EP2488672A1

EP2488672A1 - Homogenization of martensitic stainless steel after remelting under a layer of slag

Info

Publication number: EP2488672A1
Application number: EP10781969A
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: CN102575313B; JP5868859B2; BR112012008520B1; CA2777034A1; CN102575313A; RU2536574C2; US8911527B2; EP2488672B1; JP2013507530A; FR2951197A1; WO2011045513A1; RU2012119594A; US20120260771A1; CA2777034C; FR2951197B1; BR112012008520A2

Abstract

The present 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, followed by a step in which the ingot is cooled. Before the skin temperature of the ingot resulting from the slag remelting step drops below the martensitic transformation temperature M_s of the steel, the ingot is placed in a furnace, the initial temperature T₀ of which is then above the cooling‑induced pearlite transformation finish temperature Ar1 of the steel. In the furnace, the ingot is subjected to a homogenization treatment at least for a holding time t after the temperature of the coolest point in the ingot has reached a homogenization temperature T, said holding time t being equal to at least one hour and the homogenization temperature T varying between approximately 900°C and the burning temperature of the steel.

Description

HOMOGENIZATION OF STAINLESS MARTENSTIC STEELS AFTER REFUSION UNDER DAIRY

The present invention relates to a method for manufacturing a stainless martensitic steel comprising a slag remelting step of an ingot of this steel and a cooling step of this ingot

In the present invention, 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 these low values, and also 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 into which a slag (mineral mixture, for example lime, fluoride, magnesia, alumina, spath) has been poured so that the lower end of the ingot quenches in the slag. . Then we do pass an electric current 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 object is achieved by virtue of the fact that the ingot from the slag remelting is, before the skin temperature of this ingot is less than the martensitic transformation temperature Ms of the steel, placed in an oven whose initial temperature T ₀ is then greater than the end of pearlitic transformation temperature in Arl cooling of said steel, this ingot being subjected in this oven to a homogenization treatment for at least one holding time t after the temperature of the coldest point of the ingot has reached a homogenization temperature T, this holding time t being equal to at least one hour, and the homogenization temperature T varying between about 900 ° C and the burn temperature of the steel.

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.

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 the 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-like elements (application of the known rule segments on the phase diagram). An alphagene element is one that promotes a ferritic type structure (more stable structures 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 D and F curves or in the FP region. 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 to say, it will cool more slowly the ingot to obtain a ferritoperlitic structure.

Each of Figures 5 and 6 shows three cooling curves from an 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.

In fact, once the material has solidified, the dendrites 10 first turn into ferritic structures during cooling (by crossing the curves D and F of FIG. 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 structures ferritic, therefore tend to concentrate in the interdendritic regions 20. This concentration is increased by the higher content of gamma-elements in the interdendritic regions 20. At temperatures below 300 ° C, the light elements only diffuse at extremely low speeds and remain trapped in their area. 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 ES outlet (or the subsequently deformed ingot) has a large diameter (or, more generally, the larger the dimension of the ingot is large) or the ingot cooling rate is lower , 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 interdendritic regions tends to locally transform into martensite when the temperature of the steel falls below the martensitic transformation temperature Ms, which is slightly above the temperature. ambient (Figures 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 a stress in fatigue, these leaves act as sites of concentration of constraints, which are responsible for the primer premature cracking 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 is the imprint of the gaseous phase consisting of the light elements, and which is at 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 have found that when, immediately after the ESR step, a specific homogenization treatment is carried out on the ingot taken out of the ESR crucible, the formation is reduced. of gaseous phases of light elements.

Indeed, by diffusion of the alloying elements of the high concentration zones towards the low concentration zones, a reduction of the intensity of the segregations in alphagenes elements in the dendrites 10 is allowed, and a reduction of the intensity of the segregations in gammagens elements in interdendritic regions 20. The reduction of the intensity of segregations in these gammagens has the following consequences: a smaller shift towards the right of the curves D and F of transformation into a ferrito-pearlitic structure (Figure 6), a lower structural difference between the dendrites 10 and the interdendritic regions 20, and a lesser difference in solubility in light elements (H, N, O) between the dendrites and the interdendritic regions, allowing a better homogeneity in terms of structure (less cohabitation of structures austenitic and ferritic) and chemical composition including light elements.

In addition, the homogenization treatment also leads to a homogenization of the martensitic transformation temperature Ms.

When the temperature of the steel is at a temperature above 300 ° C, the diffusion of the alloying elements is far from negligible. Moreover, if the temperature gradient makes it possible to have a warmer surface That the center of the ingot, which the conditions of recovery proposed by the inventors allow, the light elements diffuse towards the surface, which reduces their overall content in the steel.

With regard to the particularities of the homogenization treatment, the inventors have found that satisfactory results are obtained when the ingot is subjected in this oven to a homogenization treatment during a holding time t after the temperature of the most The cold of this ingot has reached a homogenization temperature T, this time t being equal to at least one hour, and the homogenization temperature T varying between a temperature T _min and the burn temperature of this steel.

The temperature T _m j _n is approximately equal to 900 ° C. The burning temperature of a steel is defined as the temperature in the raw state of solidification at which the grain boundaries in the steel are transformed (or even liquefied), and is greater than T _min . This time t of maintaining the steel in the furnace therefore varies inversely with this homogenization temperature T.

For example, in the case of a stainless martensitic steel Z12CNDV12 (AFNOR standard) used by the inventors in the tests, the homogenization temperature T is 950 ° C., and the corresponding holding time t is equal to 70 hours. When the homogenization temperature T is 1250 ° C. which is slightly lower than the burn temperature, then the corresponding holding time t is equal to 10 hours.

For example, the homogenization temperature T is selected from a range selected from the group consisting of the following ranges: 950 ° C to 1270 ° C, 980 ° C to 1250 ° C, 1000 ° C to 1200 ° C.

For example, the minimum hold time t is selected from a range selected from the group consisting of the following ranges: 1 hour to 70 hours, 10 hours to 30 hours, 30 hours to 150 hours.

Furthermore, the inventors have found that satisfactory results are obtained when the ingot at the outlet of the ESR crucible is placed in an oven whose initial temperature T ₀ is greater than the end of pearlitic transformation temperature in cooling Arl of this steel. and when the skin temperature of this ingot remains higher than the martensitic transformation temperature Ms of this steel. In the case where the initial temperature T ₀ of the oven is lower than the homogenization temperature T, the temperature of the oven is, after the ingot has been placed in this oven, increased to a temperature at least equal to the temperature homogenization. Thus, during this rise in temperature, it tends to a homogeneous austenitic structure, in order to homogenize the hydrogen content, and tends to a rising temperature gradient from the center of the part to the surface. The temperature in the center of the ingot therefore remains lower than the skin temperature of the ingot throughout the rise in temperature. This allows a global degassing and more effective ingot.

Alternatively, the initial temperature T ₀ of the oven may be greater than the homogenization temperature, in which case the oven temperature is simply maintained above this homogenization temperature.

The inventors have found that the homogenization treatment is especially necessary when:

The maximum dimension of the ingot is less than approximately 910 mm, and the H content of the ingot before slag remelting is greater than 10 ppm, and

The maximum dimension of the ingot 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, and

- The minimum dimension of the ingot is greater than 1500 mm and the H content of the ingot before slag remelting is greater than 10

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;

· Immediately after slag remelting when the ingot is not hot-formed before it is subsequently cooled.

• When the ingot undergoes heat shaping after slag remelting, just before its subsequent cooling. As indicated above, the inventors have found that the concentrations of light elements may be greater (greater than 10 ppm) when the minimum dimension of the ingot or of the deformed ingot is greater than a large size threshold (in this case 1500 mm). The explanation of the existence of a high threshold (1500 mm) for the minimum dimension of the ingot is as follows: when the minimum dimension of the ingot is greater than this threshold, we approach the case of slow cooling (curve C3) in which there is almost no structural difference between dendrites and interdendritic regions during cooling. In addition, the cooling rate is sufficiently low so that the temperature is substantially homogeneous between the core of the skin of the ingot, so that the diffusion of light elements to the surface is facilitated, which allows a greater degassing. On the other hand, when the minimum dimension of the ingot is below this threshold, the core of the ingot is, during the cooling, much hotter than its surface, which favors a diffusion of the light elements towards the core and slows down the degassing.

Furthermore, it is preferable that the slag is previously dehydrated before use in the ESR crucible, because it minimizes the amount of hydrogen present in the slag, and thus minimizes the amount of hydrogen that could pass the slag ingot during the ESR process.

The inventors have carried out tests on Z12CNDV12 steels produced with the process according to the invention, that is to say with a homogenization carried out immediately after the ingress of the ingot of the ESR crucible according to the following parameters:

Test No. 1: skin temperature of the ingot at 250 ° C., put in the oven at 400 ° C., rise of the oven at the homogenization temperature of 1250 ° C., metallurgical maintenance (as soon as the coldest temperature of the ingot reaches the homogenization temperature) of 75h, cooling to room temperature.

Test No. 2: Skin temperature of the ingot at 600 ° C., put in the oven at 450 ° C., rise of the furnace at the homogenization temperature of 1000% metallurgical maintenance (as soon as the coldest temperature of the ingot reaches the temperature homogenization) of 120h, cooling to room temperature. The results of these tests are presented below,

The composition of the Z12CNDV12 steels is the following (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 the criterion 4.5≤ (Cr - 40.C - 2.Mn - 4.Ni + 6.Si + 4.Mo + ll.V - 30.N) <9

The martensitic transformation temperature Ms measured 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. The value of the number N of rupture cycles necessary to break a steel specimen subjected to a cyclic stress in tension as a function of the pseudo-alternating stress C is obtained experimentally (this is the stress experienced by the test specimen under imposed deformation. , according to Sncma DMC0401 standard 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 CN is surrounded by two curves 16 and 14 in dashed fine lines. These curves 16 and 14 are situated respectively at a distance of +3 <s _{ and -3 at _t from the first curve 15, where σ * is the standard deviation of the distribution of the experimental points obtained during these fatigue tests. and ± 3σι statistically corresponds to a confidence interval of 99.7%. The distance between these two curves 14 and 16 in dashed line is therefore a measure of the dispersion of 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. CN average curve is surrounded by two curves 26 and 24 dashed thick line, respectively located at a distance of +3 σ ₂ and -3 σ ₂ of the second curve 25, σ ₂ being the standard deviation of the distribution of points experimental results obtained during these fatigue tests. Curve 24 is the limiting factor for dimensioning a part.

It is noted 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 test pieces produced according to the invention break on average to a number N of cycles higher than that where the steel test pieces according to the prior art are broken.

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 a fatigue load of oligocyclic fatigue according to Figure 2 with a stress C _mln 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

Condil: test ions in

Steel according to 1 prior art Steel developed stil the invention tired oligocyclic

N Temperature Qnin Dispersion Qnin Dispersion

2-10 ⁵ 200 ° C 100% = M 120% M 130% M 44% M 5-10 ⁴ 400 ° C 100% = M 143% M 130% M 90% M

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 the alloying elements and a lowering of the temperatures of solution of the primary or noble carbides, which leads to a better homogenization.

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

Claims

1. A method for 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 the ingot from the slag remelting is, before the temperature of skin of said ingot is less than the martensitic transformation temperature Ms of said steel, placed in a furnace whose initial temperature T ₀ is then greater than the end of pearlitic transformation temperature in cooling Arl of said steel, said ingot being subjected in this furnace at a homogenization treatment for at least one holding time t after the temperature of the coldest point of said ingot has reached a homogenization temperature T, this holding time t being equal to at least one hour, and the temperature homogenization T varying between about 900 ° C and the burn temperature of said steel.

2. Process for manufacturing a stainless martensitic steel according to claim 1, characterized in that said initial temperature T ₀ of the furnace is lower than said homogenization temperature T, the temperature of the furnace being increased from its initial temperature T _{D to} at a temperature at least equal to the homogenization temperature T.

3. A method of manufacturing a stainless martensitic steel according to claim 1, characterized in that the homogenization temperature T is within a range selected from the group consisting of the following ranges: 950 ° C to 1270 ° C, 980 ° C at 1250 ° C, 1000 ° C to 1200 ° C.

4. A method of manufacturing a stainless martensitic steel according to claim 1 or 2, characterized in that minimum holding time t belongs to a range selected from the group consisting of the following ranges: 1 hour to 70 hours, 10 hours to 30 hours. hours, 30 hours to 150 hours.

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

6. A method of manufacturing a stainless martensitic steel according to claim 1 or 2, characterized in that said time of maintenance t varies inversely with the variation of said homogenization temperature T.

7. A method of manufacturing a stainless martensitic steel according to any one of claims 1 to 6, 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 approximately 910 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 its minimum dimension less than about 1500 mm, and the H content of the ingot before slag remelting is greater than 3 ppm.

- The minimum dimension of the ingot is greater than 1500 mm and the H content of the ingot before slag remelting is greater than 10 ppm.

8. A method of manufacturing a stainless martensitic steel according to any one of claims 1 to 7, characterized in that the carbon content of said steel is less than the carbon content below which the steel is hypoeutectoid.