On the Microstructure and Properties of the Nb-23Ti-5Si-5Al-5Hf-5V-2Cr-2Sn (at.%) Silicide-Based Alloy—RM(Nb)IC

: The microstructure, isothermal oxidation, and hardness of the Nb-23Ti-5Si-5Al-5Hf-5V-2Cr-2Sn alloy and the hardness and Young’s moduli of elasticity of its Nb ss and Nb 5 Si 3 were studied. The alloy was selected using the niobium intermetallic composite elaboration (NICE) alloy design methodology. There was macrosegregation of Ti and Si in the cast alloy. The Nb ss , α Nb 5 Si 3 , γ Nb 5 Si 3 , and HfO 2 phases were present in the as-cast or heat-treated alloy plus TiN in the near-the-surface areas of the latter. The vol.% of Nb ss was about 80%. There were Ti- and Ti-and-Hf-rich areas in the solid solution and the 5-3 silicide, respectively, and there was a lamellar microstructure of these two phases. The V partitioned to the Nb ss , where the solubilities of Al, Cr, Hf, and V increased with increasing Ti concentration. At 700, 800, and 900 ◦ C, the alloy did not suffer from catastrophic pest oxidation; it followed parabolic oxidation kinetics in the former two temperatures and linear oxidation kinetics in the latter, where its mass change was the lowest compared with other Sn-containing alloys. An Sn-rich layer formed in the interface between the scale and the substrate, which consisted of the Nb 3 Sn and Nb 6 Sn 5 compounds at 900 ◦ C. The latter compound was not contaminated with oxygen. Both the Nb ss and Nb 5 Si 3 were contaminated with oxygen, with the former contaminated more severely than the latter. The bulk of the alloy was also contaminated with oxygen. The alloying of the Nb ss with Sn increased its elastic modulus compared with Sn-free solid solutions. The hardness of the alloy, its Nb ss , and its speciﬁc room temperature strength compared favourably with many refractory metal-complex-concentrated alloys (RCCAs). The agreement of the predictions of NICE with the experimental results was satisfactory.


Introduction
Interest in RM(Nb)ICs, i.e., refractory metal (RM) intermetallic composites based on Nb, also known as Nb-silicide-based alloys or Nb in situ composites, as potential metallic ultra-high temperature materials (UHTM) to replace Ni-based superalloys in hightemperature structural applications in aero engines, has arisen from early research on binary Nb-Si alloys and experimental data that showed that such RMICs could offer an attractive combination of high temperature strength, creep resistance, and room temperature fracture toughness (e.g., see [1]). Another contributing factor was the phase equilibria of the binary system [2] that makes it possible to cast in situ composites consisting of a RM terminal solid solution, namely, the bcc Nb solid solution (Nb ss ), and the creep-resistant tetragonal Nb 5 Si 3 silicide. The latter can have the high-temperature βNb 5 Si 3 (tI32, D8 m , prototype W 5 Si 3 ) or the low-temperature αNb 5 Si 3 (tI32, D8 l , prototype Cr 5 B 3 ) structure that are isomorphous but with distinctly different lattice parameters α and c [2].
The early research focused on alloys of the Nb-Si-Ti-Hf-Cr-Al system, a typical example of which is the MASC alloy Nb-25Ti-16Si-8Hf-2Al-2Cr (at.%) [3]. The initial research High-purity (wt.%) elements (Nb 99.99%, Al 99.999%, Si 99.999%, Ti 99.99%, Cr 99.99%, Sn 99.99%, V 99.7%, and Hf 99.8%) were used as the starting materials to make button/ingots of 300 g weight of the alloy in an argon atmosphere using arc melting with a non-consumable tungsten electrode and a water-cooled copper crucible. In order to improve homogeneity, melting was repeated five times by turning over the button/ingot after each melt. The ingots were cooled down to ambient temperature inside the arc melting furnace. Cubic specimens that were cut from the bulk of the ingot were wrapped in Ta foil, placed in an alumina boat, and were heat treated at 1500 • C in a tube furnace for 100 h under a Ti-gettered argon atmosphere. The heat treatment temperature was decided following study of small-sized samples (results not shown) using differential scanning calorimetry (Stanton Redcroft STA-1500 DSC/DTA, Thermal Scientific plc., Odessa, TX, USA) to ensure avoidance of liquation during heat treatment. The DSC specimens were also cut from the bulk of the button/ingot. Standard metallographic preparation was used to prepare specimens for study by X-ray diffraction (XRD) and electron beam microprobe analysis (EPMA). A Philips X-ray diffractometer with monochromatic Cu Kα (λ = 1.5418) radiation was used for the identification of the phases, which was done using JCPDS data. Secondary electron (SE) and backscatter electron (BSE) imaging and quantitative chemical analyses were performed using a JEOL 8600 electron probe micro-analyser (JEOL, Tokyo, Japan). High purity elements and Hf 2 Si [33], Al 2 O 3 , and BN, which had been polished to a 1 µm finish, were used for standardisation purposes. Analysis was performed at 15 kV, 20 nA, and about a 1 µm diameter beam, and, at each measurement, the probe diameter was adjusted so as to achieve a dead time less than 20%. At least ten analyses were done for each phase in different parts (top, bulk, and bottom) of the button/ingot. Average concentrations and the minimum and maximum concentrations of at least ten analyses of an element in a phase or region of the ingot are given in the tables in this article. Measurements of the area fraction of Nb ss were performed using software available on the microprobe, and the contrast of the phases was adjusted under BSE imaging. Area fractions were measured for the same Metals 2021, 11, 1868 4 of 34 areas that were used for large area analyses in EPMA. At least ten measurements were taken for each alloy, all at the same magnification (X350).
Thermogravimetric analysis (TGA) of the alloy was carried out using a Stanton Redcroft thermo-balance (Thermal Scientific plc., Odessa, TX, USA) equipped with an alumina tube furnace. Small cube-shaped specimens (5 × 5 × 5 mm 3 ) were cut from the as-cast ingot. Each specimen was placed in a small cylindrical alumina crucible, which was then placed on the thermo-balance. The tube furnace was preheated to the selected temperature. The bottom of the crucible was in direct contact with a thermocouple. The isothermal oxidation of the alloy was evaluated at 700, 800, and 900 • C for up to 100 h. The oxides that formed on the surface of the specimens and spalled off were ground to powder and studied using XRD. EPMA was performed on cross sections of oxidised specimens.
The Vickers hardness of the alloy in the as-cast and heat-treated conditions was measured with a load of 10 kg. At least ten measurements were taken. The density of the alloys was measured using the Archimedes' principle and a Sartorious LA2305 electronic precision balance (Sartorius Lab Instruments GmbH & Co. KG, Göttingen, Germany), equipped with a density determination kit. The average and minimum and maximum values of the hardness and density of the alloys are given in this article.
The as-cast and heat-treated microstructures made it possible to measure the hardness and elastic moduli of Nb ss and Nb 5 Si 3 using nano-indentation. For this work, we used a Hysitron Inc. TriboScope ® Nanomechanical Testing System (Hysitron Inc., Minneapolis, MN, USA) equipped with the software package TriboScope version 3.5L and a TriboIndenter that was fitted with a Berkovich style tip, which is a flat three-sided, diamond-tipped pyramid with a total induced angle of 142.3 • and a half angle of 65.35 • [34]. The radius of curvature of the tip was 150 nm. This signified that only a low level of detail could be achieved during imaging compared with a sharper tip, owing to imaging using an atomic force microscope (AFM) that was linked to the nano-indenter. A load of 7000 µN was used for the nano-indentation measurements.

As-Cast (NV1-AC)
The actual chemical composition of NV1-AC was 51.4Nb-23Ti-6.2Si-5Al-2.1Cr-5.4Hf-2.2Sn-4.7V. This was the average composition of all the analyses taken from the bottom, bulk, and top areas of the as-cast button/ingot. The density of the alloy was 7.38 g/cm 3 ( Table 1). The microstructure in all parts of the ingot consisted of Nb ss , Nb 5 Si 3 , HfO 2 , and Nb ss + Nb 5 Si 3 fine-scale lamellar microstructure (Figures 1-3). The latter was "connected" with Nb 5 Si 3 that had Nb/(Ti + Hf) less than one and an average <Si> of about 39 at.% ( Figure 3) but not with Nb 5 Si 3 with Nb/(Ti + Hf) greater than one and an average <Si> of about 37 at.% ( Figure 2b) (<Si> = Al + Si + Sn). The XRD indicated the possible presence of Nb 3 Si. However, exhaustive study using EPMA to find Nb 3 Si, the presence or absence of which is important to understand how the lamellar microstructure formed (see discussion), did not confirm the existence of Nb 3 Si in NV1-AC.  For example, in Figure 2a,b, the areas of darker contrast near the Nb5Si3 silicide are Ti rich, and the brighter contrast Nb5Si3 grains are Hf rich (see below). The Ti-rich Nb5Si3, regarding its high Ti concentration, corresponded to the Ti-rich Nb5Si3 reported in other Ti-containing RM(Nb)ICs [35,36]. The Ti-rich Nbss was observed only in the top and bottom of NV1-AC. In Table 2, the Nb5Si3 is also designated as Ti-and Hf-rich Nb5Si3, owing to its high Hf and Ti contents (see below).   Compared with the nominal composition, the NV1-AC was slightly richer in Si with an average concentration of 6.2 at.%. The concentrations of the other alloying elements were very close to the nominal ones (see Section 2). There was macrosegregation of Ti and Si, the concentrations of which were in the ranges: 19.4 to 24.3 at.% (MACTi = 4.9 at.%) and 3.9 to 7 at.% (MACSi = 3.1 at.%), respectively. The XRD (Figure 1a) indicated the pres-  ence of tetragonal αNb5Si3 and hexagonal 5-3 silicide. There were three peaks that are characteristic of αNb5Si3 compared with six peaks for the hexagonal γNb5Si3 and four peaks for the hexagonal Ti5Si3, all shared with other phases. The latter two silicides are isomorphous [24].    In the areas close to the top of the button/ingot, the Si concentration was in the range 3.9 to 5.9 at.%, and the microstructure consisted of Nbss, Nb5Si3, and a lamellar microstructure of these phases. The Ti-rich Nbss was found either next to the Ti-and-Hf-rich Nb5Si3, or surrounded the HfO2, which was often found adjacent to the Ti-and-Hf-rich Nb5Si3. The Ti-rich areas in the Nbss and the Ti-rich Nb5Si3 exhibit a dark contrast in the microstructures shown in the Figure 2a,b. However, the Ti-and Hf-rich areas in the Nb5Si3 exhibited bright contrast. The Ti concentration varied from 18.8 to 20.8 at.% in the Ti-rich Nb5Si3 and from 22.4 to 25.2 at.% in the Ti-and-Hf-rich Nb5Si3. Despite the fact that the Ti-and-Hf-rich Nb5Si3 was richer in Ti compared with the Ti-rich Nb5Si3, it exhibited a brighter contrast owing to its very high Hf concentration, which varied from 12 to 12.7 at.% compared with   The volume fraction of Nb ss in NV1-AC was about 0.8 (Table 1). Data for the average chemical composition of the phases in the whole button/ingot are given in Table 2. Note that the very fine scale of the lamellar structure prevented reliable analysis of the chemical composition of its phases using EPMA. Strong partitioning of Cr, Ti, and V in the solid solution and of Hf and Ti in the silicide, which led to varying BSE contrast at the interface between these phases, was noted (Figures 2 and 3). For example, in Figure 2a,b, the areas of darker contrast near the Nb 5 Si 3 silicide are Ti rich, and the brighter contrast Nb 5 Si 3 grains are Hf rich (see below). The Ti-rich Nb 5 Si 3 , regarding its high Ti concentration, corresponded to the Ti-rich Nb 5 Si 3 reported in other Ti-containing RM(Nb)ICs [35,36]. The Ti-rich Nb ss was observed only in the top and bottom of NV1-AC. In Table 2, the Nb 5 Si 3 is also designated as Ti-and Hf-rich Nb 5 Si 3 , owing to its high Hf and Ti contents (see below).
Compared with the nominal composition, the NV1-AC was slightly richer in Si with an average concentration of 6.2 at.%. The concentrations of the other alloying elements were very close to the nominal ones (see Section 2). There was macrosegregation of Ti and Si, the concentrations of which were in the ranges: 19.4 to 24.3 at.% (MACTi = 4.9 at.%) and 3.9 to 7 at.% (MACSi = 3.1 at.%), respectively. The XRD (Figure 1a) indicated the presence of tetragonal αNb 5 Si 3 and hexagonal 5-3 silicide. There were three peaks that are characteristic of αNb 5 Si 3 compared with six peaks for the hexagonal γNb 5 Si 3 and four peaks for the hexagonal Ti 5 Si 3 , all shared with other phases. The latter two silicides are isomorphous [24].
In the areas close to the top of the button/ingot, the Si concentration was in the range 3.9 to 5.9 at.%, and the microstructure consisted of Nb ss , Nb 5 Si 3 , and a lamellar microstructure of these phases. The Ti-rich Nb ss was found either next to the Ti-and-Hf-rich Nb 5 Si 3 , or surrounded the HfO 2 , which was often found adjacent to the Ti-and-Hf-rich Nb 5 Si 3 . The Ti-rich areas in the Nb ss and the Ti-rich Nb 5 Si 3 exhibit a dark contrast in the microstructures shown in the Figure 2a,b. However, the Ti-and Hf-rich areas in the Nb 5 Si 3 exhibited bright contrast. The Ti concentration varied from 18.8 to 20.8 at.% in the Ti-rich Nb 5 Si 3 and from 22.4 to 25.2 at.% in the Ti-and-Hf-rich Nb 5 Si 3 . Despite the fact that the Ti-and-Hf-rich Nb 5 Si 3 was richer in Ti compared with the Ti-rich Nb 5 Si 3 , it exhibited a brighter contrast owing to its very high Hf concentration, which varied from 12 to 12.7 at.% compared with 6 to 7 at.% in the Ti-rich Nb 5 Si 3 . The aforementioned Ti-rich and Ti-and-Hf-rich Nb 5 Si 3 was observed more often as two discrete grains in the top of the ingot, unlike the bulk and bottom of the ingot where the segregation of Hf and Ti was observed mainly at the exterior or the interior of a single Nb 5 Si 3 grain. The lamellar microstructure of Nb ss and Nb 5 Si 3 was formed adjacent to the Ti-and-Hf-rich Nb 5 Si 3 . The solubility of Cr in Nb 5 Si 3 was low and exceeded 1 at.% in the Ti-and-Hf-rich Nb 5 Si 3 . Similarly, the V concentration in the Nb 5 Si 3 was low and varied from 1.4 to 2.9 at.% in the Ti-and-Hf-rich Nb 5 Si 3 and from 1.2 to 2.6 at.% in the Ti-rich Nb 5 Si 3 .
In the bulk of the button/ingot, the Si concentration was in the range 5.8 to 7 at.%. The microstructure was slightly coarser and consisted of the same phases as in the top ( Figure 2b) but with two main exceptions. Firstly, the Ti-rich parts of the Nb ss were also rich in V and Cr where the concentrations of these elements varied from 7.8 to 10.2 at.% and 3.6 to 8.4 at.%, respectively. This Nb ss is designated as Ti-, V-, and Cr-rich Nb ss in Table 2. The solubility of Al, Si, and Sn in the Nb ss was reduced. Secondly, the Nb 5 Si 3 exhibited Ti-and Hf-rich areas in different parts of the same grain. The Ti-rich Nb 5 Si 3 , unlike that observed in the top of the ingot, was also rich in Hf, and the Ti-and-Hf-rich Nb 5 Si 3 was observed with Hf concentrations as high as 17 at.%. In other words, the Hf solubility in Nb 5 Si 3 was much higher in the bulk of the ingot, compared with the top. Differences in the composition of the 5-3 silicide should also be noted (compare composition of 5-3 silicide in Figures 2b and 3). The solubility of Sn was negligible in both the Ti-and-Hf-rich and the Ti-rich Nb 5 Si 3 . The lamellar microstructure ( Figure 3) had average composition 39.6Nb-23.4Ti-18.2Si-7.7Hf-4.6Al-3.6V-1.6Sn-1.3Cr (at.%), similar to that observed in the top of the ingot.
In the bottom of the ingot, the Si concentration was in the range of 4.9 to 6.6 at.%. The Ti-rich Nb ss had a similar composition to that in the top. In the Nb 5 Si 3, the Hf segregation was not as strong as in the bulk. The solubility of Al in the Nb ss was as high as 6.1 at.%. The average Si concentration in the lamellar microstructure was 20.8 at.%, which was slightly higher compared with the top and bulk of the ingot.

Heat-Treated (NV1-HT)
The actual chemical composition of NV1-HT was Nb-23.3Ti-5Si-4.9Al-1.8Cr-5Hf-2Sn-3.7V. The XRD data ( Figure 1b) suggested the presence of αNb 5 Si 3 and hexagonal 5-3 silicides and Nb 3 Si. However, the presence of Nb 3 Si in the microstructure was not confirmed by exhaustive study using EPMA. Compared with NV1-AC, the number of peaks for αNb 5 Si 3 and γNb 5 Si 3 and Ti 5 Si 3 had increased and decreased, respectively. The typical microstructure in the bulk of NV1-HT consisted of Nb ss , Nb 5 Si 3 , and HfO 2 (see Figure 2d). There was no evidence of the lamellar microstructure seen in NV1-AC. Homogenisation had taken place in both the Nb ss and Nb 5 Si 3 ( Table 2). The HfO 2 was found at the interface of Nb ss and Nb 5 Si 3 . The vol.% of the Nb ss had not changed significantly compared with NV1-AC (see Table 1). The Cr solubility in the Nb ss increased after the heat treatment and was as high as 5.6 at.% in some parts of the solid solution. On the contrary, the average solubility of V in the Nb ss had decreased to 2.4 at.%. The Ti concentration in the Nb ss was about 23.6 at.%, close to the nominal Ti concentration in the alloy. The solubilities of Al, Sn, and Hf in the Nb ss did not change after the heat treatment. The solubilities of Ti and V in the Nb 5 Si 3 were reduced, while those of Al and Hf had not changed. However, the solubility of Cr had increased in the Nb 5 Si 3 .
In the microstructure near the surface of NV1-HT there were Ti-rich regions in the Nb 5 Si 3 (see Table 3). This was not observed in the bulk of NV1-HT (Table 2). Furthermore, the Hf concentration in the Nb 5 Si 3 was reduced compared with the Nb 5 Si 3 in the bulk of NV1-HT and NV1-AC (Tables 2 and 3), and the Ti concentration in the Ti-rich Nb 5 Si 3 was similar to that of the Ti-and Hf-rich Nb 5 Si 3 in NV1-AC, while in the Nb 5 Si 3 , the average Ti and Hf concentrations barely exceeded 15 and 1 at.%, respectively. Thus, near the surface area of NV1-HT the Nb 5 Si 3 had Nb/(Ti + Hf) greater than one, whereas in the bulk, this ratio was less than one. No Hf was analysed in the Nb ss , and the Al, V, Cr, and Sn concentrations were in the solubility range of these elements as in the Ti-rich Nb ss in NV1-AC. The vol.% and the size of the Hf oxide particles had increased compared with the bulk of NV1-HT (Figure 2c,d). Together with the phases described above, Ti nitride of irregular shape and various sizes had formed near the surfaces of NV1-HT specimens (Figure 2c), up to about 1 mm below the surface.
The Ti nitride grains were found either adjacent to Nb 5 Si 3 or between the Nb 5 Si 3 and HfO 2 . The presence of Ti nitride was confirmed only by EPMA (Table 3) owing to its small volume fraction in NV1-HT.

Oxidation
The mass change of NV1 in isothermal oxidation at 700, 800, and 900 • C was compared with the "reference" MASC alloy (see the introduction) and the alloys NV3 [37], NV4 [38], NV5 [37], NV6 [39], and NV8 [38] in Figure 4. Large button/ingots of all these alloys were made using arc melting. The oxidised specimens are shown in Figure 5, and the oxidation rate constants are given in Table 4. In the latter table, the vol.% of Nb ss in the Sn-containing alloys with similar Ti concentration to the MASC alloy is also given. Note (i) that the MASC alloy was Sn-free, (ii) that the vol.% Nb ss of the alloy NV4 [38] was about 51.3%, (iii) that the alloy NV4 was very Ti-rich (45 at.%, nominal), and (iv) that the Sn content of the alloys NV3 to NV6 and NV8 was 5 at.% (nominal) compared with 2 at.% Sn in NV1 (for nominal composition of the alloys, see the Table 1 in the Appendix A). The alloy NV9 [39] suffered catastrophic pest oxidation at all three temperatures. Data for this alloy is not included in the Figure 4. Furthermore, because of (iii), data for the oxidation rate constants of the alloy NV4 is not included in the Table 4. Table 4. Oxidation rate constants of the alloys NV1, NV3, NV4, NV5, NV6, and NV8 at 700 • C, 800 • C, and 900 • C.

Alloy
Vol.% Nb ss At 700 • C, the NV series of alloys (see the previous paragraph and the Table 1 in the Appendix A) followed parabolic oxidation kinetics, whereas the MASC alloy exhibited parabolic oxidation for the first 10 h and then linear oxidation until the end of the oxidation ( Table 4). The ranking of alloys from best to worst oxidation was NV5, NV8, NV6, NV3, NV1, and MASC. Partial pest oxidation was observed only for the MASC alloy ( Figure 4). The NV1 alloy did not pest; there was some "blistering" along its edges; and its mass change was 3.5 mg/cm 2 . At 800 • C, the alloys NV3, NV6, and MASC followed linear oxidation kinetics, and the alloys NV1, NV2, NV5, and NV8 followed parabolic oxidation kinetics. The ranking of alloys from best to worst oxidation was NV5, NV8, NV1, NV6, MASC, and NV3. Pest oxidation was observed only for the NV3 alloy. The MASC alloy formed a Maltese cross (Figure 4). Even though the specimen of the NV1 alloy did not pest, there was evidence of oxide growth around one of the specimen corners (see the bottom left-hand corner in Figure 4), and some oxide from this area spalled off. The mass change of NV1 was 10 mg/cm 2 . At 900 • C, all the alloys followed linear oxidation. The ranking of alloys from best to worst oxidation was NV1, NV3, NV5, NV6, MASC, and NV8. Pest oxidation was observed for the alloys NV6 and NV8. The MASC alloy had formed a Maltese cross that separated in six pieces and a small core (Figure 4). The mass change of NV1 was 30 mg/cm 2 after 70 h, and its scale spalled off.   The oxidised specimens of the alloy NV1 at 800 and 900 °C were studied using XRD and EPMA. The actual composition of the bulk of NV1 at 800 °C was 34.8Nb-14.9Ti-3.9Si-   The oxidised specimens of the alloy NV1 at 800 and 900 °C were studied using XRD and EPMA. The actual composition of the bulk of NV1 at 800 °C was 34.8Nb-14.9Ti-3.9Si- The oxidised specimens of the alloy NV1 at 800 and 900 • C were studied using XRD and EPMA. The actual composition of the bulk of NV1 at 800 • C was 34.8Nb-14.9Ti-3.9Si-3.6Al-2.4Cr-3.7Hf-1.4Sn-1.4V-33.9O. The alloy was heavily contaminated with oxygen. The scale that separated from the specimen during isothermal oxidation at 900 • C for 100 h (Figure 4) was crushed and then studied using XRD. The latter showed peaks that corresponded to the Nb 2 O 5 , TiNbO 4 , Nb 2 O 5 ·TiO 2 , 3Nb 2 O 5 ·TiO 2 , 5Nb 2 O 5 ·2TiO 2 , CrNbO 4 , AlNbO 4 , SiO 2 , TiO 2 , and HfO 2 oxides ( Figure 6). The same oxides were observed at 800 • C. The cross section of the oxidised specimen at 800 • C was studied by EPMA. The EPMA analysis data is given in Table 5. The microstructures of the oxide scale and bulk are shown in Figure 7. Table 6 summarises the phases and characteristic features of the oxidised alloy.
3.6Al-2.4Cr-3.7Hf-1.4Sn-1.4V-33.9O. The alloy was heavily contaminated with oxygen. The scale that separated from the specimen during isothermal oxidation at 900 °C for 100 h (Figure 4) was crushed and then studied using XRD. The latter showed peaks that corresponded to the Nb2O5, TiNbO4, Nb2O5·TiO2, 3Nb2O5·TiO2, 5Nb2O5·2TiO2, CrNbO4, AlNbO4, SiO2, TiO2, and HfO2 oxides ( Figure 6). The same oxides were observed at 800 °C. The cross section of the oxidised specimen at 800 °C was studied by EPMA. The EPMA analysis data is given in Table 5. The microstructures of the oxide scale and bulk are shown in Figure 7. Table 6 summarises the phases and characteristic features of the oxidised alloy. The scale was non-uniform, and its thickness was ≤20 μm. Cracks perpendicular to the substrate were formed in the scale, which consisted of Nb-and Ti-rich oxide (Nb/Ti ≈ 2.8), Si-rich oxide (Nb/Ti ≈ 1.2 and Si/Ti ≈ 1.5), and HfO2 as well as oxidised Nbss and Nb5Si3 (Tables 5 and 6, Figure 7b). The substrate just below the scale was also cracked with cracks running parallel and perpendicular to the surface of the specimen. In some parts of the cross section, about 50 μm below the oxide scale, there was separation of the subscale alloy from the bulk (Figure 7a). Parts of the Ti-rich Nbss in the subscale alloy were heavily oxidised (oxygen content in the range 30 to 50 at.%; see "peppered" darker contrast area on the right-hand side of analysis point 2 in Figure 7b), and the oxidised Nbss exhibited cracks running parallel to the interface (Figure 7b).
At 800 °C, and only in very few areas along the scale/substrate interface, a very thin layer was observed that exhibited a very bright contrast under BSE imaging. The thickness of these layers was too small for quantitative analysis. Line scans across such layers confirmed that they were Sn rich ( Figure 8). Analysis of such layers became possible after oxidation at 900 °C, where they became more continuous, and their thickness increased. The analyses showed that the bright contrast layer that formed at 900 °C consisted of Nb3Sn and Nb6Sn5 (see Figure 9). The average composition of Nb3Sn was 30.8Nb-14.8Ti-4.4Si-3.9Al-3.7Hf-20.2Sn-0.7V-21.5O, whereas that of Nb6Sn5 was 51.1Nb-1Ti-2.2Si-1.8Al-3.7Hf-39.9Sn-0.3V, i.e., both compounds were Cr free, and the Nb6Sn5 was not contaminated with oxygen.      The scale was non-uniform, and its thickness was ≤20 µm. Cracks perpendicular to the substrate were formed in the scale, which consisted of Nb-and Ti-rich oxide (Nb/Ti ≈ 2.8), Si-rich oxide (Nb/Ti ≈ 1.2 and Si/Ti ≈ 1.5), and HfO 2 as well as oxidised Nb ss and Nb 5 Si 3 (Tables 5 and 6, Figure 7b). The substrate just below the scale was also cracked with cracks running parallel and perpendicular to the surface of the specimen. In some parts of the cross section, about 50 µm below the oxide scale, there was separation of the subscale alloy from the bulk (Figure 7a). Parts of the Ti-rich Nb ss in the subscale alloy were heavily oxidised (oxygen content in the range 30 to 50 at.%; see "peppered" darker contrast area on the right-hand side of analysis point 2 in Figure 7b), and the oxidised Nb ss exhibited cracks running parallel to the interface (Figure 7b).
At 800 • C, and only in very few areas along the scale/substrate interface, a very thin layer was observed that exhibited a very bright contrast under BSE imaging. The thickness of these layers was too small for quantitative analysis. Line scans across such layers confirmed that they were Sn rich (Figure 8). Analysis of such layers became possible after oxidation at 900 • C, where they became more continuous, and their thickness increased. The analyses showed that the bright contrast layer that formed at 900 • C consisted of Nb 3 Sn and Nb 6 Sn 5 (see Figure 9). The average composition of Nb 3 Sn was 30.8Nb-14.8Ti-4.4Si-3.9Al-3.7Hf-20.2Sn-0.7V-21.5O, whereas that of Nb 6 Sn 5 was 51.1Nb-1Ti-2.2Si-1.8Al-3.7Hf-39.9Sn-0.3V, i.e., both compounds were Cr free, and the Nb 6 Sn 5 was not contaminated with oxygen.  The bulk microstructure of the oxidised alloy NV1 consisted of the same phases that were observed in NV1-AC (Figure 7c,d). However, the Nb ss and the Nb 5 Si 3 and the lamellar Nb ss + Nb 5 Si 3 microstructure were contaminated with oxygen. In some parts of the microstructure, the lamellar microstructure had "disintegrated" (deteriorated and fragmented) owing to the heavy oxidation of its Nb ss (analysis point 4 in Figure 7c). In other words, the contamination of NV1 with oxygen during oxidation had progressed all the way to the bulk. However, in the bulk there were also Nb 5 Si 3 grains that were not contaminated with oxygen. In the contaminated Nb 5 Si 3 , the oxygen concentration varied from about 16 at.% near the oxide scale to about 7 at.% in the bulk of the specimen. On the other hand, the contamination of the Nb ss was most severe. The oxygen content of the Nb ss exceeded 40 at.% in areas up to about 100 µm below the oxide scale, and in the bulk it was about 30 at.%. The Ti-rich Nb ss and Ti-, Cr-, and V-rich Nb ss , which were still observed in the microstructure, exhibited "peppered" darker contrast, and the latter was stronger when the oxygen concentrations approached about 50 at.% (analyses points 2 and 3 in Figure 7c and d, respectively). However, the contamination with oxygen of the bulk of NV1 did not result in a further increase in the volume fraction of HfO 2 .

Hardness and Nano Indentation
The Vickers hardness of the alloy is given in Table 1. The hardness of NV1-HT and of Nbss in NV1-HT (see below) corresponds to the bulk microstructure (the area contaminated by nitrogen was removed prior to hardness measurements). The hardness decreased slightly after the heat treatment. The average Vickers microhardness of the Nbss in NV1-AC and NV1-HT, respectively, was 531 HV (516-545) and 540 HV (529-562), where, in the parentheses, the minimum and maximum microhardness values are given.
The large vol.% of Nbss in NV1 and its characteristic microstructure in NV1-HT (Figure 2d), and in NV1-AC in which in some areas it consisted of Nb5Si3 grains between Nbss grains without a lamellar microstructure (Figure 2b), offered us the opportunity to use nanoindentation to measure the hardness and elastic moduli of said phases in NV1-AC ( Figure 10) and NV1-HT ( Figure 11). Part (a) in each figure is a 30 × 30 μm 2 AFM image of the indentation area and shows the nanoindentations running from left to right. In parts (b) and (c) in each figure, every point represents a particular nanoindentation. The first indentation on the left-hand side of (b) and (c) in each figure is number 1. In both figures the indentation number 6 is located in Nb5Si3, and the indentation number 5 in Figure 10 is on the "interface" between Nbss and Nb5Si3, whereas the indentation numbers 5 and 7 in Figure 11 are on said interface. In both conditions, there was an increase in the modulus at the interface. Additionally, the modulus of Nb5Si3 was significantly higher than that of Nbss. In NV1-AC, there was hardly a change in hardness at the interface, whereas in NV1-HT, a small increase was indicated by the indentation number 7. The hardness and moduli of the Nbss and Nb5Si3 for the as-cast and heat-treated conditions are given in Tables 7 and  8.

Hardness and Nano Indentation
The Vickers hardness of the alloy is given in Table 1. The hardness of NV1-HT and of Nb ss in NV1-HT (see below) corresponds to the bulk microstructure (the area contaminated by nitrogen was removed prior to hardness measurements). The hardness decreased slightly after the heat treatment. The average Vickers microhardness of the Nb ss in NV1-AC and NV1-HT, respectively, was 531 HV (516-545) and 540 HV (529-562), where, in the parentheses, the minimum and maximum microhardness values are given.
The large vol.% of Nb ss in NV1 and its characteristic microstructure in NV1-HT (Figure 2d), and in NV1-AC in which in some areas it consisted of Nb 5 Si 3 grains between Nb ss grains without a lamellar microstructure (Figure 2b), offered us the opportunity to use nanoindentation to measure the hardness and elastic moduli of said phases in NV1-AC ( Figure 10) and NV1-HT ( Figure 11). Part (a) in each figure is a 30 × 30 µm 2 AFM image of the indentation area and shows the nanoindentations running from left to right. In parts (b) and (c) in each figure, every point represents a particular nanoindentation. The first indentation on the left-hand side of (b) and (c) in each figure is number 1. In both figures the indentation number 6 is located in Nb 5 Si 3 , and the indentation number 5 in Figure 10 is on the "interface" between Nb ss and Nb 5 Si 3 , whereas the indentation numbers 5 and 7 in Figure 11 are on said interface. In both conditions, there was an increase in the modulus at the interface. Additionally, the modulus of Nb 5 Si 3 was significantly higher than that of Nb ss . In NV1-AC, there was hardly a change in hardness at the interface, whereas in NV1-HT, a small increase was indicated by the indentation number 7. The hardness and moduli of the Nb ss and Nb 5 Si 3 for the as-cast and heat-treated conditions are given in Tables 7 and 8.

Microstructures
The solidification of NV1-AC started with the Nb ss as the primary phase. The formation of Nb 3 Si was suppressed owing to the destabilising effect of Sn [39] or Al [36] on Nb 3 Si in the presence of Ti, and the destabilising effect of Sn on Nb 3 Si when it is in synergy with Hf or Al (alloys EZ1 and EZ7 in [28]) or in synergy with Hf and Al or Cr (alloys EZ3 and EZ4 in [28]), even though Nb 3 Si can form when Hf is in synergy with Cr in the absence of Ti (alloy YG1 in [40]) or in synergy with Ti in the absence of Al, Cr, or Sn ( [33] and alloy YG3 in [40]). Furthermore, V forms cubic V 3 Si (cP8, A15, and prototype Cr 3 Si) [24] that is isomorphous with Nb 3 Sn and A15-Nb 3 Si [41], and thus the presence of V in NV1 would be expected to destabilise on tP32 Nb 3 Si. No A15-Nb 3 X (X = Al,Si,Sn) was observed in NV1 owing to the low Sn concentration [14].
When the stable (tP32) Nb 3 Si is suppressed in the Nb-Si binary under RS conditions, the Nb ss + Nb 3 Si stable eutectic is replaced by the Nb ss + βNb 5 Si 3 metastable eutectic with Si content about 20 at.% [42]. However, the latter concentration depends on the version of the Nb-Si binary used to define the composition of the stable eutectic [43], which can be in the range 15.3 to 18.7 at.% Si [17,43]. In RM(Nb)ICs, the eutectics with Nb ss and Nb 5 Si 3 have <Si> in the range 21.6 to 24.3 at.% (<Si> = Al + Ge + Si + Sn, [43]). In the case of NV1-AC, the average <Si> content of the lamellar micro-structure was 25.1 at.% (Table 2) or 24.9 at.% for the five lamellar microstructures shown in Figure 3.
In RM(Nb)ICs, the type of Nb 5 Si 3 (meaning tetragonal α or β) and the completion or not of the transformation βNb 5 Si 3 →αNb 5 Si 3 depends on the size of the as-cast button/ingot, meaning in "large" button/ingots the above transformation can occur during solid state cooling of the as-cast alloy (see Table A3 in [28] and alloy CM1 in [27]), the method of casting (see alloy CM1 in [27]) and the DS conditions ( [33] and alloy CM1 in [27]), the heat treatment temperature and duration of treatment (see Table A3 in [28]), and the alloying addition(s). For example, the 5-3 silicides of Cr and V are isomorphous with βNb 5 Si 3 [24]; the addition of Sn in RM(Nb)ICs with/without Ti quickens but does not lead to the completion of the above transformation in as-cast large button/ingots [28,39], and the same is the case when Hf is in synergy with Al, Cr, and Ti in RM(Nb)ICs (see Table A3 in [28]). On the other hand, the synergy of Al with Ti favours the βNb 5 Si 3 in as-cast large button/ingots, and the above transformation is completed after prolonged heat treatment at 1500 • C [36], whereas when Al is in synergy with Cr and Ti, said transformation is not completed for the same heat treatment conditions (see Table A3 in [28]). When tetragonal Nb 3 Si is formed, the αNb 5 Si 3 forms from the eutectoid transformation Nb 3 Si→Nb + αNb 5 Si 3 , which is very sluggish.
In binary Nb-Si alloys, the hexagonal γNb 5 Si 3 is metastable and can be stabilised by interstitials [2]. In RM(Nb)ICs, the formation/stability of γNb 5 Si 3 depends (a) on the Hf and Ti concentrations and the Nb/(Ti + Hf) ratio in Sn-free RM(Nb)ICs [33]; (b) on the synergy of Hf with other alloying additions, in particular Al and Sn, in Ti-free RM(Nb)ICs (e.g., alloy EZ4 in [28]); and (c) on the synergy of Hf and Sn with Al and/or Cr in Ticontaining RM(Nb)ICs [29]. Nb 5 Si 3 with Nb/(Ti + Hf) < 1 is most likely to have a hexagonal structure in RM(Nb)ICs with small Nb/(Ti + Hf) ratio, and this likelihood increases as the Nb/(Ti + Hf) ratio gets smaller [33]. In the alloy NV1-AC, the EPMA analyses confirmed the presence of Nb 5 Si 3 , which according to XRD had tetragonal (αNb 5 Si 3 ) and hexagonal structures, and the EPMA analysis data indicated that the ratio Nb/(Ti + Hf) was less or greater than one (Figures 2b and 3, Table 2). The αNb 5 Si 3 and γNb 5 Si 3 were also reported in Sn-free DS Nb-Ti-Si-Hf RM(Nb)ICs with 0.85 < Nb/(Ti + Hf) < 1.95 [33]. The alloy NV1-AC had Nb/(Ti + Hf) = 1.81. Figure 3 shows that the lamellar microstructure was "connected" with ("grew" from) Nb 5 Si 3 that had Nb/(Ti + Hf) less than one, i.e., hexagonal γNb 5 Si 3 . Figures 2b and 3d provide unconvincing evidence that a lamellar microstructure with brighter contrast, similar to that of hexagonal γNb 5 Si 3 , is also associated with Nb 5 Si 3 that has Nb/(Ti + Hf) greater than one, i.e., tetragonal αNb 5 Si 3 (see brighter contrast area between silicides 2 and 3 and above and to the right of silicide 3 in Figure 2b, owing to the partitioning of Hf and Ti in Nb 5 Si 3 , and the slightly darker contrast area in the silicide 6 in Figure 3c from which "grew" the lamellar microstructure between silicides 6 and 7). Note that the lamellar microstructure was not stable in NV1-HT. Was the lamellar microstructure that was observed in NV1-AC a eutectic, a eutectoid, or a combination of eutectic and eutectoid microstructure?
In NV1-AC, owing to their average chemical compositions, both the Ti-, Cr-, and V-rich Nb ss and the lamellar microstructure had a composition that corresponds to RCCAs (i.e., they were "RCCA phases") but not the Nb ss and Ti-rich Nb ss . In other words, because of the solidification conditions and the partitioning of solutes' "conventional" phases, i.e., Nb ss and Ti-rich Nb ss , and "RCCA phases" co-existed in the aforementioned solidification paths. However, the latter phases were not stable (not observed in NV1-HT). A similar co-existence of "conventional" and RHEA alloys was reported in the arc-melted RHEA of nominal composition (at.%) 13Nb-24Ti-24Si-35Al-4Hf [44]. The formation of bcc Nb ss and eutectics with Nb ss and Nb 5 Si 3 that satisfy the "standard definition" of HEAs is not rare in RM(Nb)ICs [17,43].
In NV1-AC, the average Si solubility in Nb ss and in the Ti-rich Nb ss was similar to that reported (i) for alloys of the Nb-Ti-Si and Nb-Hf-Si systems by Bewlay et al. [45,46], (ii) for the Nb-24Ti-18Si-5Hf alloy (YG3 in [40]), and (iii) for alloys of the Nb-Ti-Si-Cr-Al system such as KZ3, KZ4, KZ5, and KZ7 (often referred to by our group as the KZ series of alloys, see [35,36]), but it was higher than that in the Nb ss in the Nb-24Ti-18Si-5Sn (alloy NV6 in [39]) and in the alloy KZ2 with higher Cr concentration [36]. The increase in the solubility of Si in the solid solution with increasing Hf concentration is supported by the work of Bewlay et al. [31]. The solubility of Si in the Nb ss was not sensitive to the V concentration, which is in agreement with [47].
The average Ti concentrations in Nb ss and in the Ti-rich Nb ss were lower than those reported in [35,45] for alloys of the Nb-Ti-Si system, for alloys of the Nb-Ti-Si-Al-Cr system [35,36], for the alloy YG3 [40], and for the alloy NV6 [39]. This difference was more significant for the Nb ss . The presence of V led to a further increase in the average Ti concentration in parts of the Nb ss where the solubilities of Cr, V, and Hf had also increased. In some of these parts, the solubility of Ti was similar to that reported for alloys of the Nb-Ti-Si and Nb-Ti-Si-Cr-Al systems [35,36,45] and alloy NV6 [39]. The increase in the solid solubilities of Al, Cr, and Hf in the Nb ss with increased Ti concentration is in agreement with the results for the KZ series of alloys [35,36]; the alloys JG2, JG3, JG4, and JG6 with/without Hf (JG series of alloys) [6,48,49]; and the alloy YG3 [40]. It was concluded (i) that the solubilities of Al, Cr, and Hf in the Nb ss increased with increasing Ti concentration in agreement with previous research and (ii) that the solubility of V in the Nb ss also increased with the Ti content.
In NV1-AC, the solubility range of Hf in the Hf rich Nb 5 Si 3 was in agreement with the data in [32,46,50] but was significantly higher than the data in [49] for the Hf-containing and Hf-and-Sn-containing alloys JG4 and JG6, respectively. Furthermore, the solubility of Hf in the Ti-rich Nb 5 Si 3 in NV1-AC was in agreement with data for the Hf-rich Nb 5 Si 3 in JG4 and JG6 [49]. The solubility of Sn in Nb 5 Si 3 was significantly lower than the data for alloys JG6 [49] and NV6 [39], and the same was the case regarding the solubility of Cr, compared to the KZ series of alloys [35,36] and the JG series of alloys [48,49]. The solubility of V in Nb 5 Si 3 did not depend on the Ti or Hf solubilities and was significantly higher than the solubility of Mo [48] but significantly lower than the solubility of Ta in Nb 5 Si 3 [36]. However, the highest concentration values for V in Nb 5 Si 3 were significantly lower than the data reported in [47]. Comparison of the solubilities of V in the Nb ss and Nb 5 Si 3 would suggest that V predominantly partitions to the former phase, which is also supported by the data for the aforementioned phases in NV1-HT and is in agreement with [47]. Finally, the solubility of Al in Nb 5 Si 3 was the same as for the KZ and JG series of alloys [35,36,48,49].
The average Si and Sn concentrations of the lamellar microstructure in NV1-AC were higher and lower than the corresponding values for the Nb ss + Nb 5 Si 3 eutectics in NV9 and NV6 [39], respectively; within the range of the Si concentrations suggested for the Nb ss + Nb 5 Si 3 eutectic in the Nb-V-Si system [47]; and higher than that of the Nb-Si eutectic in the binary system [51] and the eutectics in JG series of alloys [6]. The average Hf concentration of the eutectic was the same as for the Nb ss and Nb 5 Si 3 eutectic in the alloy YG3 [40], but the Si concentration was lower.
The suppression of Nb 3 Si resulted to equilibrium between Nb ss and γNb 5 Si 3 in the bulk and between Nb ss and αNb 5 Si 3 in the near-the-surface areas of NV1-HT, according to the XRD and EPMA data (Tables 2 and 3 and Figure 1b). In NV1-HT, the Si solubility in Nb ss was reduced compared with NV1-AC but was still higher compared with the heat-treated KZ series of alloys [35,36], the JG series of alloys [49], and the alloy NV6 [39]. Compared with NV1-AC, the solubility of V in Nb ss and Nb 5 Si 3 was reduced; the solubility of Cr in Nb 5 Si 3 was increased; and the average solubility of Hf in Nb 5 Si 3 was the same as in the Hf-rich Nb 5 Si 3 in NV1-AC but with a significantly lower Ti concentration in the Nb 5 Si 3 .
The lamellar microstructure was not observed in NV1-HT. Given that the prior eutectic microstructure was stable in heat-treated Nb-V-Si alloys [47] and not completely eliminated in NV6-HT [39] and that the prior eutectic microstructure was not observed in the heattreated KZ and JG series of alloys [35,36,48,49], it is suggested that it is the synergy of Cr and Al that has the strong detrimental effect on the stability of Nb ss + Nb 5 Si 3 lamellar microstructures. The formation of HfO 2 in the as-cast alloy and of TiN in the near-thesurface areas of the heat-treated alloy demonstrated the sensitivity of NV1 to interstitial contamination, which was also confirmed by the oxidation results.

Oxidation
Compared with the alloy ZX7 [14], which had the same Sn concentration as NV1 but which was V-free, the addition of the latter element and Hf, and the lowering of the Si concentration in NV1, had an effect on (i) the chemical composition of the scale, (ii) the thickness of the scale, (iii) the enrichment of the subscale area of the substrate with Sn, and (iv) the formation of Sn rich layer at the scale/substrate interface. The oxides in the scale of NV1 at 900 • C were the same as those reported for Sn-free and for Sn-containing RM(Nb)ICs [13][14][15][17][18][19]. The oxygen content of the scale of NV1 was in agreement with the Si-rich oxide in the alloy ZX7 and (the richer in Sn) alloy ZX8 and the Nb-and-Ti-rich oxide in the alloy ZX8 [14,15]. However, the Cr and Sn concentrations in the scale of NV1 were remarkably low compared with the alloy ZX7 [14]. Additionally, in the scale, the V content was markedly low but hafnia was present in it ( Table 5). The scale of NV1 was thinner than that of ZX7 (≤20 µm and 40 µm, respectively), but unlike the latter its thickness was not uniform.
Contrary to the alloy ZX7 [14], in which formation of an Sn-rich layer at the scale/substrate interface was not observed at 800 • C, a thin Sn-rich layer was observed in the alloy NV1 at 800 • C, the thickness of which increased at 900 • C (Figure 7b, Figure 8, and Figure 9), where it consisted of the Nb 3 Sn and Nb 6 Sn 5 compounds. Note (i) that these two compounds can be at equilibrium below about 916 • C [24], (ii) that the melting temperature of Nb 6 Sn 5 is about 916 • C, (iii) that the oxidation of NV1 at T ≥ 1000 • C was catastrophic (results not reported in this study), and (iv) that A15-Nb 3 X (X = Al,Si,Sn) was not observed in NV1-AC and was not stable in NV1-HT. In other words, in the alloy NV1, similarly to the (richer in Sn) alloy ZX8 [15], an Sn-rich layer formed at said interface. Of the two compounds that made different parts of this layer (Figure 9), the Nb 3 Sn has been reported in richer in Si and Sn but V-free RM(Nb)ICs with/without Hf, in which the compounds Nb 5 Sn 2 Si and NbSn 2 also have been observed [13,15], but not the Nb 6 Sn 5 . Remarkably, in NV1, the latter compound was not contaminated with oxygen compared with the Nb 3 Sn, the contamination of which was more severe than in the alloys JG6 and ZX8 [13,15]. Furthermore, in NV1, both the Nb 3 Sn and Nb 6 Sn 5 were Cr free, and even though their V concentration was very low, it was double in the former than in the latter compound.
The consumption of solute additions in NV1 for the formation of (a) the scale, in particular Al, Hf, Si, and Ti; and (b) the Sn-rich layer, especially Al, Si, Sn, and Ti, resulted in the bulk composition at 800 • C being poorer in these elements compared with NV1-AC. Composition changes from the bulk to the surface of NV1 at 800 • C were expected owing (i) to the tendency of Al, Cr, Si, Sn, and Ti to segregate to the surface [14] and (ii) the presence of Ti-rich Nb 5 Si 3 , Ti nitride, and hafnia in the near surface areas of NV1-HT (Table 3). In Figure 7b, the Nb 5 Si 3 grains shown with the analysis numbers 1 and 2 were not contaminated with oxygen, and their composition corresponded to hexagonal γNb 5 Si 3 ((Nb/(Ti + Hf) equal to 0.95 and 0.78, respectively, analysis number 1: 29.1Nb-20.1Ti-35.1Si-3.5Al-0Cr-10.6Hf-0Sn-1.6V and analysis number 2: 25.8Nb-21.2Ti-35.6Si-3.2Al-0.5Cr-11.8Hf-0Sn-1.9V), whereas the Nb 5 Si 3 grains shown with the analysis numbers 3 and 11 were also hexagonal γNb 5 Si 3 ((Nb/(Ti + Hf) equal to 0.94 and 0.8, respectively) and contaminated, the latter more severely owing to its proximity with the scale/substrate interface (analysis number 3: 26.5Nb-18.6Ti-33.4Si-3.1Al-0.4Cr-9.7Hf-0.3Sn-1.3V-6.7O, analysis number 11:22.2Nb-18.3Ti-29.5Si-2.6Al-0.4Cr-9.3Hf-0.3Sn-1.5V-15.9O). Further in from the scale/substrate interface and toward the bulk, the contamination of Nb 5 Si 3 was reduced, and there was oxygen-free (i.e., non-contaminated) Nb 5 Si 3 in the bulk ( Table 5). All the Nb ss grains shown with the analysis numbers 4 to 9 in Figure 7b were contaminated with oxygen, and the chemical composition of the grains shown with the numbers 4 to 8 was similar (average composition 35.4Nb-14.8Ti-1.4Si-2.7Al-1.4Cr-2.6Hf-1.5Sn-2.9V-37.3O), whereas the composition of grain 9, below the Sn-rich layer, was slightly richer in Ti and V (31.4Nb-17.8Ti-1.1Si-3.5Al-2.1Cr-2.9Hf-1.8Sn-4V-35.4O). The contamination of the microstructure with oxygen from the scale/substrate interface to the bulk of NV1 followed the Ti-rich and Ti-, Cr-, and V-rich areas of Nb ss grains on either side of Nb 5 Si 3 grains (Figures 7b-d and 9) and the Nb ss in the lamellar microstructure (Figure 7c,d). In the latter, the more severe contamination of the Nb ss resulted in "disintegration" (break-up and degeneration) of the lamellae (Figure 7c,d).
In spite of the fact that its microstructure was contaminated with oxygen from the scale/substrate interface to the bulk, and its very high vol.% Nb ss , which played the key role in the contamination, the oxidation behaviour of the alloy NV1 in the pest regime was remarkable. Indeed, compared with other Sn-containing alloys with the same (alloy ZX7 [14]) or higher Sn content (alloys ZX8 [15] and JG6 [13] and the NV series alloys in Table 4) and the MASC alloy, the alloy NV1 (a) did not suffer from catastrophic pest oxidation, (b) had a mass change at 800 • C significantly lower than those of the alloys ZX7 and ZX8 (about 60 and 40 mg/cm 2 , respectively [15]), (c) had oxidation rate constants at 800 • C similar to those of the alloys ZX7 [14] and JG6 [13], and (d) had less mass change at 900 • C than all the other alloys in the Table 4, including the alloy NV5 that had the lowest vol.% Nb ss (1.9%). However, compared with the alloy NV5, at 800 • C, the parabolic rate constant of NV1 was two orders of magnitude higher and its mass change 14 times higher.
Concerning the oxidation of RM(Nb)ICs in the pest regime, the results of this work (a) underpinned the benefits of alloying with Sn, (b) emphasised the key role of Nb ss for the contamination of the alloy, (c) provided new data that show that the chemical composition of Nb ss is as important as its vol.% for the oxidation behaviour of RM(Nb)ICs, and (d) showed that a 2 at.% Sn addition might be good enough for oxidation resistance purposes, depending on the other alloying additions and their concentrations. It is suggested that (a) to (d) are relevant to other multiphase metallic UHTMs with Nb addition, namely, RHEAs and RCCAs [10,11].
How does the isothermal oxidation of NV1 in the pest regime temperatures compare with the RCCAs in the review in [52]? At 700 • C, the mass change of NV1 (3.5 mg/cm 2 ) was lower than those of the single phase bcc solid solution RRCAs Al 0.3 HfNbTaTiZr, Al 0.5 HfNbTaTiZr, Al 0.75 HfNbTaTiZr, AlHfNbTaTiZr, and HfNbTaTiZr (14,14,11,10, and 55 mg/cm 2 after 100 h, respectively). At 900 • C, the mass change of NV1 (30 mg/cm 2 ) was lower than HfNbTaTiZr (54 mg/cm 2 ), the same as Al 0.3 HfNbTaTiZr, but higher than Al 0.5 HfNbTaTiZr (18 mg/cm 2 ), Al 0.75 HfNbTaTiZr (17 mg/cm 2 ), and AlHfNbTaTiZr (16 mg/cm 2 ). In other words, at 700 • C, the mass change of the alloy NV1 was less than all the aforementioned richer in Al, Ti, and Hf RCCAs, whereas, at 900 • C, the mass gain of the latter was less than NV1 only when their Al, Hf, or Ti concentrations were in the ranges 9.1 < Al < 16.7 at.% and 16.7 < Hf or Ti < 18.9 at.%.

Hardness and Nanoindentation
In nanoindentation, from the unloading curve, the stiffness, S, of the phase can be measured. The stiffness is correlated with the reduced modulus E r with the equation: where P is the load, h is displacement, and A is the projected surface area of the indentation. The reduced modulus E r accounts for the effects of a non-rigid indenter during loading and is given by the equation: where E s and ν s are the Young's modulus and Poisson's ratio of the phase, respectively, and E i , ν i are the parameters for the indenter [53]. A rearrangement of the last equation gives the actual modulus, E s , of the phase as: The values of E i and ν i were specified in the TriboScope manual [34] as 1140 GPa and 0.07, respectively. The E s was calculated for three different values of ν s , namely, 0.38, 0.27, and 0.2 [54,55] (Figures 10b and 11b). The average E s values of the Nb ss and Nb 5 Si 3 given in the Table 7 were from the E s calculated with ν s equal to 0.38 and 0.27, respectively.
The Young's modulus of the Nb ss in NV1-HT (Table 7) was about 10 GPa higher than the average modulus of the Nb ss in the heat-treated Sn-free alloys KZ5, KZ6, and KZ7 [56,57] (131.1 GPa, range 115.7 to 138.6 GPa) and about 8 GPa lower than the average modulus of the Nb ss in the heat-treated Sn-free and Ge-containing alloys ZF4, ZF5, and ZF6 [58] (148.1 GPa, range 142.2 to 154 GPa), all calculated from nanoindentation data for the Nb ss with ν s = 0.38. Given that the aforementioned ZF series alloys were based on KZ series alloys with the addition of Ge, the data would suggest (a) that the alloying of the Nb ss with Ge or Sn increases the elastic modulus of Nb ss and (b) that the addition of Ge has a stronger effect (plus 17 GPa) than that of Sn (plus 10 GPa). The increase in the modulus of the Nb ss near the Nb ss /Nb 5 Si 3 interface in NV1-AC and NV1-HT should be noted in Table 7.
The Young's modulus of unalloyed (binary) Nb 5 Si 3 decreases in the sequence αNb 5 Si 3 > βNb 5 Si 3 > γNb 5 Si 3 according to ab initio calculations (291, 268.9, and 188.5 GPa, respectively [54]; also see Table 3 in [59]). The alloying of Nb 5 Si 3 with Ti increases the Young's modulus of α(Nb,Ti) 5 Si 3 and γ(Nb,Ti) 5 Si 3 and decreases the modulus of β(Nb,Ti) 5 Si 3 [54]. In RM(Nb)ICs, the Si can be substituted by other simple metal and metalloid element additions and Nb by other TM and RM additions (e.g., see Table 2 in [59]). To our knowledge, there is no data from experiments or calculations that shows how other (than Ti) alloying additions that substitute Nb or additions that substitute Si in Nb 5 Si 3 affect the Young's modulus of the silicide.
The hardness of Nb ss and Nb 5 Si 3 in RM(Nb)ICs that is measured using nanoindentation (nH) can be different from the microhardness (µH). The average correction factors derived from the data for the Nb ss and Nb 5 Si 3 in heat-treated KZ series alloys [56][57][58] and NV1-HT was 0.7357 for the former and 0.6395 for the latter; in other words, µH Nbss = 0.7357 × nH Nbss and µH Nb5Si3 = 0.6395 × nH Nb5Si3 . For the particular case of KZ7-HT, where only the α(Nb,Ti) 5 (Si,Al) 3 was present [36], the correction factor was 0.6275. Furthermore, the microhardness of binary αNb 5 Si 3 and alloyed α(Nb) 5 (Si) 3 was higher than that of the binary βNb 5 Si 3 and alloyed β(Nb) 5 (Si) 3 (see Table 4 in [59]). The microhardness of Nb ss in NV1 calculated from the nanoindentation data and the aforementioned correction factor was 528.2 and 523.1 HV for NV1-AC and NV1-HT, respectively, which was in good agreement with the measured average values, particularly the former. The microhardness of Nb 5 Si 3 was 1090 and 1340.3 HV for NV1-AC and NV1-HT, respectively. The former value was lower than the hardness of Nb 5 Si 3 in the alloy NV6-HT [39] and in other Ti-containing RM(Nb)ICs (see the Figure 3 in [59]) but close to the hardness of (Nb,Ti,Hf,Cr) 5 (Si,Al,Sn) 3 (1100 HV, [29]).
In Table 9, the measured average hardness of NV1 with the calculated hardness is compared. The calculations used the data for vol.% of phases in Table 1 (the vol.% hafnia is the balance), the microhardness values of the Nb ss and Nb 5 Si 3 calculated from the nanoindentation data, and 918 HV, the microhardness of hafnia. The calculated hardness is given for the law of mixtures, a Pythagorean-type addition rule, an inverse-type addition rule [28,39] and a cube-type addition rule (HV) 3 3 where V i is the volume fraction of phase i. The calculated values from the Pythagorean-type addition and the cubetype addition are closest to the measured values for NV1-AC and NV1-HT, respectively. The measured hardness of NV1-HT was lower than that of NV1-AC owing to the higher vol.% Nb ss and the lower vol.% of Nb 5 Si 3 compared with NV1-AC. How does the room temperature hardness of NV1-AC (475 HV, Table 1) and of the bcc Nb ss in the cast alloy (531 HV or 528 HV; the latter value was calculated from nanoindentation hardness, see above) compare with that of single-phase bcc solid solution as-cast RCCAs? Table 3 in [52] reports the hardness of 12 as-cast RCCAs. Five RCCAs had a hardness in the range of 298 to 454 HV, namely, (in increasing hardness) the alloys NbTaTiV, HfNbTaTiZr, MoNbTaTiV, NbTaTiVW, and MoNbTaW, whereas the hardness of six RCCAs, namely, NbTaVW, MoNbTaV, HfMoNbTaTiZr, MoNbTaVW, HfMoTaTiZr, and HfNbTaTiVZr were in the range 493 to 558 HV, and the hardness of CrMoNBTaVW was 705 HV. In other words, the hardness of NV1-AC was higher than that of the first five aforementioned RCCAs, and the hardness of the Nb ss in NV1-AC was lower only than four of the above mentioned RCCAs, namely, the alloys CrMoNbTaVW, HfNbTaTiVZr, HfMoTaTiZr, and MoNbTaVW. Note that the mass change of NV1 at 700 and 900 • C was also better than that of HfNbTaTiZr (see previous section).
The room temperature yield stress of NV1-AC calculated from hardness was 1552.7 MPa, and the specific yield stress was 210.4 MPacm 3 g −1 . The yield stress of NV1-AC was higher and lower than the yield stress of two RCCAs with a two-phase microstructure (bcc solid solution + M 5 Si 3 silicide)-the HfNbSi 0.5 TiV (1399 MPa) and HfMo 0.5 NbSi 0.3 TiV 0.5 (1617 MPa)-whereas, the specific yield stress of NV1-AC was higher than both RCCAs (179.8 and 191.1 MPacm 3 g −1 , respectively) [52]. Furthermore, the room temperature specific yield stress of NV1-AC was higher than most of the RCCAs reviewed in [52] (see Figure 6b and data in Table 2 in the ref. [52]).

Comparison with NICE
In Section 2, where the alloy design/selection was briefly discussed, the property goal and the constraints of the design were given together with the predicted values according to NICE. The latter overestimated MACSi (3.7 versus 3.1 at.%) and the mass change at 800 • C (14.8 versus 10 mg/cm 2 ), underestimated the vol.% of Nb ss (75.5 versus 81 at.%), and predicted correctly the stable phases (meaning the Nb ss and Nb 5 Si 3 and the absence of Nb 3 Si and A15-Nb 3 X) and the room temperature yield stress from hardness. Given the uncertainties in the design of metallic UHTMs and RM(Nb)ICs [10,11,17], the agreement of NICE with the experimental results was satisfactory.

Conclusions
We studied the microstructure, isothermal oxidation, and hardness of the Nb-23Ti-5Si-5Al-5Hf-5V-2Cr-2Sn alloy and the hardness and Young's moduli of elasticity of its Nb ss and Nb 5 Si 3 . There was macrosegregation of Ti and Si in the as-cast alloy. The microstructure consisted of the Nb ss , αNb 5 Si 3 , γNb 5 Si 3 , and HfO 2 phases in the as-cast or heat-treated alloy plus TiN in the near-the-surface areas of the latter. The vol.% of Nb ss was about 80%. There were Ti-and Ti-and-Hf-rich areas in the solid solution and 5-3 silicide, respectively, and a lamellar microstructure of these two phases in the cast alloy. The V partitioned to the Nb ss , where its solubility increased with increasing Ti concentration. At 700, 800, and 900 • C, the alloy did not suffer from catastrophic pest oxidation; it followed parabolic oxidation kinetics in the former two temperatures and linear oxidation kinetics in the latter, where its mass change was the lowest compared with other Sn-containing alloys. A thin Sn-rich layer formed in the interface between the scale and the substrate, which consisted of the Nb 3 Sn and Nb 6 Sn 5 compounds at 900 • C. The latter compound was not contaminated with oxygen. Both the Nb ss and Nb 5 Si 3 were contaminated with oxygen (the former more severely than the latter). Furthermore, the bulk of the alloy was also contaminated with oxygen. However, in the bulk, some Nb 5 Si 3 grains were not contaminated. The contamination of the microstructure with oxygen from the scale/substrate interface to the bulk followed the Ti-and Ti-, Cr-, and V-rich Nb ss grains on either side of Nb 5 Si 3 grains, and the Nb ss in the lamellar microstructure, where the more severe contamination of the Nb ss resulted in "disintegration" of the lamellae. The alloying of the Nb ss with Sn increased its elastic modulus compared with Sn-free solid solutions. The hardness of the alloy and its Nb ss and its specific room temperature strength compared favourably with many RCCAs. The research answered the three questions that motivated it, namely, that an RM(Nb)IC with very high vol.% Nb ss (i) can have acceptable oxidation in the pest regime and that it (ii) will be contaminated with oxygen from the surface areas to its bulk. It also confirmed that in such an alloy both the αNb 5 Si 3 and γNb 5 Si 3 can be stable. The agreement of the predictions of the alloy design methodology NICE with the experimental results was satisfactory.

Institutional Review Board Statement: Not applicable.
Informed Consent Statement: Not applicable.
Data Availability Statement: All the data for this paper is given in the paper, other data cannot be made available to the public.
Acknowledgments: Support of parts of this work by the University of Sheffield, Rolls-Royce Plc, EPSRC (EP/H500405/1, EP/L026678/1), and the University of Surrey is gratefully acknowledged.

Conflicts of Interest:
The authors declare no conflict of interest. * JG series alloys, + NV series alloys, ** KZ series alloys, and ++ ZF series alloys.