FERRITIC STAINLESS STEEL SHEET AND ASSOCIATED PRODUCTION METHOD

Abstract

The invention relates to a ferritic stainless steel sheet the composition of which comprises, the contents being expressed by weight: C ≤ 0.03 % 0.25 % ≤ Mn ≤ 1 ⁢ % , preferably 0.3 % ≤ Mn ≤ 0.5 % 0 ⁢ % ≤ Si ≤ 0.2 % , preferably ⁢ Si ≤ 0.15 % with ⁢ Mn / Si ≥ 1.2 S ≤ 0.005 % P ≤ 0.04 % 19. % ≤ Cr ≤ 24. % Ni ≤ 0.5 % Mo ≤ 0.1 % N ≤ 0.03 % Cu ≤ 0.2 % 0.4 % ≤ Nb ≤ 1. % 0.05 % ≤ Ti ≤ 0.2 % , preferably 0.05 % ≤ Ti ≤ 0.15 % Zr ≤ 0.02 % Al ≤ 0.02 % V ≤ 0.2 % Co ≤ 0.05 % Sn ≤ 0.05 % , Rare ⁢ earth ≤ 800 ⁢ ppm it being understood that: V + Zr + Al ≤ 0.2 % Ti + V + Zr + Al ≤ 0.3 % Ti + Nb ≤ 1. % Ni + Cu + Co ≤ 0.6 % 2 × Nb - 7 × C ≥ 0.8 % 0 ⁢ % ≤ Ti - 4 × N ≤ 0.15 % 0.2 ppm ≤ Ca ≤ 20 ⁢ ppm 1 ⁢ ppm ≤ O ≤ 60 ⁢ ppm the remainder of the composition consisting of iron and unavoidable impurities resulting from the preparation, the sheet being an annealed and pickled sheet, the sheet comprising a volume fraction of Fe2Nb Laves phases of less than 0.2%.

Description

The present invention relates to a ferritic stainless steel sheet, and to an associated production method.

The development of high temperature electrochemical applications (fuel cells or solid oxide electrolyzer cells) to convert hydrogen compounds to electricity, or conversely to produce hydrogen products from carbon-free electricity, requires the use of novel materials able to be used at between 500° C. et 1000° C. for tens of thousands of hours with several tens of stop-start cycles. The constituent anodes, cathodes and electrolytes forming the cells of these systems are prepared from ceramics or cermet, the type of which can vary as a function of design. These cells are supported by or connected to metal parts called Interconnectors having the role of distributing reactive gases within the cell and collecting electrons. The low coefficient of thermal expansion of the cells, in which the electrolyte is often based on yttria-stabilized zirconia (10, 5×10⁻⁶ K⁻¹), requires the use of high-chromium ferritic stainless steels having a very close coefficient of thermal expansion and capable of resisting oxidation at these high temperatures. In addition, the interconnectors must not cause degradation of cell performance by contamination of the anode and cathode, in particular via diffusion of the chromium or excessively strong electrical resistivity of the oxide formed on the surface thereof.

The ferritic stainless steels currently used have very high chromium content with controlled, very low content of some residuals such as silicon, and also contain rare earths (lanthanum in particular). The vacuum manufacturing thereof from pure raw materials is therefore costly. Since the interconnectors account for a major proportion by weight in these systems, the competitivity of these electrolyzers is impacted. The idea therefore is to use more conventional ferritic stainless steels produced in an electric furnace such as the steels AISI 441 or 444 which have good heat-resistant properties. However, experiments show that, under some conditions of use, the performance of said steels degrades more rapidly despite an oxide layer of comparable composition and thickness. The cause of this faster degradation of performance is segregation of the silicon at the metal-oxide interface, which forms a thin film of silica (silicon oxide) that is more or less continuous and electrically highly resistive with resistivity in the region of 10⁶ Ω·cm at 850° C. This value is to be compared with the resistivity of chromium oxide Cr₂O₃, which is in the region of 100 Ω·cm at 850° C. The silica film is of very narrow thickness and typically has a thickness of approximately a few tens of nanometres at 850° C. after an operating time of 1000 hours. However, having regard to the electrical resistivity of silica, it makes a strong contribution towards increasing Area Specific Resistance (ASR). Yet, silicon cannot be fully excluded from current production methods at reasonable cost.

It is therefore one objective of the invention to propose a ferritic stainless steel sheet having conductivity properties comparable to those of steels with very high chromium content, even after use over extended times at high temperature.

Preferably, it is also sought to improve the creep resistance of the sheet since the system is also subjected to thermomechanical loads likely to generate deformations leading to modification of the contacts and functions of the interconnectors.

For this purpose, the subject of the invention is a ferritic stainless steel sheet having the following composition, the contents being expressed by weight:

$C\le 0.03\%$ $0.25\%\le \mathrm{Mn}\le 1\%,\mathrm{preferably}0.3\%\le \mathrm{Mn}\le 0.5\%$ $0\%\le \mathrm{Si}\le 0.2\%,\mathrm{preferably}\mathrm{Si}\le 0.15\%$ $\mathrm{with}\mathrm{Mn}/\mathrm{Si}\ge 1.2$ $S\le 0.005\%$ $P\le 0.04\%$ $19.\%\le \mathrm{Cr}\le 24.\%$ $\mathrm{Ni}\le 0.5\%$ $\mathrm{Mo}\le 0.1\%$ $N\le 0.03\%$ $\mathrm{Cu}\le 0.2\%$ $0.4\%\le \mathrm{Nb}\le 1.\%$ $0.05\%\le \mathrm{Ti}\le 0.2\%,\mathrm{preferably}0.05\%\le \mathrm{Ti}\le 0.15\%$ $\mathrm{Zr}\le 0.02\%$ $\mathrm{Al}\le 0.02\%$ $V\le 0.2\%$ $\mathrm{Co}\le 0.05\%$ $\mathrm{Sn}\le 0.05\%,$ $\mathrm{Rare}\mathrm{earth}\le 800\mathrm{ppm}$

• • it being understood that:

$V+\mathrm{Zr}+\mathrm{Al}\le 0.2\%$ $\mathrm{Ti}+V+\mathrm{Zr}+\mathrm{Al}\le 0.3\%$ $\mathrm{Ti}+\mathrm{Nb}\le 1.\%$ $\mathrm{Ni}+\mathrm{Cu}+\mathrm{Co}\le 0.6\%$ $2\times \mathrm{Nb}-7\times C\ge 0.8\%$ $0\%\le \mathrm{Ti}-4\times N\le 0.15\%$ $0.2\mathrm{ppm}\le \mathrm{Ca}\le 20\mathrm{ppm}$ $1\mathrm{ppm}\le O\le 60\mathrm{ppm}$

• • the remainder of the composition consisting of iron and unavoidable processing impurities,
  • the sheet being annealed and pickled, the sheet having a volume fraction of Fe₂Nb Laves phases of less than 0.2%.

The ferritic steel sheet of the invention may also have one or more of the following characteristics taken alone or in any technically possible combination:

• • the alloy has a content of rare earths of between 50 ppm and 800 ppm;
  • the niobium content of the alloy heeds the following relation: Nb-10×(C+N)≥0%.
  • the sheet has a thickness of between 0.1 mm and 2.5 mm;
  • the sheet has an average grain size of between 30 micrometres and 80 micrometres when the sheet has a thickness of between 1.2 mm and 2.5 mm, and an average grain size of between 15 micrometres and 80 micrometres when the sheet has a thickness greater than or equal to 0.1 mm and less than 1.2 mm;
  • the sheet is a cold rolled and annealed sheet;
  • when subjected to heat treatment at a temperature of 850° C. for a time of 1000 hours, the sheet comprises a volume fraction of Fe₂Nb Laves phases greater than or equal to 0.8%;
  • when subjected to heat treatment at a temperature of 850° C. for a time of 1000 hours, the sheet comprises a volume fraction of Fe₃Nb₃X cubic phases of less than 0.05%;
  • when subjected to heat treatment at a temperature of 850° C. for a time of 1000 hours, the sheet on each of its surfaces comprises a layer of oxides and, at the interface between the steel of the sheet and the layer of oxides, comprises precipitates of silicon oxide, such that the surface fraction of the precipitates of silicon oxide at the interface between the steel of the sheet and the layer of oxides is less than or equal to 0.35; and
  • the layer of oxides has a thickness of less than or equal to 10 μm.

The invention also relates to a production method of a ferritic stainless steel sheet, comprising the following steps:

• • a steel is prepared having the composition described above;
  • a semi-processed product is cast from this steel;
  • the semi-processed product is brought to a temperature higher than or equal to 1150° C. and lower than or equal to 1260° C. for a time of between 40 minutes and 60 minutes, and the semi-processed product is hot rolled to obtained a hot rolled sheet of thickness between 2.5 mm and 6 mm;
  • the hot rolled sheet is annealed;
  • the hot rolled, annealed sheet is pickled;
  • said hot rolled sheet is cold rolled at a temperature between ambient temperature and 300° C., in a single step or in several steps separated by intermediate annealing operations;
  • a final annealing of the cold rolled sheet is conducted at a temperature of between 1000° C. and 1100° C. and for a time of between 10 seconds and 6 minutes, to obtain a fully recrystallized structure.

The production method of the invention may also comprise one or more of the following characteristics taken alone or in any technically possible combination:

• • the annealing of the hot rolled sheet is conducted at a temperature of between 1000° C. and 1100° C., for a time of 30 seconds to 6 minutes;
  • the intermediate annealing operation(s) are conducted at a temperature of between 950° C. and 1100° C. for a time of 30 seconds to 6 minutes; and
  • the final annealing is conducted at a temperature of between 1050° C. and 1090° C.

The invention will be well understood and other aspects and advantages will become apparent on reading the following detailed description of an example of embodiment given with reference to the appended Figures in which:

FIG. 1 schematically illustrates a sheet of the invention in cross-sectional view after ageing heat treatment.

FIG. 2 schematically illustrates determining of the surface fraction of silica precipitates at the metal-oxide interface after said ageing treatment.

The invention concerns a ferritic stainless steel sheet having the following composition, the contents being expressed by weight:

$C\le 0.03\%$ $0.25\%\le \mathrm{Mn}\le 1\%,\mathrm{preferably}0.3\%\le \mathrm{Mn}\le 0.5\%$ $0\%\le \mathrm{Si}\le 0.2\%,\mathrm{preferably}\mathrm{Si}\le 0.15\%$ $\mathrm{with}\mathrm{Mn}/\mathrm{Si}\ge 1.2$ $S\le 0.005\%$ $P\le 0.04\%$ $19.\%\le \mathrm{Cr}\le 24.\%$ $\mathrm{Ni}\le 0.5\%$ $\mathrm{Mo}\le 0.1\%$ $N\le 0.03\%$ $\mathrm{Cu}\le 0.2\%$ $0.4\%\le \mathrm{Nb}\le 1.\%$ $0.05\%\le \mathrm{Ti}\le 0.2\%,\mathrm{preferably}0.05\%\le \mathrm{Ti}\le 0.15\%$ $\mathrm{Zr}\le 0.02\%$ $\mathrm{Al}\le 0.02\%$ $V\le 0.2\%$ $\mathrm{Co}\le 0.05\%$ $\mathrm{Sn}\le 0.05\%,$ $\mathrm{Rare}\mathrm{earth}\le 800\mathrm{ppm}$

• • it being understood that:

$V+\mathrm{Zr}+\mathrm{Al}\le 0.2\%$ $\mathrm{Ti}+V+\mathrm{Zr}+\mathrm{Al}\le 0.3\%$ $\mathrm{Ti}+\mathrm{Nb}\le 1.\%$ $\mathrm{Ni}+\mathrm{Cu}+\mathrm{Co}\le 0.6\%$ $2\times \mathrm{Nb}-7\times C\ge 0.8\%$ $0\%\le \mathrm{Ti}-4\times N\le 0.15\%$ $0.2\mathrm{ppm}\le \mathrm{Ca}\le 20\mathrm{ppm}$ $1\mathrm{ppm}\le O\le 60\mathrm{ppm}$

• • the remainder of the composition being composed of iron and unavoidable processing impurities.

Regarding the chemical composition of the steel, the carbon increases the mechanical characteristics at high temperature, in particular creep resistance. However, on account of the very low solubility thereof in ferrite, carbon tends to precipitate in the form of carbides M₂₃O₆ or M₇C₃ at a temperature lower than about 900° C. This precipitation, generally located at the grain boundaries, can lead to depletion of chromium in the vicinity of these boundaries and hence to susceptibility to intergranular corrosion. This susceptibility can be encountered in particular in welding Heat Affected Zones, heated to a very high temperature. The carbon content must therefore be limited to 0.03% at most, to obtain satisfactory resistance to intergranular corrosion and to avoid reducing workability. In addition, the carbon content must heed a relation with niobium is will be explained below.

Chromium is an essential element for stabilization of the ferritic phase, and to increase resistance to oxidation. Together with the other elements of the composition, the minimum content thereof must be higher than or equal to 19.0%, to obtain a ferritic structure at any temperature and good resistance to cyclic oxidation, in particular when the sheet is of narrow thickness (less than or equal to 0.5 mm) and this thickness limits the reservoir of chromium available for oxidation. On ageing, and as illustrated in FIG. 1, an oxide layer 2 comprising an inner layer 4 of chromium oxide Cr₂O₃ and optionally an outer layer 5 of manganese-rich chromium oxide is formed on the surface of the metal substrate 1 and protects the steel over very long periods of time and at high temperatures. The maximum content of chromium must not however exceed 24.0% otherwise this would excessively increase mechanical strength at ambient temperature, generating substantial brittleness and degrading workability.

In the context of this invention, the terms «inner» and «outer» are used in relation to closeness with the metal substrate 1, an inner layer being closer to the metal substrate 1 than an outer layer.

The alloy has a manganese content of between 0.25% by weight and 1% by weight. At these contents, manganese increases the mechanical characteristics of the alloy and additionally allows the formation of an outer layer 5 of manganese-rich chromium oxide and possibly containing iron of spinel type (Mn,Fe) Cr₂O₄. The good thermodynamic stability of these oxides allows limiting of chromium evaporation at high temperature in the presence of water vapour. This outer layer 5 of manganese-rich chromium oxide also promotes good adhesion of a protective coating e.g. a coating of LSM type (strontium-doped lanthanum manganite) or MCO type (manganese cobalt oxide, spinel). These spinels have very good electrical conductivity with resistivity in the region of 20 Ω·cm at 850° C. However, over and above 1% by weight, the oxidation kinetics under heat become too rapid and a thick, strongly adhering oxide layer is developed, making pickling operations difficult when producing the sheet. Manganese is also a gammagenic element which must be limited in ferritic steels, similar to Ni, Cu, Co. The manganese content is therefore limited to 1%.

Preferably, the manganese content is between 0.3% and 0.5%.

Silicon is a very efficient element for increasing resistance to oxidation. However, the silicon oxide (silica) 3 which forms at the metal-oxide interface has a very low coefficient of expansion in the region of 1.10⁻⁶ K⁻¹, ten times lower than that of the base metal and of chromium oxide Cr₂O₃, reducing the adhesion of the oxide layer 2 overall. Low adhesion of the oxide layer 2 leads to deteriorated conductivity properties. In addition, silicon has high electrical resistivity, this being very detrimental in the targeted application since the oxidized metal must have good electrical conductivity. Silicon is also a hardening element of ferrite when cold, reducing the ductility thereof and cold workability. The silicon content must therefore be limited to a minimum amount and must not exceed 0.20% by weight, the maximum preferably being 0.15%. The silicon content is higher than 0%, on account of the unavoidable presence of silicon in trace amount. The silicon content generally remains higher than or equal to 0.05%. Reducing the Si content below this value would require costly processing.

In addition, the Mn/Si ratio must be higher than or equal to 1:2, to promote formation of the outer layer 5 of manganese-rich chromium oxide of type (Mn,Fe) Cr₂O₄ described above to the detriment of silicon oxide.

Sulfur and phosphorus are impurities which decrease hot ductility and workability. Phosphorus segregates easily at the grain boundaries reducing the cohesion thereof. Sulfur is also harmful for oxidation by segregating at the metal-oxide interface and reducing adhesion to the metal. In this respect, the sulfur and phosphorus contents must be respectively lower than or equal to 0.005% and 0.04% by weight.

Nickel is a gammagenic element which increases the ductility of steel. To maintain a single-phase ferritic structure, the content thereof is limited. In addition, nickel does not improve the targeted properties and the intentional addition thereof would increase production costs having regard to the high price of nickel. The nickel content must also be as low as possible and lower than or equal to 0.5% by weight.

Molybdenum not only increases resistance to high temperatures but also resistance to oxidation. However, in high-chromium steels containing titanium and niobium, it generates weakness of the ferritic matrix in particular in hot rolled strips of thickness between 2.5 mm and 6 mm. Molybdenum excessively reduces ductility and workability and is a costly additional element. The content thereof must be lower than or equal to 0.10%, preferably strictly lower than 0.10%.

Similar to carbon, nitrogen increases mechanical characteristics. However, nitrogen tends to precipitate at the grain boundaries in the form of nitrides, thereby reducing resistance to corrosion. To limit problems of susceptibility to intergranular corrosion, the nitrogen content must be lower than or equal to 0.03%.

Copper has a hardening effect under heat. In excessive amount however, it reduces ductility when hot rolling. Like nickel, it is also a gammagenic element that is to be limited. In this respect, the copper content must therefore be lower than or equal to 0.20% by weight.

Niobium is an element of importance in the invention. Usually, this element can be used as stabilizing element in ferritic stainless steels: the above-mentioned phenomenon of susceptibility to intergranular corrosion can be prevented through the addition of elements forming carbides or carbonitrides that are thermally very stable. In this manner carbon and nitrogen in solution are reduced as much as is possible, thereby avoiding subsequent precipitation of chromium carbides and nitrides. Niobium as well as titanium, and to a lesser extent zirconium and vanadium, therefore provide stable fixing of carbon and nitrogen.

However, niobium also combines with iron to form some intermetallic compounds in the range 650° C.-950° C.: the inventors have evidenced that intergranular precipitation of hexagonal Fe₂Nb occurring at high temperature could be used to advantage to increase mechanical properties under heat, creep in particular, and therefore under the targeted conditions of use.

Additionally, the inventors have discovered that the precipitates of Fe₂Nb compounds of hexagonal structure (called Laves phases) capture part of the silicon and therefore minimize the formation of silicon oxides, which is desired for the application. In particular, compared with Fe₃Nb₃X cubic phases, where X designates nitrogen, oxygen or carbon, also likely to be formed, Fe₂Nb Laves phases contain about four times more silicon in weight percent. Adjustment of the initial silicon content, and the capturing of part of the silicon by said phases, allow a very significant reduction in the segregation of silicon at the metal-oxide interface and the forming of highly resistive silica. In addition, the type and intergranular spatial distribution of these precipitates i.e. at the grain boundaries, strongly promote resistance to creep up to 1000° C.

To form Fe₂Nb phases under the targeted conditions of use, several conditions must be heeded: the niobium content must be between 0.40% and 1.0% and such that 2×Nb-7×C≥0.8%, and the composition must also comprise titanium in a content of between 0.05% and 0.2%, and such that 0%≤Ti−4×N≤0.15%. Preferably, the contents of niobium, carbon and nitrogen are also such that Nb−10×(C+N)≥0%.

If the content of total Nb in the steel is lower than 0.40%, the steel will be insufficiently stabilized and the quantity of Fe₂Nb precipitates formed at high temperature will be insufficient to obtain the targeted high temperature properties. To obtain this favourable precipitation of niobium, the inventors have also evidenced the importance of the content of effective niobium: effective niobium designates the quantity of niobium in solid solution available for precipitation with iron, on the assumption that the carbon and nitrogen have fully precipitated with the niobium and titanium in the form of carbonitrides TIN, TIC and Nb(C,N). To ensure a sufficient content of effective niobium, the niobium content must heed the relation 2×Nb−7×C≥0.8%, and preferably heed the relation Nb−10×(C+N)≥0%.

Conversely, a large excess of Nb causes weakening, in particular in ferrite with very high chromium content. This excess would strongly increase hardness and slow recrystallization, which also limits the cold workability of the metal. The niobium content is therefore limited to Nb≤1.0%.

To promote the formation of Fe₂Nb phases, to the detriment of Fe₃Nb₃X phases, the titanium content must be higher than or equal to 0.05%, but lower than or equal to 0.2% to limit excessive internal oxidation at high temperature. Preferably the titanium content is higher than or equal to 0.05% and/or lower than or equal to 0.15%.

Also, the titanium must heed a relation with nitrogen of 0%≤Ti−4×N≤0.15%. If not, the niobium will precipitate not in the form of Fe₂Nb of hexagonal structure on and after 650° C., but in the form of cubic Fe₃Nb₃X compounds less efficient for the trapping of silicon.

Additionally, the sum of the contents of Nb and Ti must be lower than or equal to 1.0%, to avoid making the metal too brittle and to guarantee satisfactory toughness of welds.

Vanadium, zirconium and aluminium are nitrogen-stabilizing elements, which increase mechanical strength at high temperature, but the total content of these elements should be limited however to 0.2% at most, otherwise workability would be reduced.

Also, the content of aluminium is limited to 0.02% at most, since aluminium is likely to oxidize and form alumina that is highly electrically resistive. In particular, the electrical resistivity of alumina is in the region of 107 Ω·cm at 850° C., which means that alumina is ten times more resistive than silica (SiO₂).

The content of zirconium is at most 0.02% to avoid reducing workability and limit the risk of surface defects.

The content of vanadium is at most 0.2% to avoid reducing workability.

In addition, the inventors have evidenced that the contents of titanium, aluminium, vanadium and zirconium must jointly be restricted to limit metal brittleness. The sum of their contents must be such that: Ti+Al+V+Zr≤0.30%.

Cobalt is a hardening element under heat but degrades workability. The content thereof must therefore be lower than or equal to 0.05% by weight.

Also, the cobalt content when added to that of copper and nickel must be as low as possible to prevent an increase in electrical resistivity detrimental to the application. For this purpose, the total content Ni+Cu+Co must be lower than or equal to 0.60%.

To avoid problems related to hot forging, the tin content must be lower than or equal to 0.05%.

The alloy also contains calcium in a content of between 0.2 ppm and 20 ppm.

Calcium acts during the production process to reduce steel contents of sulfur and oxygen. The concentration of calcium should be limited to prevent internal oxidation of the metal when used at high temperature.

The oxygen content of the steel is between 1 ppm and 60 ppm, the content being approximately 20 ppm for example. As is the case for calcium, the concentration thereof should be limited to prevent the presence of oxide inclusions or internal oxidation of the metal when used at high temperature.

Rare earth elements (REE) can be added to improve adhesion of the oxide layers, making the steel resistant to corrosion. However, the content of rare earths must not exceed 800 ppm. Over and above this content, production of the metal could be made difficult on account of reactions of the rare earths with the refractories lining the casting ladle. These reactions would lead to the formation of REE oxides which would degrade the inclusion cleanliness of the steel. In addition, the efficacy of REEs is sufficient at the proposed contents, and an increase in the latter would only unnecessarily increase production costs having regard to the high price of REEs, and would lead to accelerated wear of refractories. If rare earths are added, the content is preferably at least 50 ppm.

Preferably, the rare earths are chosen from among cerium, lanthanum and yttrium or combinations thereof. In particular, rare earths comprise a mixture of cerium and lanthanum. Preferably the rare earths are a mixture of cerium and lanthanum.

The steel of the invention is generally in the form cold rolled and annealed sheet, and the thickness of the sheet is generally between 0.1 mm and 2.5 mm.

In the invention, the structure of the steel in delivery condition i.e. cold rolled and annealed sheet is fully recrystallized.

The bending creep of the sheet at high temperature is a function of the thickness of the sheet. For same mechanical load and metallurgical state, a sheet of narrower thickness will undergo greater deformation on bending.

When the thickness of the cold rolled, annealed sheet is between 1.2 mm and 2.5 mm, the average grain size of the steel is preferably between 30 micrometres and 80 micrometres i.e. of ASTM number (standard ASTM E-112) between 4 and 7. Within this range of strip thickness, a grain size larger than or equal to 30 micrometres (ASTM≤7) is advantageous since it guarantees slight deformation of the sheet by bending under creep at high temperature, relative to the thickness thereof. Also, within this range of strip thickness, the grain size is preferably less than or equal to 80 micrometres (ASTM≥4). A grain size larger than 80 micrometres (ASTM<4) leads to the onset of unattractive surface irregularities known as «orange skin» appearance when worked at ambient temperature, that are harmful for good adhesion of the protective coating.

When the thickness of the cold rolled, annealed sheet is greater than or equal to 0.1 mm and less than 1.2 mm, the average grain size of the steel is preferably between 15 micrometres and 80 micrometres, i.e. an ASTM number of between 4 and 9. Within this range of strip thickness, a grain size larger than or equal to 15 micrometres (ASTM≤9) is advantageous since it guarantees slight deformation of the sheet by bending under creep at high temperature, in relation to the thickness thereof. Also, within this range of strip thickness, the grain size is preferably less than or equal to 80 micrometres (ASTM≥4). A grain size larger than 80 micrometres (ASTM<4) leads to the onset of unattractive surface irregularities known as «orange skin» appearance when worked at ambient temperature, that are harmful for good adhesion of the protective coating.

The microstructure of the cold rolled, annealed sheet in delivery condition comprises precipitates essentially composed of intragranular titanium and niobium carbonitrides. The annealing treatments performed on the sheet have the effect of dissolving intermetallic precipitates of Fe₂Nb Laves phase type of hexagonal structure, and of Fe₃Nb₃X type of cubic structure included in the microstructure.

In particular, the volume fraction of Fe₂Nb Laves phases in the cold rolled, annealed sheet in delivery condition, is less than 0.2%.

The volume fraction of Fe₂Nb Laves phases is determined as follows.

At a first stage, standard polishing is carried out with abrasives of grain size up to 1 μm, followed by electrolytic attack in 60% nitric acid on a cross-section of a test specimen taken in orthogonal direction to the direction of rolling.

Observation is performed under electron microscope in backscattered electron mode. Minimum magnification of ×1000 is used to obtain an overall view and to determine particle size with accuracy. Five images are taken per test specimen and condition. Backscattered electron mode generates a contrast of chemical composition over a scale of 256 levels, called greyscale, ranging from white (255) to black (0).

The volume fraction of Fe₂Nb and Fe₃Nb₃X intermetallic phases is determined via image analysis, for example using Image J software, of the images obtained.

Initially, the images obtained are processed with a thresholding method so as only to maintain the whitest precipitates corresponding to the Fe₂Nb and Fe₃Nb₃X intermetallic phases. In backscattered electron mode, the precipitates high in niobium and iron, on account of their atomic number, appear as the compounds with the lightest shade in the images. Consequently, the Fe₂Nb and Fe₃Nb₃X intermetallic phases can be distinguished from niobium and titanium carbonitrides by chemical contrast and appear lighter in shade. At this thresholding step, for each image, the threshold is chosen so that the Fe₂Nb and Fe Nb₃X intermetallic precipitates can be distinguished from the remainder of the image. After thresholding, the intermetallic precipitates are shown to be white in a black matrix.

Manual filtering of the images is then performed to remove artefacts such as holes or impurities.

The surface fraction of Fe₂Nb and Fe₃Nb₃X intermetallic phases is then determined from this thresholded and filtered image. In this context, it is accepted that the volume fraction is equal to the surface fraction. The volume fraction of the Fe₂Nb and Fe₃Nb₃X intermetallic phases is thus obtained.

In addition, measurements of the Fe/Nb ratio are carried out by energy-dispersive spectroscopy (EDS) in each of the images. This ratio differing in each of these intermetallic phases, it allows a distinction to be made between cubic intermetallic Fe₃Nb₃X and hexagonal Fe₂Nb, and thereby determine the volume fraction of the Fe₂Nb Laves phases from the total volume fraction of Fe₂Nb and Fe₃Nb₃X intermetallic phases.

In the cold rolled, annealed sheet, niobium is mostly in solid solution. In particular, the weight content of Nb in solid solution in the cold rolled, annealed sheet in delivered condition is at least 0.3%.

The sheet of the invention is additionally characterized in that, when subjected to heat treatment at a temperature of between 650° C. and 1000° C. for a time longer than or equal to 30 minutes, the structure of the sheet in addition to the above-mentioned titanium and niobium carbonitrides comprises a homogeneous and intergranular precipitation of Fe₂Nb compounds of hexagonal structure (Laves phases). The structure of the sheet subjected to said heat treatment may additionally comprise an intragranular precipitation of Fe₂Nb compounds, this being decreased however when the heat treatment time increases.

In particular, after ageing heat treatment in air at a temperature of 850° C. for a time of 1000 hours, the volume fraction of Fe₂Nb Laves phases in the sheet is at least 0.8%. These Fe₂Nb Laves phases are predominantly intergranular.

Ageing heat treatment in air at a temperature of 850° C. for a time of 1000 hours is considered to represent the conditions of use of the steel, and is generally used to qualify steels.

In the invention, the Fe₂Nb precipitates are predominant among the intergranular precipitates. As mentioned above, these precipitates have the advantage of capturing part of the silicon to reduce the content thereof in the solid solution, the silicon then being less likely to segregate at the metal-oxide interface.

In particular, after said heat treatment, the volume fraction of Fe₃Nb₃X precipitates remains less than 0.05%.

In addition, together with grain size, the type and distribution of these Fe₂Nb precipitates are highly favourable for creep resistance up to 1000° C.

On account of the formation of the Fe₂Nb precipitates, segregation of silica at the metal-oxide interface is very limited compared with prior art steels.

In particular, after said ageing treatment in air at a temperature of 850° C. for 1000 hours, the surface fraction of silica segregation at the metal-oxide interface remains limited.

In particular, after said ageing treatment as illustrated in FIG. 1, the sheet on each of its surfaces comprises a layer of oxides 2 over the base metal 1. The layer of oxides 2 comprises an inner layer 4 of chromium oxide Cr₂O₃ (chromia) and an outer layer 5 of manganese-rich chromium oxide and possibly containing iron of spinel type (Mn,Fe)Cr₂O₄.

At the interface between the layer of oxides 2 and the steel of the sheet also called base metal 1, the sheet comprises precipitates of silicon oxide 3 or silica. In the sheets of the invention, after said ageing treatment, the surface fraction of the silica precipitates at this metal-oxide interface is less than or equal to 0.35, which means that the silica covers no more than 35% of the metal-oxide interface. The electrical conductivity of silica, in the region of 10⁻⁶ S·cm⁻¹ at 850° C., is 10 000 times weaker than that of chromia of approximately 10⁻² S·cm⁻¹ at 850° C., and its coefficient of thermal expansion is 10 times smaller than those of chromia, of the base metal and zirconia forming the electrolyte of the electrochemical cell. This surface fraction of silica at the metal-oxide interface is directly related to resistivity, to area specific resistance and to conductivity of the metal-oxide interface. Therefore, when the surface fraction of silica at the interface is 35%, the electrical conductivity of the interface is reduced by 35% compared with a silica-free configuration, changing from 10⁻² S·cm⁻¹ in the absence of silica to 6.5×10⁻³ S·cm⁻¹, or the resistivity of the interface comes to be increased by about 50%, increasing from 100 Ω·cm to 150 Ω·cm. Also, the area specific resistance of the interface, equal to resistivity multiplied by the thickness of the silica film, is increased in equivalent manner.

The surface fraction of the segregation of silicon oxides 3 can be determined from an image of a cross-section of a test specimen, in orthogonal direction to the direction of rolling, obtained under scanning electron microscopy with magnification of ×10 000 in backscattered electron mode.

The horizontal edges of the image are parallel to the surface of the sheet. The length L of the observed cross-section, in a plane transverse to the sheet, is 12 μm.

The backscattered electron mode generates a contrast of chemical composition on a scale of 256 levels, called greyscale, ranging from white (255) to black (0). In the images obtained, the silica on account of its composition and atomic number of the constituent elements thereof, appears as the darkest phase in contrast with the base metal which appears in very light shade and chromium oxide of intermediate grey shade.

The image thus obtained corresponds to the image denoted (a) in FIG. 2.

This image is analysed using image analysis software, for example using Image J software.

More particularly, the image initially transformed by the image analysis software via automated processing to increase the contrast between the elements of the image, shows: the base metal 1, layer of oxides 2 and segregations of silicon oxides 3. The objective of this step is that all the acquired images should have the same contrast and can evidence the segregations of silicon oxides 3 independently of the observed metal-oxide interface and analysed oxidized metal. After this treatment, an image is obtained denoted (b) in FIG. 2.

The image is then thresholded to distinguish between the segregations of silicon oxides 3 and the remainder of the image. Therefore, only two levels are maintained: black for the segregations of silicon oxides 3 and white for the remainder of the image, setting a threshold adapted for distinguishing of the segregations of silicon oxides 3. For example, the threshold is chosen to be 70, the pixels of the greyscale higher than 70 being represented in white 255 and the pixels of the greyscale lower than or equal to 70 being represented in black. Thereafter, any artefacts are filtered and the processing quality of each image is verified manually.

From this image, the sum is determined of the projected lengths Li, on a longitudinal axis, of the regions of the interface in which the precipitates of silicon oxides 3 are contained in this field of measurement, as illustrated in image (c) in FIG. 2.

The surface fraction is then calculated as the ratio between the sum of the projected lengths and the length L of the measurement field Σ Li/L. In the example illustrated in FIG. 2, the surface fraction is 77%.

In the invention, after heat treatment in air at 850° C. for 1000 hours, this ratio Σ Li/L remains less than or equal to 0.35, i.e. 35%.

Also, having regard to the chromium content of the steel of the invention, the total thickness of the layer of oxides 2 generally remains less than or equal to 10 μm after said heat treatment.

The sheet of the invention can be obtained in particular with the following method:

• • a steel is prepared having the preceding composition;
  • a semi-processed product is cast from this steel;
  • the semi-processed product is brought to a temperature higher than or equal to 1150° C. and lower than or equal to 1260° C. for a time of between 40 minutes and 60 minutes, and the semi-processed product is hot rolled to obtain a hot rolled sheet of thickness between 2.5 mm and 6 mm;
  • the hot rolled sheet is annealed for example at a temperature of between 1000° C. and 1100° C. for a time of 30 seconds to 6 minutes;
  • the hot rolled, annealed sheet is pickled;
  • said hot rolled sheet is cold rolled, at a temperature of between ambient temperature and 300° C., in a single step or in several steps, the sheet being annealed and pickled after each step. It is to be understood that by the term «step», it is meant herein cold rolling comprising either a single pass or a succession of several passes (e.g. five passes) which are not separated by any intermediate annealing; for example, a sequence of cold rolling can be envisaged comprising a first series of five passes followed by an intermediate annealing, then a second sequence of five passes; typically the intermediate annealing separating the steps is conducted at between 950° C. and 1100° C. for 30 seconds to 6 minutes;
  • final annealing is performed on the cold rolled sheet at a temperature of between 1000° C. and 1100° C., preferably between 1050° C. and 1090° C., and for a time of between 10 seconds and 6 minutes, to obtain a fully recrystallized structure with an ASTM average grain size preferably of between 4 and 7 if the sheet is of thickness between 1.2 mm and 2.5 mm and preferably between 4 and 9 if the sheet is of thickness greater than or equal to 0.1 mm and less than 1.2 mm. This heat treatment allows the placing of niobium in solid solution.

The volume fraction of Laves phases i.e. Fe₂Nb compounds of hexagonal structure, in the structure of the sheet, is very low and less than 0.2% in delivery condition i.e. after this final annealing.

A series of experiments will now be described showing the advantage of the invention. Laboratory castings were examined having the chemical analyses given in Table 1.

TABLE 1
Steel	C	Mn	Si	S	P	Cr	Ni	Mo	N	Cu	Nb	Ti	Zr	Al	V
I#1	0.009	0.34	0.16	0.001	0.002	21.9	0.20	0.002	0.018	0.10	0.69	0.120	0.002	0.009	0.11
I#2	0.015	0.36	0.14	0.001	0.023	21.3	0.24	0.026	0.010	0.07	0.54	0.089	0.002	0.008	0.12
I#3	0.014	0.34	0.16	0.001	0.002	19.1	0.20	0.002	0.015	0.05	0.69	0.130	0.002	0.009	0.13
K#1	0.018	0.34	0.15	0.001	0.004	19.2	0.20	1.920	0.015	0.05	0.70	0.140	0.002	0.009	0.12
K#2	0.017	0.34	0.39	0.001	0.004	19.0	0.20	1.910	0.015	0.10	0.59	0.001	0.001	0.005	0.14
K#3	0.017	0.29	0.45	0.001	0.003	21.0	0.20	0.020	0.023	0.40	0.39	0.130	0.001	0.010	0.10
K#4	0.016	0.27	0.26	0.0003	0.028	20.2	0.19	0.006	0.020	0.42	0.42	0.089	0.001	0.007	0.12
K#5	0.015	0.23	0/18	0.0003	0.03	17.0	0.20	0.006	0.020	0.31	0.30	0.080	0.001	0.007	0.11
					Ce +	Mn/	V +	Ti + V +	Ni +
Steel	Co	Sn	Ca	0	La + Y	Si	Zr + Al	Al + Zr	Cu + Co	2Nb − 7C	Ti − 4N
I#1	0.03	0.006	2	21	0	2.13	0.12	0.24	0.33	1.32	0.05
I#2	0.03	0.005	1	25	0	2.57	0.13	0.21	0.33	0.98	0.05
I#3	0.02	0.005	2	19	0	2.13	0.14	0.27	0.27	1.28	0.07
K#1	0.02	0.005	2	18	0	2.27	0.13	0.27	0.27	1.27	0.08
K#2	0.03	0.005	3	19	0	0.87	0.15	0.15	0.33	1.06	−0.06
K#3	0.01	0.005	2	20	0	0.64	0.11	0.24	0.61	0.66	0.04
K#4	0.03	0.005	2	21	0	1.05	0.13	0.22	0.64	0.74	0.01
K#5	0.03	0.005	2	22	0	1.28	0.13	0.21	0.54	0.49	0.01

For each of the steels in Table 1, the remainder is iron and unavoidable processing impurities.

The cast test specimens were transformed with the following method:

• • the cast semi-processed products were brought to a temperature of 1220° C. for 40 minutes, and hot rolled to obtain sheet of thickness 5 mm;
  • the sheets were annealed at 1080° C. for 6 minutes and pickled;
  • the hot rolled sheets were cold rolled at ambient temperature to obtain sheet of thickness 1.5 mm;
  • final annealing was performed on the sheets at a temperature of 1080° C. for 4 minutes.

The volume fraction of Laves phases was determined, i.e. Fe₂Nb compounds of hexagonal structure, in the structure of each sheet in delivery condition i.e. after annealing of the cold rolled sheet. Table 2 under the column «Fe₂Nb Laves phases in delivery condition» gives the volume fraction of Laves phases determined in the sheet in delivery condition.

The sheets were then subjected to heat treatment at 850° C. for a time of 1000 hours.

After this heat treatment, the precipitates in the structure were determined. Table 2 below under the column «Fe₂Nb Lave phases after treatment at 850° C. for 1000 hours», gives the volume fraction of Laves phases after this heat treatment.

Also determined was the volume fraction of Fe₃Nb₃X intermetallic precipitates in the sheet after this heat treatment; this fraction is given under the column «Fe₃Nb₃X cubic phases after treatment at 850° C. for 1000 hours» in Table 2.

For each of the sheets, the surface fraction was then determined of the segregation of silicon oxides 3 from an image obtained under electron microscopy as described above. The determined surface fraction is given under the column «Silica surface fraction after treatment at 850° C. for 1000 hours» in Table 2.

Also determined was the toughness of the annealed and pickled hot strip, obtained after hot rolling at 1220° C. to a thickness of 5 mm, with a Charpy impact test on a notched specimen following standard NF EN ISO 148-1 (March 2017 version): for this test, the energy absorbed by the test specimen is measured when impacted to breakage point with a 3-point flexure test, as a function of temperature (between −10° C. and 80° C.). It is considered that the strip is ductile if it has an impact strength higher than 30 J/cm² at the temperature of 20° C. (the closest temperature to ambient temperature).

TABLE 2
		Volume	Volume fraction	Volume fraction
		fraction of	of Fe2Nb	of Fe3Nb3X	Silica surface
		Fe2Nb Laves	Laves phases	cubic phases	fraction after
		phases in	after treatment	after treatment	treatment	Charpy
		delivery	at 850° C.	at 850° C.	at 850° C. for	toughness
		condition	for 1000 hours	for 1000 hours	1000 hours	of hot strip
Test	Steel	(%)	(%)	(%)	(%)	(J/cm2)
E1	I#1	0.06	1.5	0	7.6	35
E2	I#2	0.04	1.0	0	8.3	40
E3	I#3	0.05	1.25	0	19.4	50
E4	K#1	0.04	1.7	0	16.0	20
E5	K#2	0.02	0.7	0.10	74.1	10
E6	K#3	0.08	0.5	0	76.7	35
E7	K#4	0.10	0.6	0	39.0	40
E8	K#5	0.09	0.4	0	43.5	55

In Table 2 the comparative tests are underlined.

It is seen that, for the tests according to the invention (E1 to E3), in which:

• • the contents of niobium and titanium heed the above-described conditions, namely:
  • Nb between 0.40% and 1.0% and 2×Nb−7×C≥0.8%, and
  • Ti between 0.05% and 0.2%, and 0%≤Ti−4×N≤0.15%, and
  • the ratio Mn/Si is higher than 1:2, and
  • the sum Ni+Cu+Co is between 0 and 0.60%, the volume fraction of Fe₂Nb intermetallic precipitates of hexagonal structure is higher than or equal to 0.8% after treatment at 850° C. for 1000 hours, and the surface fraction of silicon oxide 3 after treatment at 850° C. for 1000 hours represents a surface fraction less than or equal to 35%.

In this manner, sheet is obtained having good performance in terms of electrical conductivity under conditions of use, and in particular comparable with that of steels with very high chromium content.

On the contrary, in the comparative tests E5 to E8, the contents of titanium and/or niobium do not heed the above conditions. In these tests, the volume fraction of Fe₂Nb intermetallic precipitates of hexagonal structure is less than 0.8% after treatment at 850° C. for 1000 hours. The silicon is therefore less likely to be trapped by the Fe₂Nb Laves phases and hence limit segregation of silica even if the silicon content is less than 0.2% as in the K5 alloy. In addition, in these comparative tests, the Mn/Si ratio does not heed the above condition. It will be noted that; in these tests, the surface fraction of silicon oxide 3 after treatment at 850° C. for 1000 hours represents a surface fraction greater than 35%, which leads to degraded performance in terms of electrical conductivity, with interface resistivity increased by more than 50%.

In addition, in tests E1 to E3, the hot strip is ductile since it exhibits an impact strength greater than 30 J/cm². On the contrary, in the comparative tests E4 and E5 in which the steel has a molybdenum content higher than the limits described in the invention, the hot strip obtained is not ductile since it has an impact strength of less than 30 J/cm².

As explained in the foregoing, according to one optional aspect:

• • the average grain size of the steel is between 30 micrometres and 80 micrometres, i.e. an ASTM number of between 4 and 7 when the thickness of the cold rolled and annealed sheet is between 1.2 mm et 2.5 mm; and
  • the average grain size of the steel is between 15 micrometres and 80 micrometres, i.e. an ASTM number of between 4 and 9 when the thickness of the cold rolled and annealed sheet is greater than or equal to 0.1 mm and less than 1.2 mm.

To confirm the technical effect of this optional property, the inventors subjected sheets having compositions of the invention to a creep test under their own weight at 850° C. for 200 hours.

These sheets were obtained from hot rolled and pickled sheet, prepared with the above-described method, that was cold rolled at ambient temperature to a final thickness given in Table 3 for each test, followed by final annealing under annealing conditions given in Table 3. The sheets were then pickled and the grain size measured with the circular intercept procedure described in standard ASTM E112. Table 3 also specifies the average grain size for each test.

Creep was measured with a creep test known as the «Sag Test». The Sag Test is not normalized but is used to characterize creep. This test is described in the article by Faria, Geraldo Lucio de; Melo, Denilson Pereira de; Moreira, Paulo Sérgio. UTILIZAçÃO DA METODOLOGIA SAG TEST PARA AVALIAR O COMPORTAMENTO EM FLUÊNCIA DOS AçOS INOXIDÁVEIS AISI 321 E AISI 441, p. 34-44. In: 75° Congresso Anual da ABM, São Paulo, 2022. ISSN: 2594-5327, DOI 10.5151/2594-5327-34135.

For the creep test, strips of metal are cut from sheet in delivery condition (cold rolled, annealed and pickled), the strips having a length of 205 mm and cut in the direction of rolling, width 25 mm and thickness corresponding to the final thickness of the strip (1.5 mm or 0.5 mm depending on the test under consideration).

These planar strips were suspended on two supports with an air gap of 200 mm, in a furnace at 850° C. for a determined time. Measurements of sagging, characterizing deformation under creep, were regularly carried out at 1 hour, 25 hours, 50 hours, 100 hours and 200 hours. Measurement of sagging of the test specimen was performed at ambient temperature on a planar surface, typically marble, using a comparator having accuracy and resolution of less than 0.05 mm. For each test, three specimens were tested.

Table 3 below gives the mean sag values of the test specimen after an exposure time of 200 hours.

TABLE 3
		Final					Mean sag
		thickness					measured after
		after cold	Annealing	Annealing	Average	Average	exposure time
		rolling	temperature	time	grain size	grain size	of 200 hours
Test	Steel	(mm)	(° C.)	(min)	(μm)	(ASTM)	(mm)
E9	I#1	1.5	1080	4	40	6.3	0.9
E10			1030	4	11	10	13.1
E11	I#2		1080	4	60	5.1	0.4
E12			1030	4	20	8.2	3.8
E11	I#2	0.5	1080	3	107	4	4.0
E12	I#3		1030	3	10	10.2	>50
E13			1050	3	36	6.5	8.4

These tests show that a grain size of less than 30 μm for a sheet thickness of 1.5 mm degrades creep properties. Sagging of more than 3 mm was observed after 200 hours at 850° C. Similarly, an average grain size of less than 15 μm for a sheet thickness of 0.5 mm degrades creep properties. Sagging of more than 9 mm was observed after 200 hours at 850° C.

The embodiment in which the average grain size heeds the conditions specified above is therefore particularly advantageous in terms of creep resistance.

Claims

1. Ferritic stainless steel sheet, the composition of which comprises, C ≤ 0.03 % 0.25 % ≤ Mn ≤ 1 ⁢ %, 0 ⁢ % ≤ Si ≤ 0.2 %, with ⁢ Mn / Si ≥ 1.2 S ≤ 0.005 % P ≤ 0.04 % 19. % ≤ Cr ≤ 24. % Ni ≤ 0.5 % Mo ≤ 0.1 % N ≤ 0.03 % Cu ≤ 0.2 % 0.4 % ≤ Nb ≤ 1. % 0.05 % ≤ Ti ≤ 0.2 %, Zr ≤ 0.02 % Al ≤ 0.02 % V ≤ 0.2 % Co ≤ 0.05 % Sn ≤ 0.05 %, Rare ⁢ earth ≤ 800 ⁢ ppm V + Zr + Al ≤ 0.2 % Ti + V + Zr + Al ≤ 0.3 % Ti + Nb ≤ 1. % Ni + Cu + Co ≤ 0.6 % 2 × Nb - 7 × C ≥ 0.8 % 0 ⁢ % ≤ Ti - 4 × N ≤ 0.15 % 0.2 ppm ≤ Ca ≤ 20 ⁢ ppm 1 ⁢ ppm ≤ O ≤ 60 ⁢ ppm the remainder of the composition consisting of iron and unavoidable processing impurities, the sheet being an annealed and pickled sheet, the sheet comprising a volume fraction of Fe2Nb Laves phases of less than 0.2%.

the contents being expressed by weight:

it being understood that:

2. The ferritic stainless steel sheet according to claim 1, wherein: 50 ⁢ ppm ≤ Rare ⁢ earths ≤ 800 ⁢ ppm.

3. The ferritic stainless steel sheet according to claim 1, wherein: Nb - 10 × ( C + N ) ≥ 0 ⁢ %.

4. The ferritic stainless steel sheet according to claim 1, wherein the sheet has a thickness of between 0.1 mm and 2.5 mm.

5. The ferritic stainless steel sheet according to claim 4, wherein the sheet has an average grain size of between 30 micrometres and 80 micrometres when the sheet has a thickness of between 1.2 mm and 2.5 mm, and an average grain size of between 15 micrometres and 80 micrometres when the sheet has a thickness greater than or equal to 0.1 mm and less than 1.2 mm.

6. The ferritic stainless steel sheet according to claim 1, wherein the sheet is cold rolled and annealed sheet.

7. The ferritic stainless steel sheet according claim 1, which, when subjected to heat treatment at a temperature of 850° C. for a time of 1000 hours, comprises a volume fraction of Fe2Nb Laves phases higher than or equal to 0.8%.

8. The ferritic stainless steel sheet according to claim 1, which, when subjected to heat treatment at 850° C. for a time of 1000 hours, comprises a volume fraction of Fe3Nb3X phases of less than 0.05%.

9. The ferritic stainless steel sheet according to claim 1, wherein, when subjected to heat treatment at a temperature of 850° C. for a time of 1000 hours, the sheet on each of its surfaces comprises a layer of oxides (2) and, at the interface between the sheet steel and the layer of oxides (2), comprises precipitates of silicon oxide (3) such that the surface fraction of the precipitates of silicon oxide (3) at the interface between the sheet steel and layer of oxides (2) is less than or equal to 0.35.

10. The ferritic stainless steel sheet according to claim 9, wherein the layer of oxides (2) has a thickness less than or equal to 10 μm.

11. A method to produce ferritic stainless steel sheet, characterized in that:

a steel is prepared having the composition according to claim 1;

a semi-processed product is cast from this steel;

the semi-processed product is brought to a temperature higher than or equal to 1150° C. and lower than or equal to 1260° C. for a time of between 40 minutes and 60 minutes, and the semi-processed product is hot rolled to obtain hot rolled sheet of thickness between 2.5 mm and 6 mm;

the hot rolled sheet is annealed;

the hot rolled and annealed sheet is pickled;

said hot rolled sheet is cold rolled at a temperature of between ambient temperature and 300° C., in a single step or in several steps separated by intermediate annealings;

final annealing of the cold rolled sheet is performed at a temperature of between 1000° C. and 1100° C. for a time of between 10 seconds and 6 minutes, to obtain a fully recrystallized structure.

12. The method according to claim 11, characterized in that the annealing of the hot rolled sheet is conducted at a temperature of between 1000° C. and 1100° C., for a time of 30 seconds to 6 minutes.

13. The method according to claim 11, characterized in that the intermediate annealing operation(s) are conducted at a temperature of between 950° C. and 1100° C. for a time of 30 seconds to 6 minutes.

14. The method according to claim 11, characterized in that the final annealing is conducted at a temperature of between 1050° C. and 1090° C.