The role of Sn and Ti additions in the microstructure of Nb–18Si base alloys

Intermetallics 15 (2007) 1518e1528 www.elsevier.com/locate/intermet

The role of Sn and Ti additions in the microstructure of Nbe18Si base alloys N. Vellios a, P. Tsakiropoulos a,b,* b

a School of Engineering, University of Surrey, Guildford, Surrey GU2 5XH, UK IMMPETUS, Department of Engineering Materials, The University of Sheffield, Sir Robert Hadfield Building, Mappin Street, Sheffield, Yorkshire S1 3JD, UK

Received 7 February 2007; received in revised form 29 May 2007; accepted 5 June 2007 Available online 8 August 2007

Abstract The effects of Sn and Ti on the microstructure and hardness of the as cast and heat treated Nbe18Sie5Sn (NV9) and Nbe24Tie18Sie5Sn (NV6) alloys were studied. In both alloys the phases present in the as cast and heat treated microstructures were Nbss, Nb3Sn and Nb5Si3. In NV9, Sn suppressed the formation of Nb3Si, partitioned in Nbss stronger than in Nb5Si3 and did not affect significantly the solubility of Si in the Nbss. In NV6, the solubility of Ti in (Nb,Ti)ss increased in the presence of Sn, the concentration of Ti in Nb5Si3 was sensitive to cooling rate and the solubility of Sn in Nb5Si3 decreased as the concentration of Ti increased. The Ti controlled the partitioning of Si between (Nb,Ti)ss and Nb3Sn and was considered responsible for the macrosegregation of Si in the as cast ingot. The transformation of b to a Nb5Si3 was enhanced by the synergy of Sn and Ti. The addition of Ti did not destabilise the Nb3Sn. Silicon increased the hardness of Nb3Sn significantly, Sn did not affect the hardness of Nb5Si3 and Ti reduced the hardness of Nb3Sn and Nb5Si3 significantly. The hardness of NV9 and NV6 decreased and increased, respectively, by heat treatment. The reduction of the hardness of NV6-AC compared to NV9-AC is attributed to the strong effect of Ti on the hardness of Nb3Sn and Nb5Si3. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: A. Silicides, various; B. Phase transformations; B. Phase identification; C. Casting; D. Microstructure

1. Introduction New high temperature structural materials are required for the next generation of advanced aircraft engines. Many intermetallic compounds have been studied as candidate materials that could replace the conventional Ni superalloys. Niobium silicide base alloys (or Nb silicide base in situ composites) are very promising candidates for future application in the hottest airfoil parts [1,2]. To date, the most studied alloys belong to the NbeSieTie HfeCreAl system ([1e3] and Refs. within). These alloys exhibit excellent creep strength at elevated temperatures, but * Corresponding author. IMMPETUS, Department of Engineering Materials, The University of Sheffield, Sir Robert Hadfield Building, Mappin Street, Sheffield, Yorkshire S1 3JD, UK. Tel.: þ44 (0)1142225960; fax: þ44 (0)1142225943. E-mail address: [email protected] (P. Tsakiropoulos). 0966-9795/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2007.06.002

they lack oxidation resistance and suffer from poor fracture toughness at room temperature [2]. Tin has been considered as an alloying addition to improve oxidation resistance, in particular, to manage pest oxidation that has been eliminated in the temperature range 750e950  C [2,4]. However, Sn had a minimal effect on oxidation at a higher temperature (T ¼ 1200  C) [4]. The research described in this paper is part of a wider programme that aims to understand the effects of transition metal, refractory metal and free electron metal additions on phase selection and microstructure development in multi-component alloys based on Nbe18Si and the oxidation of these alloys [4e11]. In [5e7] the effects of Cr, Al and Ta were described while [8,10] have considered the effects of alloying with Mo, Hf and Sn. To date, the understanding of how Sn affects microstructural development and oxidation in Nb silicide base alloys is still poor. The motivation of this work was to study the

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effects of Sn and Ti on the microstructure of Nb silicide base alloys in the absence of Cr, Al, Hf and refractory metal additions. 2. Experimental The nominal compositions of the Nb silicide base alloys studied in this work are given in Table 1. The alloys NV9 and NV6 were designed to understand the effect of Sn and Ti additions on phase selection in Nbe18 Si in which the equilibrium phases at room temperature are solid solution and Nb5Si3 silicide [1]. To the authors’ knowledge, only Waterstrat and Mueller [12] have provided phase equilibrium data for the NbeSie Sn system in the form of an isothermal section at 1600  C according to which (i) the microstructure of Nbe18Sie5Sn would consist of Nbss, the body-centred tetragonal Nb3Sn (A15) phase and the Nb5Si3 silicide and (ii) the presence of Si causes a strong deviation from the ideal stoichiometry of the A15 phase with the concentrations of Sn and Si in the ranges 25e12 at% and 0e5 at%, respectively. The only other piece of work where the formation of Nb3Sn has been reported in NbeSi base alloys is the recent work of Geng et al. [4] who studied the oxidation of the Nbe24Tie18Sie5Cre5Ale5Hfe 2Moe5Sn alloy at 1200  C and reported the formation of Nb3Sn in the oxide/alloy interface. In this work, the Sn addition was set at 5 at% in order to study the effect of Sn in the presence or not of Ti on phase selection regarding the intermetallics Nb3Sn, Nb5Si3 and Nb3Si and the Nb solid solution. In alloy NV6 the Ti concentration was set at 24 at%, the same as for alloy KZ3 in [5]. Niobium (99.99 wt.%), Si (99.999 wt.%), Ti (99.99 wt.%) and Sn (99.99 wt.%) were used as the starting materials for alloy preparation. The alloys were melted in an argon atmosphere using a non-consumable tungsten electrode and a water-cooled copper crucible as described in [5]. Cubic specimens for heat treatments at 1200  C, 1500  C and 1600  C were wrapped in Ta foil, placed in an alumina boat and annealed under argon atmosphere in a tube furnace for 100 h or 10 h. Another boat, filled with titanium sponge that was used as a getter was placed at the entrance of the argon flow in the tube furnace. A Philips X-ray diffractometer with a monochromatic Cu ˚ ) radiation was used for the identification Ka (l ¼ 1.5418 A of the phases which was done using JCPDS data. A JEOL 8600 electron probe micro-analyser was used. Analysis was performed at 15 kV and at each measurement the probe diameter was adjusted so as to achieve dead time less than 20%. Under the voltage used the spatial resolution is about 1e1.5 mm. Thus, analysis was performed for phases of size Table 1 Nominal compositions (at%) of the alloys of this study Alloy

Si

Ti

Nb

Sn

NV9 NV6

18 18

e 24

77 53

5 5

1519

>5 mm diameter. At least 10 analyses were done for each phase in different parts (top, centre and bottom) of the ingot. The average concentrations of at least 10 analyses of an element in a phase or region of the ingot are given in Tables 3 and 4. Measurements of the area fraction of phases were performed using software available on the microprobe. Area fractions were measured for the same areas that were used for large area analyses. At least 10 measurements were taken for each alloy, all at the same magnification (350). In the alloy NV9 the area fraction of Nbss could not be measured. In the alloy NV6 it was not possible to measure the fractions of the 5-3 silicide with different concentrations of Ti. Thus, the fraction of the area occupied by the 5-3 silicide can be calculated from the data in Table 2. Vickers hardness of the alloys in the as cast and heat treated condition was measured with a load of 10 kg on a Vickers Instruments hardness machine. At least 10 measurements were taken for each alloy. The Vickers microhardness of phases present in the alloys was measured using a Mitutoyo hardness machine with a load of 0.05 kg. The density of the alloys was measured using the Archimedes’ principle and a Sartorious electronic precision balance equipped with a density determination kit. The density of the alloys is given in Table 2. In Tables 2, 3 and 4 the average values are given in bold. 3. Results 3.1. Nbe18Sie5Sn (alloy NV9) 3.1.1. As cast (NV9-AC) Data for the compositions and microstructures of the ingot and phases in NV9-AC are given in Figs. 1a and 2 and in Table 3. The data in Table 3 summarise the compositions of the phases in the whole ingot. Compared to the nominal composition (Table 1), the NV9-AC was poor in Sn and rich in Si with the actual concentrations of these elements being 4.2 and 19 at%, respectively. Large area analyses of different parts of the NV9-AC ingot showed no macrosegregation for Sn and Si. According to the XRD data (Fig. 1a) the Nb3Sn phase as well as the solid solution and the 5-3 silicide were present in NV9-AC. Thus, the EPMA data for NV9-AC has been summarised in Table 3 for a microstructure consisting of Nbss, Nb3Sn and Nb5Si3 silicide. In the areas close to the top of the ingot there was primary Nb5Si3, Nbss and Nb3Sn with a fine lamellar eutectic of Nb5Si3 silicide with Nbss. The eutectic was formed either adjacent to the 5-3 silicide or the Nb3Sn phase. The composition of the eutectic was 79.3Nbe18Sie 2.7Sn. In Nb3Sn phase the Si þ Sn concentration was w17.8 at% with Si/Sn z 1 while in Nbss the Si þ Sn concentration was w5.7 at% and the Si/Sn ratio was w0.3. The XRD data indicated the presence of both a and b Nb5Si3. The Nb3Sn phase (Fig. 2a and b) was found next to the blocky Nb5Si3 and in some cases surrounded the 5-3 silicide. There were no Sn rich areas in the Nb5Si3. Similar microstructure was observed in the centre of the ingot and in areas close to the bottom of the ingot. Fig. 2c shows an area near the bottom of the ingot where

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Table 2 Density and hardness data of the alloys and areas of phases Alloy

Density (g/cm3)

Hardness (Hv)

Area of Nb3Sn (%)

Area of Nbss (%)

Area of Nb5Si3 (%)

NV6eAC NV6eHT NV9eAC NV9eHT1c NV9eHT2c NV9eHT3c

6.95  0.03 e 7.9  0.01 e e e

638  22 683  26 814  27 761  22 675  19 687  26

20.9  3.3 25.0  2.6 16.9  1 17  2.5 17.2  4 17.1  3.5

33.5  1.5 29.5  1.8 39.8b 45b 40.3b 40.2b

45.6a 45.5a 43.3  1.8 38.0  1.6 42.5  1.9 42.7  1.9

a b c

Difference from 100%, see text. Area of eutectic/lamellar structure calculated as difference from 100%. See Table 3.

the Nb3Sn phase was not formed. In the latter areas the scale of the microstructure was significantly finer. The composition of the eutectic in the bottom of the ingot was the same as above. 3.1.2. Heat treated (NV9-HT) The alloy was given three separate heat treatments, namely 1200  C/100 h, 1500  C/100 h and 1600  C/10 h. Following the heat treatment at 1500  C for 100 h the Nb3Sn, Nb5Si3 and Nbss phases were present in the microstructure (Figs. 1b and 3b). In some areas the eutectic had almost been replaced by small islands of Nb5Si3 dispersed in the Nbss. In Nb3Sn, the Si þ Sn concentration was w17.2 at% with Si/Sn z 0.8. In Nbss the Si þ Sn concentration was w 3.4 at% with Si/ Sn z 0.1. There were also areas of Nbss with the composition 96.5Nbe0.6Sie2.9Sn that exhibited a darker grey contrast

compared to the majority of Nbss (Fig. 4). Furthermore, there was a small volumed fraction of the Nb3Sn phase that exhibited a brighter contrast under back scattered (BSE) imaging conditions owing to its higher Sn concentration (w13 at%). In Table 3, these regions are identified as Sn rich Nb3Sn. These regions were found between the Nb5Si3 islands and the Nbss (Fig. 4d) and had Si þ Sn z 18.2 at% and Si/Sn z 0.4. There was precipitation of fine Nbss particles in some 5-3 silicide grains. According to the XRD results (Fig. 1b) both the a and b Nb5Si3 were present in NV9-HT. A slight contrast between various areas of the silicide was observed in the BSE images (Fig. 4). In some of these silicide grains, a distinct boundary between the two areas of different contrast could be seen (Fig. 4). EPMA indicated that there were no Sn rich areas in the 5-3 silicide, as was the case in the as cast alloy.

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Fig. 1. X-ray diffractograms of (a) the as cast and (b)e(d) the heat treated NV9 alloy; (b) 1500 C/100 h, (c) 1200  C/100 h and (d) 1600  C/10 h.

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Fig. 2. BSE images of areas near (a) the surface (1000), (b) the centre (350) and (c) the bottom (350) of NV9-AC ingot.

The microstructure after the heat treatment at 1200  C for 100 h was similar to the microstructure described above for the 100 h heat treatment at 1500  C (Figs. 1c, 3a, b and Table 3). Both a and b Nb5Si3 were present in the microstructure. The eutectic microstructure seen in NV9-AC was still present, and had a finer scale compared to the one after the 1500  C heat treatment. In Nb3Sn the Si þ Sn concentration was w17.8 at% with Si/ Sn z 1 and in Nbss the Si þ Sn concentration was w2.4 at% with Si/Sn z 0.04. The Sn rich Nb3Sn that was present in NV9-HT (1500  C/100 h) was also present between Nb5Si3 and Nbss but the area fraction was significantly less. Similar to the two previous heat treatments, the microstructure after 10 h at 1600  C consisted of Nb5Si3, Nb3Sn and Nbss (Fig. 1d) and remaining eutectic of Nb5Si3 and Nbss (Fig. 3c). In Nb3Sn the Si þ Sn concentration was w17.3 at% with Si/ Sn z 0.8 and in Nbss the Si þ Sn concentration was w3.4 at% with Si/Sn z 0.03. In Sn rich Nb3Sn, the Sn concentration had increased further and Si þ Sn z 17.8 at% with Si/Sn z 0.5. Compared to the as cast alloy, the fraction of aNb5Si3 increased at the expense of bNb5Si3 after the heat treatments

at 1200, 1500 and 1600  C (Fig. 1). The area fraction of Nb5Si3 had not changed significantly after heat treatment at 1500 and 1600  C but was slightly reduced at 1200  C (see Table 2). 3.2. Nbe24Tie18Sie5Sn (alloy NV6) 3.2.1. As cast (NV6-AC) Data for the compositions of the phases and microstructure of the ingot of NV6-AC are given in Figs. 5a, 6aec and in Table 4. On the basis of the XRD data for NV6 (Fig. 5a, b) and the XRD data for NV9 (Fig. 1), it is concluded that the Nb3Sn phase was present in NV6. The data in Table 4 summarises the compositions of the phases in the whole ingot. Large area analyses of different parts of the NV6-AC ingot showed macrosegregation for Ti and Si. Indeed, between the top and bottom parts of the ingot there was a variation of the Si and Ti concentrations of 14.8e 20.6 at% and 21.8e30.2 at%, respectively. Overall, the microstructure of NV6-AC consisted of (Nb,Ti)ss, (Nb,Ti)3Sn phase and (Nb,Ti)5Si3 silicide. In the latter there were areas rich in Ti

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Table 3 EPMA data of the as cast and heat treated NV9 alloy Si (at%)

Nb (at%)

19.0  0.9 18.6  0.6 19.3  1.0 1.3  0.2 8.8  0.4 35.0  0.3

76.9  0.6 77.1  0.6 76.4  0.9 94.3  0.3 82.2  0.3 63.8  0.2

4.1  0.3 4.3  0.4 4.3  0.2 4.4  0.3 9.0  0.4 1.2  0.2

NV9-HT1 1200  C/100 h Bulk 18.1  0.1 Nbss 0.1 8.7  0.3 Nb3Sn Nb5Si3 35.2  0.2

77.6  0.1 97.6  0.2 82.2  0.2 63.4  0.2

4.3  0.2 2.3  0.2 9.1  0.3 1.4  0.1

NV9-HT2 1500  C/100 h Bulk 18.8  0.7 0.4  0.4 Nbss Nb3Sn 7.6  0.2 Sn rich Nb3Sn 5.3  0.1 35.4  0.3 Nb5Si3

77.0  0.7 96.6  0.6 82.8  0.3 81.8  0.1 63.3  0.1

4.2  0.2 3.0  0.2 9.6  0.4 12.9  0.2 1.3  0.2

NV9-HT3 1600  C/10 h Bulk Nbss Nb3Sn Sn rich Nb3Sn Nb5Si3

76.0  0.4 96.6  0.3 82.7  0.3 82.2  0.2 63.3  0.2

4.0  0.2 3.3  0.1 9.5  0.4 11.6  0.1 1.2  0.2

NV9-AC Surface Bulk Bottom Nbss Nb3Sn Nb5Si3

20.0  0.4 0.1 7.8  0.2 6.2  0.1 35.5  0.3

Sn (at%)

was surrounded by (Nb,Ti)ss and was found in the areas where the eutectic had developed. In (Nb,Ti)ss and (Nb,Ti)3Sn the Si þ Sn concentrations and the Si/Sn ratio were, respectively, w7.7 and 18.2 at% and w0.3 and 0.7 at%. 3.2.2. Heat treated (NV6-HT) The microstructure of the heat treated alloy consisted of (Nb,Ti)5Si3 of various sizes distributed unevenly in the solid solution (Fig. 6d). According to the XRD data (Fig. 5b) only the a (Nb,Ti)5Si3 was present, which would suggest that the transformation of b to a (Nb,Ti)5Si3 had been completed after 100 h at 1200  C. After this heat treatment there was still evidence of Ti rich areas in the 5-3 silicide. The solid solution and Nb3Sn had similar Ti concentrations. The area of the (Nb,Ti)ss had decreased from 33.5% to 29.5% and that of the (Nb,Ti)3Sn had increased from 20.9% to 25%, see Table 2. Furthermore, in (Nb,Ti)ss the Si concentration was low, and in agreement with previous work [5,6,8]. In the very Ti rich areas of the 5-3 silicide, the Ti concentration was also reduced (w32 at% as compared to w40 at% in the as cast alloy). In some areas there was still some of the eutectic microstructure seen in NV6-AC. In (Nb,Ti)ss the Si þ Sn concentration was w4.3 at% and the Si/Sn ratio was w0.08. In (Nb,Ti)3Sn the corresponding values were Si þ Sn z 16.3 at% and Si/Sn z 0.5. 3.3. Hardness

(Table 4). The (Nb,Ti)ss, regarding its high Ti concentration, corresponded to the Ti rich (Nb,Ti)ss reported in [5,6,8]. In the areas close to the top of the ingot there was primary (Nb,Ti)5Si3, (Nb,Ti)ss and (Nb,Ti)3Sn phase with the latter distributed unevenly between the (Nb,Ti)5Si3 dendrites and a fine lamellar eutectic of (Nb,Ti)5Si3 and (Nb,Ti)ss. The average composition of the eutectic was 49.8Nbe29.7Tie16.3Sie 4.2Sn. The XRD data (Fig. 5a) indicated the presence of both the a and b (Nb,Ti)5Si3. There was no evidence for the existence of Ti3Si or Ti5Si3. The (Nb,Ti)3Sn was often found next to the blocky (Nb,Ti)5Si3 as for NV9-AC (see Fig. 6a). There were Ti rich areas in the (Nb,Ti)5Si3, and these exhibited dark and strong dark contrast, see Fig. 6a. Thus, the 5-3 silicide is designated as Ti rich and very Ti rich (Nb,Ti)5Si3 in Table 4 with the Ti concentration in the latter varying from 29.8 to 37 at% at the top of the ingot. The Ti rich areas of the (Nb,Ti)ss surrounded the very Ti rich (Nb,Ti)5Si3 particles. Titanium rich areas were observed in the 5-3 silicide in the centre of the ingot and in areas close to the bottom of the ingot, but in the centre of the ingot the very Ti rich 5-3 silicide was not found. In the centre of the ingot the volume fraction of (Nb,Ti)3Sn increased, there were more blocky 5-3 silicides and fine eutectic of the 5-3 silicide and (Nb,Ti)ss, the latter often forming near the facets of the 5-3 silicides. The average composition of the eutectic was as above. The microstructure near the bottom of the ingot (Fig. 6c) consisted of a fine eutectic of (Nb,Ti)5Si3 and (Nb,Ti)ss, separated by (Nb,Ti)3Sn from blocky (Nb,Ti)5Si3 dendrites. The very Ti rich 5-3 silicide

The hardness values of the alloys in the as cast and heat treated conditions are given in Table 2. The microhardness values of phases present in the alloys are given in Table 5. 4. Discussion 4.1. Nbe18Sie5Sn (alloy NV9) According to [13,14], binary NbeSi alloys with 17  Si  18.6 at% would be hyper-eutectic alloys that would solidify with Nb3Si as the primary phase, while according to [15] some alloys in this composition range would be hypoeutectic with the Nbss as the primary phase. According to [13,15], binary NbeSi alloys with Si > 18.6 at% would solidify with bNb5Si3 as the primary phase and the Nb3Si would then form via the peritectic reaction L þ bNb5Si3 / Nb3Si. In the aforementioned alloys, as the temperature is decreased in equilibrium solidification, the Nbss and aNb5Si3 would form from the eutectoid transformation of Nb3Si. In practice, this reaction is very sluggish in binary NbeSi alloys [16]. In the NbeSn binary phase diagram reported by Shunk [17] the Nb3Sn compound is assumed to be unstable below 775  C. Other researchers [18e20] have considered the Nb3Sn phase to be stable at low temperatures and reported its existence in the temperature range 400e600  C. Toffolon et al. [21] have recently reassessed the NbeSn system considering the Nb3Sn phase to be stable below room temperature. In the microstructure of NV9-AC the Nb3Si was not observed, instead the Nb5Si3 silicide was present as the

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Fig. 3. BSE images (350) of the bulk of NV9-HT. (a) 1200  C/100 h, (b) 1500  C/100 h and (c) 1600  C/10 h.

primary phase. It is therefore concluded that the Sn addition had a destabilising effect on the 3-1 silicide in the solidification of Nbe18Sie5Sn. The Si concentration of the eutectic formed between Nb5Si3 and Nbss was 18 at%, slightly higher than the eutectic concentrations in [13,14] for NbeSi binaries and very close to the eutectic concentration in [15]. The concentration of Sn in Nb5Si3 did not change significantly between the as cast and heat treated conditions which would suggest that the Sn solubility in Nb5Si3 is w1.3 at%. The concentration of Si in the Nbss was reduced after heat treatment to 0.4 at%, in agreement with [5,6,8e10], while the concentration of Sn was reduced slightly. The data would suggest that the solubility of Sn in the Nbss of NV9 is w3 at%. In Nb3Sn there were no significant changes in the concentration of Si and Sn after heat treatment at 1200  C for 100 h, but there was a small reduction in the concentration of Si after the heat treatments at 1500 and 1600  C. The Si þ Sn concentration changed only slightly from w17.8 at% in NV9-AC to w17.2 at% in NV9-HT (1500  C) which would suggest that this is the equilibrium Si þ Sn concentration in Nb3Sn in the Nbe18Sie5Sn alloy.

In Nb5Si3 the Si þ Sn concentration was w36.2 at% in NV9AC and increased slightly after the heat treatments to 36.6 at% (at 1200  C) and 36.7 at% (at 1500 and 1600  C), very close to the stoichiometric composition in the NbeSi binary [14]. It is concluded that the equilibrium Si þ Sn concentration is 36.7 at% in the 5-3 silicide in the NbeSieSn system. The results for the Si þ Sn concentrations in Nbss, Nb3Sn and Nb5Si3 would suggest that the three-phase equilibrium between the aforementioned phases at the three temperatures (1200, 1500 and 1600  C) would not change significantly compared to the 1600  C section of the NbeSieSn by Waterstrat and Mueller [12]. However, equilibrium had not been reached in NV9-HT (at 1200, 1500 and 1600  C) as suggested by the formation of Sn rich Nb3Sn in the microstructure of the heat treated alloys. The solidification route of NV9-AC was L / L þ bNb5Si3 / L þ bNb5Si3 þ Nb3Sn / Nb5Si3 þ Nb3Sn þ (Nb5 Si3 þ Nbss)eutectic with bNb5Si3 as the primary phase and some bNb5Si3 / aNb5Si3 transformation occurring during the cooling of the ingot to room temperature, as suggested by the presence of peaks of both b and a Nb5Si3 in the XRD data (Fig. 1a). As the 5-3 primary phase was formed in the melt,

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Fig. 4. BSE images of the bulk of NV9-HT (1500  C/100 h) showing the Nbss with the average composition 96.5Nbe2.9Sne0.6 Si (a, c and d), the contrast exhibited by the Nb5Si3 (b and c) and Sn rich Nb3Sn (d).

the latter became leaner in Si but richer in Sn (owing to the rejection of Sn to the melt) near the Nb5Si3. Thus, in the areas near the primary 5-3 silicide the melt reached a composition with Si/Sn z 1 and the Nb3Sn was formed. As a result, the melt in front of the Nb3Sn became leaner in Si and Sn and the eutectic composition was reached leading to the eutectic reaction L / Nb5Si3 þ Nbss. This solidification route would be affected by high cooling rates (effect on rejection of Sn in the melt), which would be the case near the bottom of the ingot, where the area fraction of the Nb3Sn surrounding the Nb5Si3 primary phase had been reduced. There were also parts of the melt near the primary 5-3 silicide where the Si concentration reached the eutectic composition quickly and thus the eutectic reaction L / Nb5Si3 þ Nbss occurred first. 4.2. Nbe24Tie18Sie5Sn (alloy NV6) Compared to the alloy Nbe24Tie18Si (KZ3 in [5]), in which there was no macrosegregation of Ti and Si, the addition of Sn caused macrosegregation of Si and Ti in NV6AC. Compared to NV9, the addition of Ti caused

macrosegregation of Si in NV6-AC. Considering the effects of Sn and Ti on Nb3Si formation in NV9-AC and KZ3-AC, it is concluded that Sn has a very strong destabilising effect on the 3-1 silicide. In NV6-AC, the Ti concentration in (Nb,Ti)ss corresponded to that of the Ti rich (Nb,Ti)ss in KZ3, and the concentration of Sn had increased compared to NV9-AC. It is concluded that in the presence of Sn the solubility of Ti in the solid solution is increased. In (Nb,Ti)3Sn the Sn concentration had increased slightly compared to NV9. In NV6-HT, the concentration of Si in (Nb,Ti)ss was reduced to w0.3 at%, slightly lower than for the heat treated Nbe24Tie18Si, and in Nb3Sn the concentration of Si was also reduced slightly as for the Sn rich Nb3Sn in NV9-HT. Compared to NV9-HT the presence of Ti in NV6 had a significant effect on the solubility of Sn in (Nb,Ti)ss, which was reduced to 4 at%. It is concluded that (i) Ti controls the partitioning of Si between (Nb,Ti)ss and Nb3Sn, (ii) the solubility of Sn in Nbss increases with the solubility of Ti and (iii) Ti is responsible for the macrosegregation of Si in NV6-AC.

N. Vellios, P. Tsakiropoulos / Intermetallics 15 (2007) 1518e1528

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30

40

50

60

70

80

90

Degrees 2-Theta NV6-HT 1200 °C 3000

Nbss γ-Nb5Si3 α-Nb5Si3

b

2500

Nb3Sn + Specimen holder

Counts

2000 1500 1000 +

500

+

0 20

30

40

50

60

70

80

90

Degrees 2-Theta Fig. 5. X-ray diffractograms of (a) the as cast and (b) the heat treated NV6 alloy.

The microstructural data would suggest that the microstructures of NV9-HT (1600  C) and NV6-HT correspond to the same homologous temperature, i.e., 1473/TNV6 ¼ 1873/TNV9 L L NV6 NV9 or TL /TL z 0.8. In NV6 liquation was observed when the alloy was heat treated at 1500  C and this is attributed to the effect of Ti on the liquidus temperature TL of the alloy. Thus, it is assumed that TNV6 > 1500  C. Therefore, it is deL duced that the liquidus temperature of NV9 exceeds 1971  C (2216 K). XRD did not provide conclusive evidence for the presence of the 3-1 silicide. Furthermore, systematic examination of the microstructure of the alloy by EPMA did not reveal the presence of the 3-1 silicide. Thus, the (Nb,Ti)3Si that was formed in the as cast Nbe24Tie18Si (alloy KZ3 in [5]) was suppressed by the addition of Sn everywhere in the microstructure of the NV6 ingot and was replaced by the (Nb,Ti)5Si3 silicide (not present in KZ3 [5]), which exhibited Ti rich areas like the 3-1 silicide in the alloy KZ3. These areas were particularly evident in the top and bottom parts of the ingot where very Ti rich 5-3 silicide was observed. There were no Sn rich areas in the 5-3 silicide in which the Sn concentration was low, as for NV9. Considering the data for the alloys NV9-AC and NV6-AC and the alloy JG6 in [10], it is concluded that Sn controls the suppression of the 3-1 silicide in the solidification of Nb silicide base alloys (without additions of Fe, see [22]). The data for the 5-3 silicide would suggest that in the presence of Sn the solubility of Ti in the latter phase is sensitive to cooling

1525

rate and that as the concentration of Ti increases that of Sn decreases. According to the XRD data (Fig. 5) and compared to NV9-HT, in the presence of both Ti and Sn in the alloy, the transformation of b to a Nb5Si3 was enhanced. As the 5-3 primary phase was formed in the melt, the latter became leaner in Si but richer in Sn and Ti near the Nb5Si3. Thus, in the areas near the primary 5-3 silicide where the melt reached a composition with Sn þ Si concentration and Si/Sn ratio as in NV9-AC, the (Nb,Ti)3Sn phase was formed. Thus, the melt in front of the (Nb,Ti)3Sn became leaner in Si and Sn and the eutectic composition was reached leading to the eutectic reaction L / Nb5Si3 þ (Nb,Ti)ss. As for NV9, this solidification route would be affected by cooling rate, which was the case near the bottom of the ingot, where the volume fraction of (Nb,Ti)3Sn had been reduced. There were also parts of the melt near the primary 5-3 silicide where the Si concentration reached the eutectic composition quickly and thus the eutectic reaction L / Nb5Si3 þ (Nb,Ti)ss occurred first (Fig. 6b). Thus, in NV6-AC, the (Nb,Ti)3Sn was formed as Sn and Ti partitioned during solidification. Of these two elements the segregation of Ti played the dominant role. The equilibration of the Ti concentration after heat treatment was accompanied by a decrease of the volume fraction of (Nb,Ti)ss and an increase of the volume fraction of (Nb,Ti)3Sn (Table 2). It is suggested that the solidification path of NV6 was L / L þ b(Nb,Ti)5Si3 / L þ b(Nb,Ti)5Si3 þ (Nb,Ti)3Sn / (Nb, Ti)5Si3 þ (Nb,Ti)3Sn with a eutectic of the Nb5Si3 and (Nb, Ti)ss forming in the microstructure and some bNb5Si3 / aNb5Si3 transformation occurring during cooling.

4.3. Hardness The hardness data in Table 2 show that (i) the addition of Ti in NV6 was accompanied by a reduction of w176 Hv compared to NV9-AC, and (ii) the hardness of NV9 and NV6, respectively, decreased and increased after heat treatment. The hardness values of Nb3Sn and Nb5Si3 are, respectively, 450 Hv [23] and 1360 Hv [24]. The data in Table 5 would thus suggest that (i) the hardness of Nb3Sn is more than doubled by the addition of Si, (ii) the addition of Sn does not affect the hardness of Nb5Si3 significantly, and (iii) alloying with Ti causes a reduction of the hardness of Nb3Sn and Nb5Si3 by w33% and 27%, respectively, in the as cast microstructure. Thus, the significant reduction in the hardness of NV6-AC compared to NV9-AC is attributed to the strong effect that Ti has on the hardness of Nb3Sn and Nb5Si3. The data in Tables 3 and 5 would also suggest that the hardness of Nb3Sn increases as the solubility of Ti increases. Indeed, as the concentration of Ti in Nb3Sn increased from w20 at% to w28 at% in NV6-HT, the hardness of Nb3Sn increased by w17%. The hardness of (Nb,Ti)5Si3 in the heat treated microstructure also increased compared to NV6-AC but was still lower by w12 at% than the hardness of Nb5Si3. We attribute this change in the hardness of (Nb,Ti)5Si3 to a reduction in the Ti rich and very Ti rich areas in (Nb,Ti)5Si3 in NV6-HT.

N. Vellios, P. Tsakiropoulos / Intermetallics 15 (2007) 1518e1528

1526

Fig. 6. BSE images of different areas of (aec) the as cast and (d) the heat treated (1200  C/100 h) NV6 alloy; (a) near the surface of the ingot (1000), (b) near the centre of the ingot (1000), (c) near the bottom of the ingot (1000) and (d) centre (800).

Table 4 EPMA data of the as cast and heat treated NV6 alloy

NV6-AC Surface Bulk Bottom (Nb,Ti)ss (Nb,Ti)3Sn (Nb,Ti)5Si3 Ti-rich (Nb,Ti)5Si3 Very Ti-rich (Nb,Ti)5Si3 NV6-HT 1200  C/100 h Bulk (Nb,Ti)ss (Nb,Ti)3Sn (Nb,Ti)5Si3 Ti-rich (Nb,Ti)5Si3 Very Ti-rich (Nb,Ti)5Si3

Si (at%)

Ti (at%)

Nb (at%)

16.5  1.4 19.3  1.8 18.6  1.0 1.6  0.4 7.6  0.4 35.3  0.7 35.5  0.4 34.2  0.9

27.6  1.9 24.2  1.3 24.8  0.6 35.6  2.1 20.2  1.0 16.6  0.5 21.0  1.2 30.4  3.8

51.1  0.9 4.8  0.5 51.6  0.5 4.9  0.3 51.6  0.4 5  0.4 56.7  2.1 6.1  0.3 61.6  0.7 10.6  0.7 46.6  0.7 1.5  0.4 42.3  1.2 1.2  0.2 34.0  4.1 1.4  0.6

18.9  1.0 0.3  0.3 5.3  0.5 34.7  0.1 35.0  0.3 34.9  1.1

24.4  0.6 29.2  0.7 28.2  1.1 16.2  0.2 21.9  2.4 32.1  1.2

51.7  0.3 66.5  0.7 55.5  1.5 47.3  0.2 41.9  2.3 32.7  1.5

Sn (at%)

5.0  0.3 4.0  0.2 11  0.8 1.8  0.1 1.2  0.2 0.3  0.2

We have used the data in Tables 2 and 5 and the hardness values of 250 Hv and 230 Hv for the solid solutions, respectively, in NV6-AC and NV6-HT (the latter based on data for the hardness of Nb-TM binary alloys P [25]), and calculated the hardness of alloys using Hv ¼ Vi Hv1 (law of mixtures) Table 5 Microhardness of phases in alloys NV9 and NV6 and hardness of eutectic/ lamellar structure in NV9 Alloy þ condition

NV9-AC NV9-HT1b NV9-HT2 NV9-HT3 NV6-AC NV6-HT a b

See text. See Table 3.

Phase Nb5Si3

Nb3Sn

Eutectica

1351  30 1303  25 1305  56 1377  40 990  61 1198  40

943  65 889  40 930  31 919  34 636  31 745  66

681  36 647  47 545  56 609  28 e e

N. Vellios, P. Tsakiropoulos / Intermetallics 15 (2007) 1518e1528

1527

Table 6 Measured and calculated hardness values of alloys NV9 and NV6

Acknowledgements

Alloy þ condition

We would like to thank the EPSRC and Dstl for financial support under Grant GR/S81759/01. Collaboration with the Faraday Powder-Matrix partners HIP e Bodycote, Qinetiq, PSI and Rolls Royce is gratefully acknowledged.

NV9-AC NV9-HT1c NV9-HT2c NV9-HT3c NV6-AC NV6-HT

Hardness Measureda Calculatedb

814 761 675 687 638 683

A

B

C

(A þ B)/2 (B þ C)/2 (A þ C)/2

1014 937 933 992 668 779

661 594 614 659 478 580

921 849 803 868 470 504

834 766 774 826 573 690

791 722 709 764 474 542

968 893 868 930 569 642

A e law of mixtures, B e Pythagorean type addition rule, C e an inverse type addition rule. a Average value from Table 2. b See text. c See Table 3.

P 2 or Hv2 ¼ P ðVi Hvi Þ (Pythagorean type addition rule) or 1=Hv ¼ Vi =Hv1 (an inverse type addition rule), where Vi is the volume fraction of a phase and Hvi is its hardness. In Table 6 the data calculated using the three equations are given in the columns A, B and C, respectively. The calculated hardness values for the alloy NV9 are in good agreement with the measured values of NV9-AC and NV9-HT1 (hardness ¼ (A þ B)/2), and of NV9-HT2 and NV9-HT3 (hardness ¼ (B þ C)/2 and B, respectively). For the alloy NV6 the calculated hardness values are in good agreement with the measured values of NV6-AC (hardness ¼ A) and NV6-HT (hardness ¼ (A þ B)/2).

5. Conclusions We have studied the effects of Sn and Ti on the microstructure and hardness of as cast and heat treated Nbe18Sie5Sn (NV9) and Nbe24Tie18Sie5Sn (NV6) alloys. The conclusions of this work are as follows. The microstructure of both alloys in the as cast and heat treated conditions consisted of the Nbss, Nb3Sn and Nb5Si3 phases. In NV9, the Sn suppressed the formation of Nb3Si, did not affect the solubility of Si in the Nbss and partitioned in the Nbss stronger than in the Nb5Si3. In NV6, the Ti was responsible for the macrosegregation of Si and the synergy of Sn and Ti enhanced the transformation of bNb5Si3 to aNb5Si3. The solid solubility of Ti in (Nb,Ti)ss increased in the presence of Sn, and the Ti controlled the partitioning of Si between (Nb,Ti)ss and Nb3Sn and did not destabilise Nb3Sn in the alloy. The Sn suppressed the formation of (Nb,Ti)3Si, which is stabilised by Ti in Nbe24Tie18Si. The Si increased the hardness of Nb3Sn significantly, the Sn did not affect the hardness of Nb5Si3 and the Ti reduced significantly the hardness of Nb3Sn and Nb5Si3. The hardness of NV9 and NV6 decreased and increased, respectively, by heat treatment. The significant reduction of the hardness of NV6-AC compared to NV9-AC is attributed to the strong effect of Ti on the hardness of Nb3Sn and Nb5Si3.

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