequent cooling. Preferably, the slag is dehydrated before use in the ESR crucible. Indeed, it is possible that the concentration of H in the steel ingot from the ESR slag remelting is greater than the concentration of H in this ingot before its slag remelting. In this case, hydrogen can pass from slag to ingot during the ESR process. By dehydrating the slag beforehand, the amount of hydrogen present in the slag is minimized, and thus the amount of hydrogen that could pass from slag to ingot during the ESR process is minimized.

The inventors have carried out tests on Z12CNDV12 steels developed according to the following parameters:

Test No. 1: Cooling of the ingot at the ESR crucible outlet (8.5ppm H content) when the skin temperature is 250 ° C, put in the oven at 690 ° C and metallurgical maintenance (as soon as the coldest temperature of the ingot reaches the temperature of homogenization) of 12h, cooling to room temperature.

Cooling after discharge operation diameter between 910 and 1500mm, when the skin temperature is 300 ° C, baked at 690 ° C and metallurgical maintenance of 15h, cooling to room temperature.

Cooling after stretching operation to a diameter of less than 900 ° C to room temperature.

Test No. 2;

Cooling of the ingot at the ESR crucible outlet (content H of 7 ppm) when the skin temperature is 270 ° C., put in the oven at 700 ° C. and maintained metallurgically (as soon as the coldest temperature of the ingot reaches the homogenization temperature ) 24h, cooling to room temperature.

Cooling after discharge operation diameter between 910 and 1500mm, when the skin temperature is 400 ° C, baked at 690 ° C and metallurgical maintenance of 10h, cooling to room temperature.

Test n ° 3:

Cooling of the ingot at the ESR crucible outlet (8.5ppm H content) when the skin temperature is 450 ° C., put in the oven at 1150 ° C. for delivery. Cooling after discharge operation diameter between 910 and 1500mm, when the skin temperature is 350 ° C, baked at 690 ° C and metallurgical maintenance of 15h, cooling to room temperature.

Test n ° 4:

Cooling of the ingot at the ESR crucible outlet (content H of 12 ppm) when the skin temperature is 230 ° C., put in the oven at 690 ° C. and maintained metallurgically (as soon as the coldest temperature of the ingot reached the homogenization temperature) 24h, cooling to room temperature.

- Cooling after discharge operation diameter between 910 and 1500mm, when the skin temperature is 270 ° C, baked at 690 ° C and 24h metallurgical maintenance, cooling to room temperature.

- Cooling after stretching operation to a diameter of less than 900 ° C when the skin temperature is 650 ° C, baked at 1150 ° C for a second stretch.

- At cooling, when the skin temperature is 320 ° C, put in the oven at 690 ° C and metallurgical maintenance of 15h, cooling to room temperature. At this point, the measurement of hydrogen gave 1.9 ppm

Test n ° 5:

- Cooling ingot output ESR crucible (8.5ppm H content) when the skin temperature is 450 ° C, baked at 1150 ° C for discharge.

- Cooling after discharge operation diameter between 910 and 1500mm, when the skin temperature is 350 ° C, baked at 690 ° C and metallurgical maintenance of 15h, cooling to room temperature.

- Cooling after stretching operation to a diameter less than 900 ° C to room temperature.

The results of these tests are presented below.

The composition of the Z12CNDV12 steels is as follows: (standard

DMD0242-20 index E):

C (0.10 to 0.17%) - If (<0.30%) - Mn (0.5 to 0.9%) - Cr (11 to 12.5%)

- Ni (2 to 3%) - Mo (1.50 to 2.00%) - V (0.25 to 0.40%) - N ₂ (0.010 to 0.050%) - Cu (<0.5%) - S (<0.015%) - P (<0.025%), and satisfying criterion 4.5 <(Cr - 40.C - 2.Mn - 4. i + 6.51 + 4.Mo + l IV - 30.N The measured Martensitic transformation temperature Ms is 220 ° C.

The amount of Hydrogen measured on the ingots before slag remelting varies from 3.5 to 8.5 ppm.

Figure 1 qualitatively shows the improvements made by the method according to the invention. We obtain experimentally the value of number N of rupture cycles necessary to break a steel specimen subjected to a cyclic stress in tension as a function of the alternating pseudo stress C (this is the stress experienced by the specimen under imposed deformation, according to the standard DMC0401 of Snecma used for these tests).

Such a cyclic bias is shown schematically in FIG. 2. The period T represents a cycle. The constraint evolves between a maximum value C _ma x and a minimum value C _min .

By fatigue testing a statistically sufficient number of test pieces, the inventors obtained points N = f (C) from which they drew a mean statistical curve C-N (stress C as a function of the number N of fatigue cycles). The standard deviations on the stresses are then calculated for a given number of cycles.

In FIG. 1, the first curve 15 (in fine lines) is (schematically) the average curve obtained for a steel produced according to the prior art. This first average curve C-N is surrounded by two curves 16 and 14 in dashed fine lines. These curves 16 and 14 are located respectively at a distance of +3 σ, and -3 σι of the first curve 15, σ, being the standard deviation of the distribution of the experimental points obtained during these fatigue tests, and ± 3σι corresponds in statistics to a confidence interval of 99.7%. The distance between these two dashed lines 14 and 16 is therefore a measure of the dispersion of the results. Curve 14 is the limiting factor for dimensioning a part.

In FIG. 1, the second curve 25 (in thick line) is

(schematically) the average curve obtained from the results of fatigue tests carried out on a steel produced according to the invention under a load according to FIG. 2. This second average curve CN is surrounded by two curves 26 and 24 in dashed thick lines , located respectively at a distance of +3 σ ₂ and -3 σ ₂ from the second curve 25, σ ₂ being the standard deviation of the distribution of the experimental points obtained during these fatigue tests. Curve 24 is the limiting factor for dimensioning a part.

Note that the second curve 25 is located above the first curve 15, which means that under fatigue stress at a stress level C, the steel specimens produced according to the invention breaks on average at a higher number N of cycles than that in which the steel specimens according to the prior art break.

In addition, the distance between the two curves 26 and 24 in thick dashed line is smaller than the distance between the two curves 16 and 14 in dashed fine lines, which means that the dispersion in fatigue resistance of the developed steel according to the invention is lower than that of a steel according to the prior art,

Figure 1 illustrates the experimental results summarized in Table 1 below.

Table 1 gives the results for an fatigue load of oligocyclic according to Figure 2 with a strain C _m m zero, at a temperature of 250 ° C, at N = 20 000 cycles, and N = 50 000 cycles. Oligocyclic fatigue means that the bias frequency is of the order of 1 Hz (the frequency being defined as the number of periods T per second).

Table 1

It should be noted that for a given value of the number N of cycles, the minimum value of fatigue stress required to break a steel according to the invention is greater than the minimum fatigue stress value M (set at 100%) necessary to break a steel according to the prior art. The dispersion (= 6 σ) of the results at this number N of cycles for a steel according to the invention is less than the dispersion of the results for a steel according to the prior art (dispersions expressed as a percentage of the minimum value M).

Advantageously, the carbon content of the stainless martensitic steel is lower than the carbon content below which the steel is hypoeutectoid, for example a content of 0.49%. Indeed, such a low carbon content allows a better diffusion of alloying elements and a lowering of the resetting temperatures of the primary or noble carbides, resulting in better homogenization.

For example, martensitic steel, before its slag remelting, was made in the air.

The first embodiment according to the invention can also be applied to the ingot during its cooling at the outlet of the ESR crucible, the ingot being then subjected to no austenitic thermal cycle.

Claims

1. A method for manufacturing a stainless martensitic steel comprising a slag remelting step of an ingot of said steel then a cooling step of said ingot and then at least one austenitic thermal cycle consisting of heating said ingot above its austenitic temperature followed by a cooling step, characterized in that during each of said cooling steps;

- If said cooling step is not followed by an austenitic thermal cycle, said ingot is maintained at a holding temperature included in the ferrito-pearlitic transformation nose during a holding time longer than the time required to transform the most completely possible austenite in ferrito-pearlitic structure in this ingot at said holding temperature, said ingot being maintained at this holding temperature as soon as the temperature of the coldest point of the ingot has reached the holding temperature,

- If said cooling step is followed by an austenitic thermal cycle, said ingot is, before its minimum temperature is lower than the martensitic transformation start temperature s, be maintained, for the duration between these two austenitic thermal cycles at a temperature greater than the austenitic transformation end temperature in Ac3 heating, is maintained at said holding temperature included in the ferrito-pearlitic transformation nose as above.

2. Method of manufacturing a stainless martensitic steel according to claim 1, characterized in that it is carried out on said steel in one of the following cases:

The maximum dimension of said ingot before cooling is less than about 910 mm or the minimum dimension is greater than 1500 mm, and the H content of the ingot before slag remelting is greater than 10 ppm,

- The maximum dimension of said ingot before cooling is greater than about 910 mm and the minimum dimension of the ingot is less than about 1500 mm, and the H content of the ingot before slag remelting is greater than 3 ppm.

3. A method of manufacturing a stainless martensitic steel according to claim 1 or 2, characterized in that the slag used in said remelting step has been previously dehydrated.

4. A method of manufacturing a stainless martensitic steel according to any one of claims 1 to 3, characterized in that the carbon content of said steel is less than the carbon content below which the steel is hypoeutectoid.

5. A method of manufacturing a stainless martensitic steel according to any one of claims 1 to 4, characterized in that the holding of said ingot at a temperature is carried out by placing it in an oven,

6. A method of manufacturing a stainless martensitic steel according to claim 5, characterized in that the ingot is placed in an oven before the skin of the ingot skin is lower than the end of ferrito-pearlitic transformation Arl cooling.

7. A method of manufacturing a stainless martensitic steel comprising a slag remelting step of an ingot of said steel and a cooling step of said ingot, characterized in that during said cooling step said ingot is maintained at a holding temperature. included in the ferrito- pearlitic transformation nose during a holding time longer than the time necessary to transform the austenite as completely as possible into a ferrito-pearlitic structure in this ingot at said holding temperature, said ingot being maintained at this temperature of maintaining as soon as the temperature of the coldest point of the ingot has reached the holding temperature, the ingot does not undergo an austenitic thermal cycle after said slag remelting step.