Austenitic stainless steel with aluminum oxide formation and its uses

Abstract

An austenitic stainless steel with medium Manganese content and aluminum oxide formation is provided which is used in the nuclear field, preferably in a nuclear plant and / or in a process for the production of nuclear energy, more preferably in a fourth-generation nuclear plant, wherein the reactor is lead-cooled.

Description

[0001] AUSTENITIC STAINLESS STEEL WITH ALUMINUM OXIDE FORMATION

[0002] AND ITS USES

[0003] Cross-Reference to Related Patent Applications

[0004] This patent application claims priority of the Italian patent application No . 102024000014233 filed on June 20 , 2024 , the content of which is incorporated by reference herein .

[0005] Field of the Invention

[0006] The present invention refers to the use of an austenitic stainless steel with aluminum oxide formation in the nuclear field at high temperatures and high neutron irradiation .

[0007] Description of the Prior Art

[0008] Technological progress in fourth-generation nuclear power poses unprecedented engineering challenges , especially with regard to the energy ef ficiency and operational safety of lead-cooled fast reactors ( LFRs ) .

[0009] Lead-cooled reactors , as well as associated apparatuses and equipment , are designed to operate at signi ficantly elevated temperatures , above 450 ° C, in order to maximi ze the energy ef ficiency of the nuclear system . Such an operating condition requires the use of new structural materials endowed with high metallurgical stability in the critical temperature range of 500-700 ° C, while ensuring mechanical resistance , corrosion and prolonged oxidation over time , at the most extreme operating conditions .

[0010] In the energy field, as well as in the chemical and petrochemical field, where operating temperatures are between 600 ° C and 900 ° C, a new type of austenitic stainless steel with aluminum oxide formation ( alumina- forming austenitic, AFA, stainless steel ) has been developed .

[0011] As described by M . P . Brady et al . in "Co-optimi zation of wrought alumina-forming austenitic stainless steel composition ranges for high-temperature creep and oxidation / corrosion resistance", Material Science & Engineering A, 590 (2014) , 101-115, and in US 8,431,072, the main characteristic of such steels is to form, on the free surface of the component, a continuous layer of Aluminum oxide (AI2O3) starting from an austenitic microstructure free of delta phase and bcc phase. The range of application in temperature starts from 550 °C onwards.

[0012] In this regard, WO22211709 describes the use of such a family of austenitic stainless steels with aluminum oxide formation in lead-cooled nuclear reactors. However, these austenitic stainless steels are either particularly rich in Nickel (in the range 20-32 wt%) , thus being very expensive and particularly subject to activation under neutron irradiation or, in the configuration with the lowest Nickel content (12%) , have reinforcement mechanisms through precipitation based on complex metal carbides (e.g. M23C6) , notoriously subject to microstructure embrittlement due to their dimensional growth over time and with increasing temperature (also called coarsening) , i.e. more subject to intergranular corrosion in a mixed and welded zone, due to the phenomenon of sensitization induced by high Carbon contents (as in this case) . Furthermore, such steel is simply used as a contact coating with liquid lead or eutectic leadbismuth mixtures and therefore not for structural applications. It should also be noted that in WO22211709 the presence of a minimum fraction (5%) of delta / bcc phase is a required condition.

[0013] In parallel, austenitic stainless steels with aluminum oxide formation, but with a high Manganese content, as described by Y. Yamamoto in "Evaluation of Mn substitution for Ni in alumina-forming austenitic stainless steels", Material Science and Engineering A 524 (2009) 176-185, and in US 2008 / 0292489 have been developed. These, unlike the parallel austenitic stainless steel with aluminum oxide formation, are less expensive since Manganese replaces Nickel, while maintaining good mechanical, creep and corrosion / oxidation properties.

[0014] The rather high content of Manganese in the product analysis (up to 15 wt% or higher, hence high-Mn AFA) , i.e. the combination of Manganese with particularly high contents of Chromium (up to 17 wt%) , on the other hand, leads to the formation of M23C6, which are notoriously subject to embrittlement of the microstructure due to their dimensional growth over time and with increasing temperature (also called coarsening) and capable of significantly reducing the creep resistance of stainless steel, even for not particularly high operating temperatures (also called low-temperature ) . In addition, the combination of Manganese and Chromium in high percentages under analysis can compromise the formation of the AI2O3 layer. Last but not least, higher and higher levels of Manganese stabilize the sigma phase, a strong weakener of the microstructure responsible for the rapid decrease in the mechanical and ductility properties of steel.

[0015] There is also the difficulty of industrially producing austenitic stainless steels with high Manganese content, due to the high emission of Manganese fumes in the steel plant (low-melting metal) , which can only be contrasted with the use of pressurised plants (with a considerable increase in manufacturing costs) , i.e. of fume extraction and abatement systems specifically sized for these steels, a condition (to our knowledge) unrelated to the vast majority of world steel production . To date , austenitic stainless steels with aluminum oxide formation and with a high or medium content of Manganese in the nuclear field are not known . There is therefore a need to develop new austenitic stainles s steels with aluminum oxide formation in order to enhance and commerciali ze the nuclear technology and especially the fourth-generation one .

[0016] Description of the Invention

[0017] Accordingly, an aim of the present invention consists in providing an austenitic stainless steel with aluminum oxide formation with structural and functional properties suitable for overcoming the above-described problems of the prior art .

[0018] In particular, an aim of the present invention cons ists in providing an austenitic stainless steel for use in the nuclear field, i . e . in an environment where operating temperatures are high (between 500 ° C and 700 ° C ) , where neutron irradiation is intense and where oxidation and corrosion phenomena are very common .

[0019] A speci fic aim of the invention consists also in providing an austenitic stainles s steel that is able to guarantee good mechanical properties thanks to a high metallurgical stability, a high resistance to creep, to damage induced by neutron irradiation and to oxidation and corrosion phenomena, and at the same time to be economical and manufacturable .

[0020] A further aim of the present invention consists in providing an austenitic stainless steel capable of ensuring the safety and ef ficiency of apparatuses , equipment or devices ( for example of a lead-cooled nuclear reactor ) in which it is employed in the nuclear field and, preferably, in the fourth-generation nuclear field . Accordingly, the present invention refers to an austenitic stainless steel with aluminum oxide formation for use in a nuclear plant and / or in a process for the production of nuclear energy, as defined in claim 1 .

[0021] The invention further refers to an apparatus , equipment or device , in particular of a nuclear plant or used in a process for the production of nuclear energy, comprising at least a part made of corros ion-resistant austenitic stainless steel , as defined in claim 15 .

[0022] The invention further refers to a nuclear plant and a process for the production of nuclear energy comprising at least one such apparatus , equipment or device having at least a part made of an austenitic stainless steel as defined, respectively, in claims 17 and 19 ; and to an austenitic stainless steel with medium Manganese content and aluminum oxide formation in a lead-cooled nuclear reactor as defined in claim 21 .

[0023] Preferred and advantageous features of the invention form the subj ect-matter of the secondary claims .

[0024] The austenitic stainless steels of the invention are speci fically characteri zed by a chemical composition and a microstructure such as to give them the mechanical and chemical-physical properties necessary to be used in the nuclear field and, preferably, in fourth-generation nuclear technology . In detail , said austenitic stainless steel combines a high content of Manganese with speci fic amounts of Cr, Ni and Al : in fact , it has been acknowledged that these elements , used together according to speci fic composition rules , have a very speci fic and predictable combined ef fect on metallurgical stability, durability, structural reliability, and corros ion resistance as well as on other favourable material properties . In fact , it has been found that these main elements (Mn, Ni , Al , Cr ) , combined also with other secondary elements , ef fectively increase the mechanical and physicalchemical properties of an austenitic stainless steel with aluminum oxide formation (having the particular composition of the invention) i f each element is used in a speci fic content field and the Manganese content is increased with respect to the known art in the nuclear field .

[0025] In other words , the inventors have found that the austenitic stainless steels having the particular compositions of the invention, and in particular, having a higher Manganese content than the known austenitic steels in the nuclear field, simultaneously show :

[0026] - excellent metallurgical stability within a range of operating temperatures between 500 ° C and 700 ° C, maintaining this stability for extended periods of time (up to 40 years ) . The continuous exposure to these temperatures can in fact induce precipitation of brittle phases within the steel matrix, i . e . nucleation and growth of ferrite and sigma phase in the prevailing austenitic microstructure , critically compromising the mechanical properties , such as creep resistance , and technological properties , such as corrosion resistance . Therefore , reducing such deterioration allows achieving a good structural integrity and, therefore , a good and lasting functionality of apparatuses , equipment or devices in the nuclear field, for example of the reactor, comprising said steel . Replacement and / or maintenance intervals are therefore minimi zed for the benefit of a greater reactor productivity and lower operating costs .

[0027] - excellent resistance to neutron radiation when such steels are used in the nuclear field . In the context of , for example , fourth-generation nuclear reactors , the materials used, for example steels , are exposed to high intensity neutron fluxes . This interaction between the neutrons and the atomic structure of the materials causes signi ficant alterations within their crystal lattice , mani festing themselves through the introduction of point defects , the generation of voids and the accumulation of dislocation rings . These phenomena alter the macroscopic properties of the materials ( such as , for example , ductility and toughness ) , inducing a process of embrittlement and a volumetric increase of the components ( swelling) . Reducing such damages therefore leads to an improved durability and structural reliability of apparatus , equipment or device in which steel is used .

[0028] - a high resistance to corrosion by contact , for example , with molten lead . The interaction between molten lead and the structural materials , such as the known advanced steels and the nickel-based alloys , can accelerate the corrosion processes , leading to the dissolution of key elements of the alloy, to the degradation of the exposed surface and, ultimately, to the deterioration of the mechanical properties of the material . To mitigate corrosion in molten lead, the present austenitic stainless steel exhibits a reduced sensitivity to liquid metal embrittlement (LME ) mechanisms and has an intrinsic ability to form protective barriers at the interface , with high sel f-healing capabilities .

[0029] Thermodynamic and kinetic simulations and experimental tests confirm that the austenitic stainless steels according to the invention, i . e . having a speci fic combined content of Mn, Ni , Al and Cr and other elements lead to the formation of an austenitic metastable matrix with low Stacking Fault Energy ( SEE ) and devoid, for kinetic reasons , of sigma phase and ferrite / bcc in the temperature range between 500 ° C and 700 ° C, able to resist corrosion by molten lead and capable of producing only very fine carbides of Niobium / Titanium and not of complex Chromium .

[0030] This result cannot be predicted in consideration of the teachings of the prior art .

[0031] On the contrary, the inventors of the present invention have understood that a certain amount of Manganese , higher than the austenitic steels that are customary in the nuclear field, has in fact a good impact on the resistance to corrosion, to the damage induced by neutron irradiation and to the microstructural deformations of an austenitic stainless steel , i f Manganese is associated with Chromium, Nickel and Aluminum and other elements according to speci fic rules .

[0032] In short , the austenitic stainless steels with aluminum oxide formation in the nuclear field of the invention have a Manganese content of at least 5 . 0% (herein and in the following, all percentages are understood, unless otherwise speci fied, as percentages by weight with respect to the total weight of the steel ) .

[0033] Presentation of the Figures

[0034] Further characteristics , obj ects and advantages of the invention will emerge from the fol lowing description, which is purely illustrative and non-limiting and supported by the attached figures , in which :

[0035] Figure 1 shows the equilibrium phases in the temperature range between 550 ° C and 650 ° C calculated for some austenitic steels according to the present invention ( Figure la, lb, 1c, Id, 1g) and for two comparative austenitic steels ( Figure le and I f ) . - Figure 2 shows the molar percentage amount of carbides of type M23C6 ( in dots ) and carbides of Niobium or Titanium ( grey) at equilibrium at 600 ° C for some steels according to the invention and for three comparative steels ;

[0036] Figure 3 shows the TTT ( Time Temperature Trans formation) diagram for the sigma phase in di fferent molar fraction percentages in di f ferent austenitic stainless steels according to the present invention;

[0037] - Figure 4 shows the thickness of the surface layer comprising aluminum oxides ( figure 4a ) and chromium oxides ( figure 4b ) in some variants according to the present invention;

[0038] - Figure 5 shows the precipitation kinetics of two phases : of the sigma phase in a first comparative steel with respect to a steel according to the present invention ( Fig . 5a ) ; of ferrite in a second comparative steel with respect to a steel according to the present invention ( Fig . 5b ) ;

[0039] Figure 6 shows the free energy of formation of aluminum oxide (AGf ) and of the actual concentration of valence electrons (Veff ) for the austenitic stainless steels according to the present invention in relation to the oxidation rate ( kt) ;

[0040] Figure 7 shows the calculated SFE ( acronym for Stacking Fault Energy) value of a steel according to the present invention and of the austenitic matrix after heat treatment at 600 ° C ;

[0041] - Figure 8 shows results of air aging tests after 2160 hours at 650 ° C obtained by X-ray di f fraction (XRD) for three austenitic steels according to the present invention;

[0042] Figure 9 shows results of the lead corrosion resistance tests obtained by SEM-EDS analysis for an austenitic steel according to the present invention; Figure 10 shows results of the lead corrosion resistance tests obtained by SEM-EDS analysis for a further austenitic steel according to the present invention; and

[0043] Figure 11 shows results of the lead corrosion resistance tests obtained by SEM-EDS analysis for a subsequent austenitic steel according to the present invention .

[0044] Detailed Description

[0045] Austenitic stainless steels with aluminum oxide formation, preferably used in the nuclear field, for example in a nuclear power plant and / or in a process for the production of nuclear energy, have a Manganese content of at least 5.00%, preferably between 5.00% and 14.00 wt%, more preferably between 5.50% and 13.0 wt%, even more preferably between 6.00% and 12.00 wt% (inclusive) .

[0046] Manganese (Mn) is in fact an element that forms austenite and a certain amount of Manganese is required to ensure the stabilisation of the austenitic phase. In addition, Manganese strongly contributes to reducing stacking fault to promote the twinning deformation mechanism.

[0047] On the other hand, Manganese promotes the formation of sigma brittle phases that negatively impact the mechanical properties (e.g. in relation to the creep phenomena) of apparatuses, equipment or devices and comprising said steel. In fact, this sigma phase tends to coalesce at the triple grain edges, representing the cause of triggers of microcracks .

[0048] According to the invention, Manganese is thus used in combination with specific amounts of other elements, including Niobium (Nb) , Titanium (Ti) and Carbon (C) , with the aim of maximising the formation of nano-precipitates of secondary carbides of Niobium and Titanium, to the detriment of the complex chromium carbides of the M23C6 type. In this regard, the addition of Copper also leads to the formation of nano-precipitates with action complementary to the previous ones.

[0049] Niobium, in fact, is present in the stainless steel according to the invention in an amount equal to or less than 1.00%, preferably between 0.00 and 0.60 wt% . In addition to minimizing sigma phase formation, the formation of Niobium nano-precipitates increases creep resistance and hinders the movement of the defects under irradiation.

[0050] As for Niobium, also the content of Titanium in the steel of the present invention is equal to or less than 1.00%, preferably between 0.10% and 0.80%, more preferably between 0.15% and 0.60 wt%. Titanium in these amounts plays a crucial role as a stabilizer: by binding with Carbon to form Titanium carbides, it prevents the precipitation of Chromium carbide at the edges of the grain during the welding process and heating to high temperatures, and therefore contributes to maintaining the resistance to intergranular corrosion of steel unchanged.

[0051] While Carbon is present in an amount equal to or lower than 0.15%, preferably between 0.05 and 0.14 wt%, more preferably between 0.10% and 0.14 wt% . In these amounts, Carbon contributes to increasing the mechanical resistance of steel through the formation of carbides that strengthen the austenitic matrix.

[0052] The present austenitic steel further comprises Chromium in amounts between 10.50% and 15.00%, preferably between 10.70% and 14.00 wt%, more preferably between 11.00% and 13.50 wt%. In these amounts, Chromium is advantageous for ensuring corrosion resistance, e.g. the corrosion induced by contact with molten lead when the present steel is used in the fourth-generation nuclear field, wherein the reactor is cooled with molten lead. This corrosion resistance is achieved through Aluminum oxide formation on the free surface of an apparatus, equipment or device made with the present steel. In fact, the Aluminum oxide formation is linked not only to the amount of Aluminum, but also to the amount of a third element, such as precisely Chromium (it is known as Third Element Effect) . On the other hand, Chromium is a ferritizer and contributes to the creation of the sigma phase. However, the combination of Chromium with specific amounts of other elements, such as Ti, Nb and C, ensures the formation of an austenitic matrix free of sigma brittle phase .

[0053] In addition to Chromium, to achieve optimal corrosion resistance, for example to molten lead, specific amounts of other elements are added, including Yttrium (Y) , Zirconium (Zr) , Lanthanum (La) and Scandium (Sc) . These elements, by binding to the oxygen present in the steel even at very low concentrations, work as a getter, that is, these elements are strongly related to oxygen, thus being able to oxidize quickly by sequestering oxygen and delaying the establishment of the internal oxidation regime. This allows the atomic aluminum to diffuse towards the free surface to form the protective oxide layer.

[0054] Therefore, these reactive elements, such as Zirconium, Yttrium, Lanthanum, and Scandium are present in the stainless steel according to the invention each in an amount equal to or less than 1.00%, preferably between 0.00 and 0.50 wt% of Zirconium and / or Yttrium and / or Scandium, more preferably 0.00% and 0.30 wt% of Zirconium and Yttrium and / or Scandium, and preferably between 0.00% and 0.60 wt% of Lanthanum. These elements promote precisely the formation of a compact oxide layer and, in addition, act as oxygen getters by delaying internal oxidation. Yttrium and Zirconium, for example, are elements that oxidize very easily compared to all other metals present in the composition of the present austenitic stainless steel and extend their effect to the highest operating temperatures.

[0055] In one embodiment of the present invention, Yttrium is not present in the austenitic stainless steel. Yttrium, in fact, can contribute, together with Nickel and Iron, to forming an unfavourable eutectic phase. This eutectic phase may adversely affect the subsequent thermomechanical processes due to its low melting point. Furthermore, austenitic stainless steels according to the invention, comprising Yttrium, can increase B2-BCC phase precipitation.

[0056] In a further embodiment of the invention, the austenitic stainless steel comprises Zirconium and / or Scandium. Preferably, the present austenitic stainless steel contains Zirconium and / or Scandium in a range between 0.04 and 0.50%, preferably between 0.05 and 0.30 wt%, more preferably 0.05 wt% of Zirconium and 0.10 wt% of Scandium. The presence of Scandium and / or Zirconium in the austenitic stainless steel according to the invention avoids the formation of a eutectic phase, limits the B2-bcc phase precipitation and favours the formation of carbides with fee structure. Furthermore, austenitic stainless steels according to the invention comprising Scandium and / or Zirconium may show increased resistance to lead corrosion.

[0057] The present austenitic stainless steel further contains Nickel (Ni) in a range between 9.00 and 19.00%, preferably 10.00 and 17.00 wt%, preferably between 10.50 and 16.50 wt%, more preferably between 11.00 and 16.30 wt%, still more preferably between 11.50 and 16.00 wt%. Like Manganese, Nickel also plays a fundamental role in stabilising the austenitic phase. At the same time, Nickel must be limited and optimized to ensure a low SFE of the matrix and avoid the formation of brittle phases.

[0058] The Nickel content is limited to 19.00% to reduce the swelling phenomenon (i.e. less activation) under neutron irradiation and for reasons of cost and sustainability of the present steel . In addition, a high Nickel content would make the steel subject to a stronger dissolution phenomenon.

[0059] Therefore, Nickel is partially replaced by Manganese with advantages in terms of reduction of swelling, cost and diffusion in molten lead.

[0060] The present austenitic stainless steel also comprises Cobalt in an amount equal to or less than 5.00%, preferably in a range between 0.00% and 2.80%. Advantageously, Cobalt further increases the stability of the austenitic matrix and contributes to increasing the creep resistance. Like Manganese, Cobalt also consistently reduces the value of SFE (total and austenitic phase) .

[0061] It should be pointed out that the introduction of elements that lower the value of SFE, such as Manganese and Cobalt, in the present steel makes the formation of partial dislocations energetically favourable. Such partial dislocations promote the activation of metallurgical mechanisms that promote extensive steel ductility. In fact, for partial dislocations, one of the movement mechanisms (e.g. cross slip) is statistically unlikely and this makes such movement more difficult.

[0062] Low and intermediate SFE values also activate further plastic deformation mechanisms, such as Twinning-Induced Plasticity (TWIP) or Transformation-Induced Plasticity (TRIP) deformation.

[0063] In order for the addition of such elements with low SFE to have an effect on the mechanical properties of the present steel, the latter are part of the chemical elements that form the matrix consisting of a cubic solid solution with centred faces (fee, austenite) .

[0064] At the same time, these elements (Ni, Mn and Co) selectively inhibit other metallurgical mechanisms, such as the mobility of the reticular defects, so as to minimize the creep phenomenon.

[0065] Disadvantageously, however, Cobalt is subject to activation under irradiation. For this reason, its use in the nuclear field is subject to limitation and therefore cannot be higher than 5.00%.

[0066] The present austenitic stainless steel obviously contains Aluminum in an amount between 2.00% and 4.00%, preferably between 2.20% and 3.50 wt%, more preferably between 2.20% and 3.00 wt%. Aluminum in fact forms aluminum oxide on the free surface of an apparatus, equipment or device made with the present steel. Advantageously, the specific quantitative range avoids the formation of ferritic phase in the operating temperature range between 500 °C and 700 °C.

[0067] However, Aluminum disadvantageously has a high deformation energy density value. Therefore, the selected Aluminum content allows an optimal balance between the positive effect of AI2O3 formation and the reduced formation of partial dislocations.

[0068] The austenitic stainless steel of the invention further contains the following elements.

[0069] Tantalum (Ta) . Tantalum is responsible for the formation of stable carbides, such as Tantalum Carbide (TaC) , which contribute to the microstructural stability of steel at high temperatures. Tantalum can also improve the corrosion resistance of steel, especially in aggressive environments. The amount of Tantalum is equal to or less than 1.00%, preferably between 0.00% and 0.60 wt%.

[0070] Silicon (Si) . Silicon is used as a f errite-forming element and for the deoxidation phase in steelworks, i.e. in the manufacturing process of the austenitic stainless steels. High amounts of Silicon are avoided in order to reduce the possibility of precipitation of intermetallic phases. Thus, the amount of Silicon is equal to or less than 0.10%, preferably from 0.10% to 0.80 wt%, preferably between 0.15% and 0.60 wt%.

[0071] Copper (Cu) . Copper is an element capable of stabilising the austenitic matrix. In general, the addition of Copper leads to the formation of nano-precipitates with complementary action with respect to the nano-precipitates of secondary carbides of Niobium and Titanium. However, in too high concentrations, Copper can lead to the formation of brittle phases especially under irradiation. Therefore, the Copper content is equal to or less than 3.00%, preferably from 0.00% and 2.80 wt%.

[0072] The steel compositions of the invention may also comprise unavoidable impurities.

[0073] By impurities is meant the set of elements and compounds that are not intentionally added to the steel formulation but are nevertheless present in small amounts that are contained in the raw materials used for the manufacture of austenitic stainless steel.

[0074] Impurities may include phosphorus (P) and sulfur (S) , as well as extremely low levels of Antimony (Sb) and Tin (Sn) . However, the content of P and S should be kept as low as possible. Therefore, the S content should be less than

[0075] O.005%, and the P content should be less than 0.025%. Typical amounts are less than 0.005% for S and less than 0.020% for

[0076] P.

[0077] Therefore, the content of the impurities may be represented by a range not exceeding 0.025 wt%.

[0078] Summarizing, the austenitic stainless steel with medium Manganese content and aluminum oxide formation, mainly used in a nuclear plant, in particular in at least a part of an apparatus, equipment or device, in direct contact with a primary cooling fluid, circulating in a reactor, and / or in a process for the production of nuclear energy, comprises in weight percentage:

[0079] Mn from 5.00 to 14.00

[0080] C from 0.05 to 0.15, preferably from 0.05 to 0.14

[0081] Cr from 10.50 to 13.50, preferably from 10.70 to 13.50

[0082] Ni from 9.00 to 19.00, preferably from 10.00 to 17.00 Al from 2.00 to 4.00, preferably from 2.20 to 3.50 Si from 0.10 to 0.80

[0083] Ti from 0.10 to 0.80

[0084] Nb 1.00 or less, preferably from 0.00 to 0.60

[0085] Co 5.00 or less, preferably from 0.00 to 2.80

[0086] Cu 3.00 or less, preferably from 0.00 to 2.80

[0087] Zr and / or Sc from 0.00 to 0.50

[0088] La 1.00 or less, preferably from 0.00 to 0.60

[0089] Ta 1.00 or less, preferably from 0.00 to 0.60

[0090] Y 1.00 or less, preferably from 0.00 to 0.50; impurities 0.025 or less; the balance being Fe .

[0091] The invention thus provides improved formulations of austenitic stainless steels, completely suitable for use in highly corrosive conditions, i.e. for example in contact with molten lead, even at temperatures ranging from 500 ° C to 700 ° C .

[0092] The austenitic stainless steels with medium manganese content and aluminum oxide formation of the invention are suitable for use in any nuclear field, i . e . in any type of nuclear plant / process for the production of nuclear energy, and speci fically in a nuclear reactor, preferably in a lead- cooled nuclear reactor, operating at high temperatures (between 500 ° C and 700 ° C ) , in the presence of a neutron flux typical of a fast reactor, and an induced damage equal to or higher than 1 dpa ( dpa stands for displacement per atom) , and in contact with molten lead, used in a fourthgeneration nuclear plant and in the relative process for the production of nuclear energy .

[0093] Thus , the austenitic stainless steels of the invention are especially useful for the manufacture of apparatuses , equipment and devices ( or parts thereof ) , in direct contact with the primary cooling fluid of the reactor, circulating in the reactor, which are exposed to high temperatures , to an intended neutron flux and, in some cases , to molten lead such as the reactor and / or parts connected thereto .

[0094] The steels of the invention are therefore suitable for use in a nuclear plant of any type , including in particular the most demanding conditions in terms of corrosion of a fourth-generation nuclear plant comprising a lead-cooled reactor .

[0095] The invention therefore refers to the use of austenitic stainless steel as described herein in a nuclear plant , and speci fically in an apparatus , equipment or device ( or a part thereof ) , in direct contact with the primary cooling fluid of the reactor, circulating in the reactor, which is exposed to high temperatures , an intense neutron flux and, in some cases, molten lead.

[0096] The invention further refers to an apparatus, equipment or device, in particular of a nuclear plant or used in a process for the production of nuclear energy, comprising at least a part made of austenitic stainless steel described herein, in direct contact with the primary cooling fluid of the reactor, circulating in the reactor.

[0097] The invention further refers to a nuclear plant and a process for the production of nuclear energy comprising at least one apparatus, equipment or device having at least a part made of austenitic stainless steel as described herein.

[0098] Finally, the invention also refers to an austenitic stainless steel with medium Manganese content and aluminum oxide formation in a lead-cooled nuclear reactor, comprising in weight percentage (wt%) : Mn from 5.00 to 14.00, preferably from 5.50 to 13.00 C from 0.05 to 0.15, preferably from 0.05 to 0.14 Cr from 10.50 to 13.50, preferably from 10.70 to 13.50 Ni from 9.00 to 19.00, preferably from 10.00 to 17.00 Al from 2.00 to 4.00, preferably from 2.20 to 3.50 Si from 0.10 to 0.80

[0099] Ti from 0.10 to 0.80

[0100] Nb 1.00 or less, preferably from 0.00 to 0.60

[0101] Co 5.00 or less, preferably from 0.00 to 2.80

[0102] Cu 3.00 or less, preferably from 0.00 to 2.80

[0103] Zr and / or Sc from 0.00 to 0.50 La 1.00 or less, preferably from 0.00 to 0.60

[0104] Ta 1.00 or less, preferably from 0.00 to 0.60

[0105] Y 1.00 or less, preferably from 0.00 to 0.50

[0106] Impurities 0.025 or less; the balance being Fe .

[0107] In addition to the advantages already found, such as avoiding the formation of brittle phases and ensuring metallurgical stability, ensuring better resistance to irradiation and corrosion by molten lead, the austenitic stainless steels of the invention also favour processes of self-generation of the surface oxide layer, even in the presence of damage and erosive phenomena (self-healing) . The ability of aluminum oxide to self-generate derives from the ease of its formation (determined by the high free energy of formation of the aluminum oxide) and the formation of protective oxide layers with a high number of vacations, which allow an optimal diffusion of oxygen through the incipient oxide layer. Optimized oxygen diffusion allows the oxide to grow faster. The formation of oxide with high vacancies is optimized by calculating the number of valence electrons (effective electrons) of the alloy elements (Veff) .

[0108] Preferably, the present austenitic stainless steel comprises an amount of Yttrium equal to 0.00 wt%, an amount of Scandium equal to 0.10 wt% and / or an amount of Zirconium equal to 0.05 wt% . This embodiment of the invention may have further advantages, including the absence of formation of a eutectic phase, a decrease in the dimension of the aforementioned B2-bcc phase, the formation of carbides with fee structure and greater resistance to lead corrosion. EXAMPLES

[0109] The inventive and comparative examples are given herein by way of illustration and are not intended to limit the invention .

[0110] Twelve austenitic stainless steel compositions marked with "AFA_R1"," AFA_R2 " , "AFA_R3", "AFA_R4", "AFA_R5",

[0111] "AFA_R6", , "AFA_R7", "AFA_R8" e "AFA_R9", "AFA_R4_C_016 " , "AFA HB-1" and "AFA HC-4" were examined according to table 1 below. The components shown in Table 1 are indicated in weight percentage value of the total weight of the steel (the amount of Iron, which constitutes the remainder, is not specified) .

[0112] 5 The austenitic stainless steel according to the invention is indicated with "AFA_R1"," AFA_R2", "AFA_R3", "AFA_R4", "AFA_R5", "AFA_R6", "AFA_R7 " , "AFA_R8" and

[0113] "AFA_R9" .

[0114] While two comparative steels, described by Y. Yamamoto 10 in "Evaluation of Mn substitution for Ni in alumina-forming austenitic stainless steels", Material Science and Engineering A 524 (2009) 176-185, are indicated with "AFA HB- 1" and "AFA HC-4".

[0115] A further comparative steel, indicated with 15 "AFA_R4_C_016" , instead has a carbon content higher than that of the steels according to the invention.

[0116] Table 1 Some steels of table 1 have been subjected to thermodynamic simulations performed with Thermocalc® Matcalc® and Dictra® programs, well known and used in this technical field.

[0117] In detail, the characteristics of the steels of Table 1 were validated through the use of thermodynamic simulations under equilibrium conditions, kinetic analyses to quantify the precipitation and evolution of the various phases, diffusion studies in order to evaluate the oxidative properties of the aforementioned innovative material, and, finally, through the determination of the Stacking Fault Energy (SEE) , essential to facilitate the activation of the twinning-induced plasticity (TWIP) mechanism. Computational results regarding thermodynamic stability and kinetic precipitation are obtained with Matcalc@.

[0118] The results of the thermodynamic and kinetic simulations for the steels of Table 1 are shown in Figures 1-7.

[0119] Figure 1 shows the molar fraction of the equilibrium phases (BCC-A2 phase, i.e. ferrite; BCC-B2 phase; sigma phase; Niobium carbide (NbC) ; and Chromium complex carbides of the M23C6 type) in the temperature range between 550 °C and 650 °C. Figures la, lb, 1c, Id, and 1g relate to the inventive steels "AFA_R2 " , "AFA_R3", "AFA_R4", "AFA_R5" and "AFA_R1", respectively.

[0120] Figures le) and If) concern the comparative steels "AFA HC-4", and "AFA HB-1".

[0121] From what is reported in Figures la) , lb) , 1c) , Id) , and 1g) it emerges that the steels according to the invention have a certain molar fraction of sigma phase that is stable from a thermodynamical point of view. However, it is clear that the steels according to the invention consist of an austenitic matrix with ordered B2-BCC phase and a minimum fraction of secondary carbide nano-precipitates.

[0122] For comparison, as can be noted in Figure If) , the comparative steel "AFA HB-1" in the temperature range of interest and in particular between 550 °C and 650 °C, comprises a non-negligible and thermodynamically stable amount of ferritic BCC-A2 phase, much higher than the steels according to the invention.

[0123] Still in comparison, Figure le) shows the results of the thermodynamic calculations inherent in the comparative steel "AFA HC-4". As can be seen in Figure le) within the temperature range of interest and in particular between 550 °C and 650 °C, it is thermodynamically stable and therefore there is a non-negligible amount of sigma phase, much higher than the steels subject-matter of the present invention.

[0124] Figure 2 compares the formation of M23C6 carbides (in dots) and Niobium or Titanium carbides (in grey) at equilibrium for a temperature of 600 °C for the inventive steels and the comparative steels of Table 1; it is evident that in the steels sub ect-matter of the present invention the formation of unwanted M23C6 carbides is absent or limited compared to the comparative steels; moreover, the greater amount of Carbon in the comparative steel "AFA_R4_C_016" , compared to the corresponding inventive steel "ARA_R4", leads to the formation of M23C6 carbides. The latter result justifies the Carbon range up to 0.15% selected for the inventive steels.

[0125] Figure 3 instead shows the TTT diagrams relative to the sigma phase of some steels according to the invention "AFA_R2" (in Fig. la) ) , "AFA_R4" (in Fig. lb) ) , "AFA_R5" (in Fig. 1c) ) , "AFA_R6" (in Fig. Id) ) indicated in Table 1 , in which the sigma phase formation iso-molar curves are below the operating temperature expected for this steel for the AFA_R2 and AFA_R4 steels , and for the AFA_R5 and AFA_R6 variants , the sigma phase formation is limited and also characteri zed by very slow kinetics .

[0126] Figure 4 instead reports the thickness of the oxide layer of Aluminum ( Fig . 4a ) ) and o f Chromium ( Fig . 4b ) ) which is formed over time thanks to the di f fusion of Aluminum and Chromium, respectively, in the material at 600 ° C ( austenitic phase ) for the steels according to the invention "AFA_R2" , "AFA_R3" , "AFA_R4 " , "AFA_R5" . The thickness and kinetics of formation of the oxide layers are in line with the optimal values to ensure adequate protection from dissolution due to corrosion, given by contact with molten lead . The results of Fig . 4a ) and Fig . 4b are obtained with simulations carried out using Dictra@ .

[0127] Figure 5 reports the results obtained from simulations of precipitation of ferritic and sigma phases at 550 ° C : Figure 5a ) compares the precipitation of the sigma phase in the comparative steel "AFA HC-4" and in the steels according to the invention from "AFA_R1" to "AFA_R6" , while Figure 5b ) compares the kinetics of ferrite formation between the comparative steel "AFA HB- 1" and the steels according to the invention from "AFA_R1" to "AFA_R6" .

[0128] It can be seen both in Figure 5a ) and in Figure 5b ) that only in the comparative steels there is , also from a kinetic point of view, the beginning of precipitation of the ferritic phase in one case and of the sigma phase in the other, demonstrating the greater metallurgical stability of the inventive steel .

[0129] Figure 6 shows the positions of some steels according to the invention, represented by geometric figures and grouped in the circle, with regard to the calculation of the free energy of formation of Aluminum oxide (AGf) and the actual concentration of valence electrons (Veff) : more negative free energies of formation and low Veff values, where the steels according to the invention are actually located, favour the formation and stability of the protective layer of AI2O3.

[0130] Figure 7 instead reports the calculated value of the SFE of the steel according to the invention "AFA_R2 " and of the austenitic matrix after heat treatment at 600 °C. These SFE values guarantee the promotion of the twinning deformation mechanism or twinning phenomenon.

[0131] Some steels in accordance with the inventions of Table 1 have also been subjected to experimental tests, including air aging tests and static liquid lead corrosion experiments.

[0132] Aging and lead corrosion resistance were studied by optical microscope inspection, X-ray diffraction (XRD) pattern analysis, Vickers hardness measurement, SEM-EDS (Scanning Electron Microscopy - Energy Dispersive Spectroscopy) analysis and using TEM-STEM (Transmission Electron Microscopy - Scanning Transmission Electron Microscopy) technique.

[0133] Regarding the aging tests, the samples, in the form of 1.5 / 2.5 cm high cylinders, were obtained by cutting EDM (Electrical Discharge Machining) from larger cylinders, after the solubilization treatment at high temperature; the oxidized surface was worked before the annealing treatment in order to be able to study the air oxidation process. These samples were placed in temperature-controlled muffles via RID (Proportional-Integral-Derivative) regulator for 1000 hours or for 2160 hours. Specifically, the samples were annealed in air at 500 °C, 600 °C and 650 °C. After aging, a section was obtained from each sample using a metallographic cutting machine, for measurements of the XRD diffraction pattern.

[0134] The XRD diffraction patterns were obtained using a Coalpha (or Co-Ka) source in Bragg-Brentano configuration.

[0135] The remainder of the sample was assembled and polished for both optical microscope (OM) analysis and SEM-EDS characterization. Prior to OM characterization, an electrochemical attack was performed to highlight the formation of carbides and precipitates. The Vickers hardness (500 g load) was measured close to the average radius of the polished sample.

[0136] Figure 8 shows the results of air aging tests after 2160 hours at 650 °C obtained by X-ray diffraction (XRD for the sample AFA_R4 (Figure 8a) and for the samples AFA_R8 and AFA_R9 (Figures 8b) and 8c) , respectively) . The XRD diffraction patterns (Figure 8) confirm the presence of an fee matrix (main phase or f cc-austenite ) , with additional peaks attributable to Niobium carbide with fee structure and a peak associated with the precipitation of a fine bcc / B2- bcc phase rich in Al and Ni .

[0137] The formation of Al- and Ni-rich precipitates after aging at 600 °C was also confirmed by TEM-STEM-EDS analysis, while the actual formation of NbC secondary nanoprecipitates was only hypothesized based on the interpretation of a STEM image of a Sc-containing sample (AFA_R8) after aging.

[0138] In particular, the aging tests at 500 °C, both theoretical and experimental, have highlighted that the temperature at 500 °C represents the limit for the start of precipitation of the B2-BCC Al- and Ni-rich phase. Precipitation becomes significantly more pronounced at 600 °C and 650 °C, a range in which both the dimensions and the amount of the Al- and Ni-rich precipitate are affected. The formation of Al- and Ni-rich precipitates after aging at 600 °C was also confirmed by TEM-STEM-EDS analysis, while the actual formation of NbC secondary nano-precipitates was only hypothesized based on the interpretation of a STEM image of a Sc-containing sample after aging.

[0139] Regarding the effect of Zr in AFA_R9 and Sc in AFA_R8, the experimental results, in line with the simulations, indicate that their main contribution consists in the inclusion of these reactive elements, i.e. Sc and Zr, in the formation of carbides with fee structure.

[0140] Furthermore, it appears that the addition of reactive elements, such as Zr and Sc, also influences the precipitation kinetics of the B2-BCC phase. The presence of Y in AFA_R4 and of the eutectic phase favours instead the precipitation of the Al- and Ni-rich phase, making the dimensions of the precipitate more visible compared to AFA_R8 containing Sc, which shows smaller dimensions of the B2-BCC precipitate .

[0141] The hypothesized precipitation mechanism foresees that the separation of Al and Ni originates in the form of an fee solid solution, progressively transforming into a bcc structure with an increasing level of order. From both the XRD diffraction patterns and the SEM data it clearly emerges that the growth of the B2-BCC phase takes place in a preferential direction.

[0142] With specific reference to the X-ray diffraction results for the sample according to the invention AFA_R9 containing Zr, the following emerges:

[0143] - at 500 °C: no phase transformation or precipitation occurs; as observed after 1000 hours, even after 2160 hours the composition shows no microstructural changes. The XRD pattern shows the presence of peaks belonging to the fee solid solution and to the NbC carbide, as in the samples after the solubilization treatment;

[0144] - at 600 °C: the phase transformation begins to manifest itself, in particular the precipitation of a darker phase in the matrix of the fee solid solution, visible by SEM imaging. The phase transformation is confirmed by XRD diffraction, where, together with the peaks of the fee, a peak attributable to a B2-BCC phase begins to appear. The EDS analysis of the darker phase, barely visible in the SEM image matrix, was not possible due to the very small dimensions of the precipitate. Nb-Zr-rich carbides have been identified;

[0145] - at 650 °C (Figure 8 c) ) : the phase transformation takes place as at 600 °C, but with faster kinetics, and larger precipitates with dark contrast, detected by SEM-EDS, are formed. The larger dimensions of the precipitates allows an EDS analysis to be carried out, which reveals that the darker precipitates have a higher content of Al and Ni . Consistently, the XRD pattern is compatible with a B2-BCC phase incorporated into an fee solid solution, making it highly likely that the phase present at 600 °C is also an Al and Ni-rich phase. Carbides rich in Nb-Zr and Fe-Zr were identified.

[0146] Regarding the lead corrosion resistance tests, no significant penetration of lead was observed in the samples according to the invention. A protective layer of aluminum oxide, characterised by a dense, compact structure and a limited thickness (< 1 pm) , has developed on the surface. This layer offers passivation and protection against further corrosive phenomena. However, it was found that the samples AFA_R8 and AFA_R9, containing Sc and Zr, respectively, further promote the formation of this protective layer . In fact , it has been observed that in the samples free of Sc and Zr such as , for example , in AFA_R1 , AFA_R2 and AFA_R7 , under the same corrosion test conditions , the oxide layer is thicker and contains other elements such as Mn and / or Cr .

[0147] In detail , the lead corrosion tests consist in immersing the sample in static liquid lead at 600 ° C for 1000 hours , with an O2content equal to 10~7wt% .

[0148] The results of the SEM-EDS analysis are reported in Figures 9- 11 for the samples AFA_R8 , AFA_R9 and AFA_R4 , respectively . For all the areas analysed, it is shown a BSE (Backscattered Electrons ) image near the surface in contact with the lead and subsequently an EDS scan carried out through the edge of the sample .

[0149] From the oxygen and aluminum EDS spectra of Figures 9 and 10 , an initial peak is observed which highlights the formation of a protective Al- and O-rich oxide layer along the edge of the sample .

[0150] From Figure 11 relative to the sample AFA_R4 , free of Sc and Zr, but with Y, it was observed that the protective oxide layer is quite thick and also enriched in Cr and / or Mn and not only in Al , as shown by the EDS spectra of Figure 11 .

[0151] The results on aging and corrosion resistance show how the presence of Sc and Zr ( in the samples according to the invention AFA_R8 and AFA_R9 ) af fects the precipitation kinetics of B2 phase .

[0152] It also emerges how the presence of Y ( in the samples according to the invention as AFA_R4 ) tends to stabili ze the Ni-Fe-Y eutectic phase which, due to its low melting point , generates negative ef fects during subsequent thermomechanical processes . In addition to the precipitation of B2 phase, the presence of Sc and Zr in AFA_R8 and AFA_R9 alters the composition of the metal carbides, resulting, for example, in the formation of Nb-Fe-rich carbides in addition to the traditional Nb-Ti-rich carbides of the samples, such as AFA_R4, containing Y. For example, the sample AFA_R9, has a higher concentration of carbides rich in Nb-Zr and Fe-Zr compared to the most common carbides rich in Nb-Ti. However, the effect of the inclusion of Zr in the carbides remains poorly understood and needs further investigation.

[0153] Based on the Vickers hardness analysis, it emerges that the presence of Sc and / or Zr favours a precipitation strengthening effect compared to the AFA samples free of Zr and / or Sc.

[0154] The results of the static corrosion test confirm the formation of a protective Al-rich oxide layer with the thickness of hundreds of nanometres on the surface of all samples in accordance with the invention, without f erritization or penetration of Pb .

[0155] In addition to the aluminum-rich oxide layer, Cr- and Mn-rich oxide nodules were observed in some regions of AFA_R4 containing Y.

[0156] In addition, the same test conditions were also applied to samples with minimal amount of Y (0.06%) and with massive amount of Y (0.10%) . No significant performance alterations were observed in these variants compared to AFA_R4 with average Y content (0.08%) . The precipitation start temperature of B2 phase remains at 600 °C. All conditions confirm the presence of a stable austenitic matrix, of the Ni-Fe-Y eutectic phase and of metal carbides rich in Nb, Ti and Fe even after 2160 hours of aging in air, as envisaged by the precipitation kinetics simulations. As in AFA_R4, a thick oxide layer was observed on the surface of these variants after the static corrosion test . Instead of being rich in Al , the oxide layer is characteri zed by a mixture of Al , Cr and / or Mn . Neither the ferritic phase nor the sigma phase was observed, in any of the compositions subj ect-matter of the invention, respectively, in the AFA variants susceptible to the ferritic and sigma phase , not even after aging at 650 ° C for 2160 hours or after the Pb corrosion test at 600 °C for 1000 hours . The experimental results obtained are in excellent agreement with the predictions provided by the s imulations , confirming the validity of the model and of the approach followed for the design of the innovative austenitic steel with medium manganese content .

[0157] The results therefore confirm that the austenitic stainless steels according to the invention satis fy all the purposes of the present illustration .

[0158] Finally, it is understood that further modi fications and variants can be made to the steels according to the invention that do not fall outside the scope of the appended claims .

Claims

CLAIMS1. Use of an austenitic stainless steel with medium Manganese content and aluminum oxide formation for making at least a part of an apparatus, equipment or device of a nuclear plant, in direct contact with a primary cooling fluid, circulating in a reactor, wherein the austenitic stainless steel comprises in weight percentage (wt%) :C from 0.05 to 0.15Cr from 10.50 to 13.50Ni from 9.0 to 19.0Al from 2,0 to 4,0Si from 0.10 to 0.80Ti from 0.10 to 0.80Nb 1.00 or lessCo 5.00 or lessCu 3.00 or lessZr and / or Sc from 0.00 to 0.50 La 1.00 or lessTa 1.00 or lessY 1.00 or less impurities 0.025 or less; wherein Mn is from 5.00 to 14.00; the balance being Fe .

2. Use as claimed in claim 1, wherein the austenitic stainless steel contains 5.50-13.00 wt% of Mn, preferably 6.00-12.00 wt%.

3. Use as claimed in claim 1 or 2, wherein the austenitic stainless steel contains 10.70 -13.50 wt% of Cr, preferably 11.00-13.50 wt%.

4. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel contains 10.00-17.00 wt% of Ni, preferably 10.50-16.50 wt%, more preferably 11.00-16.30 wt%, even more preferably 11.50-16.00 wt%.

5. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel contains 2.20-3.50 wt% of Al, preferably 2.20-3.00 wt%.

6. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel contains 0.15-0.60 wt% of Si and / or Ti.

7. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel contains 0.00-2.80 wt% of Co and / or Cu.

8. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel contains 0.00-0.60 wt% of Ta and / or La and / or Nb .

9. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel contains 0.00-0.50 wt% of Zr and / or Y, preferably 0.00-0.30 wt%.

10. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel contains 0.05-0.14 wt% of C, preferably 0.10-0.14 wt%.

11. Use as claimed in one of the preceding claims, wherein the austenitic stainless steel is configured to make at least a part of an apparatus, equipment or device, in direct contact with the primary cooling fluid, circulating in the reactor, which performs a process for the production of nuclear energy or a nuclear plant.

12. Use as claimed in claim 11, wherein the device is a nuclear reactor, preferably a lead-cooled reactor.

13. Use as claimed in claim 12, wherein the nuclear reactor is at an operating temperature between 500 °C and 700 °C.

14. Use as claimed in claim 12 or 13, wherein the austenitic stainless steel is configured to be exposed to aneutron flux in the nuclear reactor, in particular a fast reactor .

15. Apparatus, equipment or device of a nuclear plant, in direct contact with a primary cooling fluid, circulating in a reactor, or used in a process for the production of nuclear energy, comprising at least a part made of austenitic stainless steel with medium manganese content and aluminum oxide formation, wherein said austenitic stainless steel comprises in weight percentage (wt%) :C from 0.05 to 0.15Cr from 10.50 to 13.50Ni from 9.00 to 19.00, preferably from 10.00 to 17.00 Al from 2.00 to 4.00, preferably from 2.20 to 3.50 Si from 0.10 to 0.80Ti from 0.10 to 0.80Nb 1.00 or less, preferably from 0.00 to 0.60Co 5.00 or less, preferably from 0.00 to 2.80Cu 3.00 or less, preferably from 0.00 to 2.80Zr and / or Sc from 0.00 to 0.50 La 1.00 or less, preferably from 0.00 to 0.60Ta 1.00 or less, preferably from 0.00 to 0.60Y 1.00 or less, preferably from 0.00 to 0.50 impurities 0.025 or less; wherein Mn is from 5.00 to 14.0; the balance being Fe .

16. Apparatus, equipment or device as claimed in claim 15, wherein the device is a nuclear reactor, preferably a lead-cooled reactor.

17. Nuclear plant comprising at least one apparatus, equipment or device as claimed in claim 15.

18. Nuclear plant as claimed in claim 17, wherein the device is a nuclear reactor, preferably a lead-cooledreactor .

19. Austenitic stainless steel with medium Manganese content and aluminum oxide formation in a lead-cooled reactor, comprising in weight percentage (wt%) : Mn from 5.00 to 14.00, preferably from 5.50-13.00%C from 0.05 to 0.15, preferably from 0.05 to 0.14Cr from 10.50 to 13.50, preferably from 10.70 to 13.50 Ni from 9.00 to 19.00, preferably from 10.00 to 17.00 Al from 2.00 to 4.00, preferably from 2.20 to 3.50 Si from 0.10 to 0.80Ti from 0.10 to 0.80Nb 1.00 or less, preferably from 0.00 to 0.60Co 5.00 or less, preferably from 0.00 to 2.80Cu 3.00 or less, preferably from 0.00 to 2.80Zr and / or Sc from 0.00 to 0.50 La 1.00 or less, preferably from 0.00 to 0.60Ta 1.00 or less, preferably from 0.00 to 0.60Y 1.00 or less, preferably from 0.0 to 0.50 Impurities 0.025 or less; the balance being Fe .

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