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Article

Effect of Ta and W Additions on Microstructure and Mechanical Properties of Tilt-Cast Ti-45Al-5Nb-2C Alloy

by
Juraj Lapin
* and
Kateryna Kamyshnykova
Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Dúbravská cesta 9, 845 13 Bratislava, Slovakia
*
Author to whom correspondence should be addressed.
Metals 2021, 11(12), 2052; https://doi.org/10.3390/met11122052
Submission received: 22 November 2021 / Revised: 11 December 2021 / Accepted: 14 December 2021 / Published: 18 December 2021
(This article belongs to the Special Issue TiAl-Based Alloys and Their Applications)
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Abstract

:
The effect of Ta and W additions on microstructure and mechanical properties of tilt-cast Ti-45Al-5Nb-2C (at.%) alloy was investigated. Three alloys with nominal composition Ti-45Al-5Nb-2C-2X (in at.%), where X is Ta or W, were prepared by vacuum induction melting in graphite crucibles followed by tilt casting into graphite moulds. The microstructure of the tilt-cast alloys consists of the α2(Ti3Al) + γ(TiAl) lamellar grains, single γ phase, (Ti,Nb,X)2AlC particles with a small amount of (Ti,Nb,X)C, and β/B2 phase identified only in W containing alloy. The EDS analysis shows that Ta segregates into the carbide particles and reduces dissolution of Nb in both (Ti,Nb,Ta)C and (Ti,Nb,Ta)2AlC phases. The alloying with W reduces Nb content in both carbide phases and leads to stabilisation of β/B2 phase in the lamellar α2 + γ regions. The alloying with Ta and W does not affect the volume fraction of the carbide particles but influences their size and morphology. While the alloying with Ta and W has no significant effect on Vickers hardness and the indentation elastic modulus of the studied alloys, the addition of Ta affects the nanohardness and elastic modulus of the (Ti,Nb,Ta)2AlC phase. The addition of W significantly increases the Vickers microhardness of the lamellar α2 + γ regions.

1. Introduction

Lightweight alloys based on a TiAl-Nb system are of great interest for applications in aerospace, power engineering, and automotive industry due to their high melting point, low density, high specific strength, and good creep and oxidation resistance [1,2,3,4,5]. Nb is a common alloying element in numerous TiAl-based alloys with high solubility and different partition between γ(TiAl) and α2(Ti3Al) phases as well as different influences on the lattice parameters of both phases [6]. The effects of Nb on the microstructure, creep properties, oxidation resistance, and deformation mechanisms of TiAl-based alloys have been well studied [7,8,9,10,11]. Further expansion of the usability of TiAl-Nb alloys can be achieved by improving their properties through additional alloying elements such as Ta, W, and C.
Ta, as the slower diffuser than Nb in TiAl-based alloys, is considered to be the most effective solute atom to promote the formation of massive γm(TiAl) phase in air-cooling conditions that have a great opportunity to improve the properties of TiAl-Ta alloys [12,13]. Zhang et al. [14,15] have reported that the addition of 0.5–1 at.% Ta promotes the formation of ultrafine γ(TiAl) grains through the formation of massive γm phase during quenching from α (Ti-based solid solution with a hexagonal crystal structure) phase field and leads to an improvement of mechanical properties. Vojtěch et al. [16,17] have shown that the addition of Ta to TiAl improves both creep strength and high-temperature cyclic oxidation resistance compared to those of TiAl-Nb alloy. Lapin et al. [18,19] have reported that Ta retards the transformation of α2 lamellae to γ phase and τ(Ti3Al2Ta) precipitates and stabilises the lamellar α2 + γ + τ type of microstructure of TiAl-Ta alloy during long-term ageing at 750 °C up to 10,000 h or creep exposure at 700 °C up to 30,000 h.
The positive effect of W addition on mechanical properties of TiAl-W alloys has been attributed to a reduction in the dislocation mobility and pinning of lamellar α2 + γ interfaces by fine B2 (ordered Ti-based solid solution with a cubic crystal structure) precipitates [20]. The high capability of W to significantly improve creep resistance of TiAl-W alloys has been demonstrated by Beddoes et al. [20,21,22], Couret et al. [23], and Lapin and Nazmy [24,25,26]. However, a significant softening has been observed at temperatures of 900–1000 °C due to instability of the lamellar α2 + γ microstructure [27,28].
Alloying with C below its solubility limit contributes to the solid solution hardening of TiAl-Nb alloys. As shown for several TiAl-based alloys [29,30,31,32], cubic perovskite needle-like Ti3AlC carbide particles, which are thermodynamically stable at medium temperatures, enhance strength and creep resistance due to precipitation hardening. On the contrary, fine hexagonal plate-like Ti2AlC carbide particles, which are thermodynamically stable at higher temperatures, are less efficient to impede dislocation mobility [33,34,35]. It has been reported that the alloying with C above its solubility affects the solidification path, changes the primary solidification phase, and leads to the formation of coarse primary Ti2AlC particles during solidification of TiAl-Nb alloys [32,36,37,38,39,40,41,42].
Various methods such as powder metallurgy, mechanical alloying, vacuum induction melting, etc., were applied for the processing of TiAl-based alloys. Among them, vacuum induction melting (VIM) combined with precise casting is identified to be one of the most cost-effective processing routes for the production of near net shape components. Several ceramic materials have been evaluated for melting and casting of TiAl-based alloys but none of them were inert against their melts. VIM in relatively cheap graphite crucibles and casting into graphite moulds represents a cost-effective way for processing TiAl-based castings. This process leads to an increase in C in TiAl-based alloys. However, the content of C can be controlled by an appropriate selection and optimisation of processing parameters, such as heating rate, melt temperature, holding time at the melt temperature, arrangement of the charge in a graphite crucible, chemical composition of master alloys, and vacuum pressure [43,44,45,46,47,48].
Despite previous studies on microstructure and mechanical properties of carbon-containing TiAl-Nb alloys reinforced with coarse primary Ti2AlC particles [39,49,50,51], only very limited information has been published about the effect of Ta and W on their microstructure, solidification behaviour, and properties [52]. This study aims to investigate the effect of Ta and W additions on the tilt-cast microstructure and room temperature mechanical properties of Ti-45Al-5Nb-2C (at.%) alloy. The effect of alloying with Ta and W on the chemical composition of coexisting phases, solidification behaviour, elastic modulus, hardness, microhardness, and nanohardness is investigated and discussed.

2. Materials and Methods

Three alloys with nominal compositions Ti-45Al-5Nb-2C-2X (at.%), where X is Ta or W, were prepared from a master alloy with nominal composition Ti–43.5Al–1Nb (at.%). The cylindrical pieces of the master alloy with a diameter of 60 mm and height of 15 mm were mixed with additions of pure elements in the form: sponge of Ti, rectangular pieces of Al, Nb sheet, Ta chips, W powder, and TiC powder. The charge with a weight of 490 g was placed into a graphite crucible with an inner diameter of 65 mm and a length of 135 mm. The graphite crucible was put into a protective alumina-based crucible equipped with a pouring cup, which was connected to a cylindrical graphite mould with an inner diameter of 40 mm and length of 230 mm. Before the heating, the vacuum chamber of the induction melting furnace was evacuated to a vacuum pressure of 6.5 Pa and flushed with argon (purity 99.9995%) three times. After increasing the vacuum pressure to 1 × 104 Pa by a partial filling of the vacuum chamber with high purity argon, the charge was heated to a melt temperature of 1650 °C and held at this temperature for 30 s. The temperature of the melt was measured by a pyrometer. The tilt casting was carried out into a cold graphite mould under the vacuum pressure corresponding to that of VIM. The tranquil filling of the mould with minimum surface turbulence was achieved by a control of the tilting speed. The cylindrical tilt-cast samples were cut transversally into smaller pieces for metallographic observations and mechanical testing using an electric spark wire cutting machine.
Metallographic preparation of the samples consisted of standard grinding using abrasive papers, polishing on diamond pastes with various grain sizes up to 0.25 μm, and etching in a solution of 100 mL H2O, 6 mL HNO3, and 3 mL HF. Microstructure evaluation was performed by scanning electron microscopy (SEM) and scanning electron microscopy in back-scattered electron (BSE) mode using JEOL 7600F field emission scanning electron microscope. The X-ray diffraction (XRD) analysis of coexisting phases was carried out by a diffractometer Bruker D8 equipped with an X-ray tube with a rotating Cu anode operating at 12 kW. The chemical composition of the alloys was analysed by energy-dispersive spectrometry (EDS) calibrated using the standards (TiC, TaC, and Ti2AlC) for measurements of the composition of carbides. The average content of carbon in the tilt-cast alloys was measured by LECO CS844 elemental analyser. The average contents of oxygen and nitrogen were measured by a LECO ONH836 elemental analyser. Grain size, interlamellar spacing, volume fraction, and size of particles were measured by computerised image analysis using digitalised micrographs. The measured microstructural data were treated statistically.
Instrumented hardness measurements were performed by a universal testing machine. Instrumented Martens hardness measurements were carried out at an applied load of 298 N, holding time at the point of load application of 0.1 s, and rate of load application of 15 N/s. Vickers hardness measurements were performed at an applied load of 298 N and holding time at the point of load application of 10 s. Vickers microhardness measurements were carried at an applied load of 0.49 N and a dwell time of 10 s on polished and slightly etched sections of the tilt-cast samples. Instrumented nanoindentation measurements were carried out at an applied load of 0.01 N and holding time at the point of load application of 0.1 s on polished and slightly etched samples using a nanoindenter with Berkovich tip of the indenter.

3. Results

3.1. Chemical Composition

Table 1 summarises nominal and measured chemical compositions with standard errors of the measurements of the studied tilt-cast alloys. It is clear from Table 1 that the measured average contents of the alloying elements in the tilt-cast alloys designated as Nb, NbTa, and NbW correspond well to their designed nominal compositions. In addition, the applied cost-effective processing route consisting of VIM in graphite crucibles followed by tilt casting into cold graphite moulds leads to an acceptable increase in gases significantly affecting solidification behaviour [53,54,55,56]. The initial contents of oxygen and nitrogen of 510 and 170 wt.ppm measured in the charge increased up to 870 and 235 wt.ppm in the tilt-cast Nb, NbTa, and NbW alloys, respectively. The present results indicate that VIM in graphite crucibles followed by casting into graphite moulds can be applied for processing of TiAl-based alloys containing minor additions of strong carbide forming elements such as Nb, Ta, and W.

3.2. Microstructure and Phase Analysis

Figure 1 shows the typical microstructures of the tilt-cast Nb, NbTa, and NbW alloys. The microstructure of the alloys consists of the matrix composed of two single phases (1 and 5) and lamellar regions (2). The matrices of the alloys are reinforced with relatively uniformly distributed plate-like and irregularly shaped particles (3 and 4), as seen in Figure 1a,c and e. Figure 2 shows XRD diffraction patterns and summarises four identified coexisting phases in the studied alloys: α2, γ, β/B2, and Ti2AlC. Table 2 shows the average chemical composition with standard error of the measurement of the coexisting phases and microstructural regions measured by EDS. Based on the measured chemical composition and XRD analyses, the microstructure of the studied Nb, NbTa, and NbW alloys consist of equiaxed α2 + γ lamellar grains and grain boundaries containing a single γ phase, as seen in Figure 1b,d and f. In addition, the matrix of the NbW alloy contains an irregular network enriched in Nb and W that belongs to β/B2 phase (disordered/ordered Ti-based solid solution with a cubic crystal structure), as seen in Figure 1f and Table 2. Figure 3, Figure 4 and Figure 5 show the distribution of alloying elements in the reinforcing carbide particles and the surrounding matrix. In all studied alloys, cores of the coarse irregular shaped (Ti,Nb,X)2AlC particles, where X is Ta or Nb, are enriched in Ti but depleted in Al and Nb compared to those of the surrounding (Ti,Nb,X)2AlC phase or the matrix. The measured chemical compositions of these particles correspond to the (Ti,Nb,X)C phase but their amount in the studied Nb, NbTa, and NbW alloys are under the detectable limit of the XRD analysis. The EDS map and point analyses indicate that Ta as a strong carbide forming element segregates intensively to both carbide phases in the NbTa alloy, as seen in Figure 4 and Table 2. However, its content is slightly higher in (Ti,Nb,Ta)C than in (Ti,Nb,Ta)2AlC. Alloying with Ta reduces the dissolution of Nb in both (Ti,Nb,Ta)C and (Ti,Nb,Ta)2AlC phases compared to that in the Nb alloy. The EDS maps in Figure 5 and Table 2 indicate that W in the NbW alloy segregates predominantly to the β/B2 phase and its dissolution is reduced in (Ti,Nb,W)C and (Ti,Nb,W)2AlC phases. Alloying with W leads to a decrease in Nb content in both (Ti,Nb,W)C and (Ti,Nb,W)2AlC phases and increases its content in the lamellar α2 + γ regions compared to those in the Nb alloy, as shown in Table 2.
The results of the statistical evaluation of quantitative image analysis data of carbide particles, such as mean length of major axis Lmaj, mean length minor axis Lmin, volume fraction VC, and shape factor defined as SF = 4πA/P2, where A is the area and P is the perimeter of the carbide particle, are summarised in Table 3. It is clear from Table 3 that the alloying with Ta and W does not statistically affect the volume fraction of the primary carbide particles, which is measured to be about 7.4 vol.% in the studied alloys. The increase in the shape factor SF in the NbW alloy indicates an increased number of irregularly shaped particles at the expense of a decreasing number of plate-like ones compared to those in the Nb and NbTa alloys. The alloying with Ta and W decreases the size of carbide particles, indicated by a decrease in mean lengths of major and minor axes, as shown in Table 3.
The microstructure of the tilt-cast Nb, NbTa, and NbW alloys consists of equiaxed grains with nearly lamellar microstructure, as seen in Figure 1. Figure 6 shows the statistical evolution of the measured grain size (about 1000 measurements for each alloy) and interlamellar spacing (about 2000 measurements for each alloy). The correlation coefficients of the log-normal regression curves of grain size are better than r2 = 0.98, as seen in Figure 6a. The statistical mean values of the grain size d are summarised in Table 3. The measured mean grain size ranging from 33.6 to 42.5 μm indicates that the formation of about 7.4 vol.% of primary carbide particles very efficiently refines the grain structure of the studied alloys compared to that observed in numerous as-cast TiAl-based alloys without or low volume fraction primary carbide particles [46,57,58]. Although there are clear statistical differences between the mean grain size measured in the studied alloys, indicating the more efficient effect of the carbide particles on the nucleation rate and grain growth in the Nb alloy, this effect can be appropriately evaluated only after full heat treatments required for microstructure optimisation [41,44,59].
Figure 6b shows log-normal distribution curves of α2-α2 interlamellar spacing λ. It should be noted that the correlation coefficients of the log-normal regression curves are better than r2 = 0.97. It is clear from Table 3 that the alloying with W significantly refines the lamellar microstructure and leads to a significantly lower mean interlamellar spacing λ in the tilt-cast NbW alloy compared to that of the Nb and NbTa ones.

3.3. Nanohardness, Microhardness, and Hardness

Figure 7 shows the effect of Ta and W on Vickers hardness and indentation elastic modulus of Ti-45Al-5.6Nb-2C (at.%) alloy. The alloying with Ta and W has no significant effect on average Vickers hardness HV30 measured on the tilt-cast samples. All variations in the Vickers hardness values fall well into the experimental errors of the measurements, as seen in Figure 7a. Similarly, Ta or W additions have no significant effect on the elastic modulus of the studied alloys. All variations in the measured elastic modulus fall into the experimental error of the measurements, as seen in Figure 7b. It should be noted that the average values are calculated from 30 independent measurements carried out for each alloy. The applied load of 298 N was sufficiently high to deform a representative volume and create the indentations with dimensions covering nine or more equiaxed grains in each alloy. Figure 8 shows nanohardness and elastic modulus of the reinforcing (Ti,Nb,X)2AlC particles. Even though the segregation behaviour of Ta differs significantly from that of W in the studied alloys, the measured nanohardness values and indentation elastic modulus for the (Ti,Nb,X)2AlC phase are lower in the NbTa and NbW alloys than those in the Nb alloy, as seen in Figure 8a,b. Figure 9 shows Vickers microhardness of lamellar α2 + γ regions. The Vickers microhardness HV0.05 of the NbW alloy is higher than that of the Nb and NbTa alloys, as shown in Figure 9a. The Vickers microhardness increases with the decreasing of the measured mean interlamellar α2–α2 spacing, as shown in Figure 9b.

4. Discussion

4.1. VIM in Graphite Crucibles and Casting into Graphite Moulds

The VIM in graphite crucibles is connected with an increase in carbon content that depends on alloy chemistry, melt temperature, reaction time, quality of graphite, the volume of the melt, and contact area between the melt and graphite crucible [45,46,47,60]. Szkliniarz and Szkliniarz [43] have reported that the VIM of Ti-47Al-2W-0.5Si (at.%) alloy in graphite crucibles at 1650 °C leads to an increase in C content by 0.35 at.%. Lapin and Klimová [44] have reported an increase in C content by 0.9 at.% in Ti-xAl-8Nb-3.6C-0.8Mo (at.%) alloys, where x = 38, 42 and 45 at.%, during their VIM in graphite crucibles at 1700 °C. The increase in C content can be mainly related to the interactions between carbide forming elements in the melt and graphite walls. Assuming the chemical composition of the studied alloys and Gibbs free energy ΔG of carbide formation at 1650 °C [61], the following exothermic reactions can contribute to the increase in C content in VIM alloys: Ti + C = TiC (ΔG = −162 kJ/mol), Ta + C = TaC and (ΔG = −142 kJ/mol), Nb + C = NbC (ΔG = −127 kJ/mol), 2Nb + C = Nb2C (ΔG= −184 kJ/mol), and 2W + C = W2C (ΔG = −35 kJ/mol). The geometrical arrangement of the charge in the graphite crucible and electromagnetic steering of the melt can supply high convective forces for mechanical detachment of carbide particles from the growing reaction layer and their redistribution into the melt. The solid carbide particles can be either preserved, partially dissolved, or fully dissolved in the melt. It should be noted that the extent of the reaction between the graphite crucible and the melt strongly depends on the quality of the graphite crucible. In the case of the studied Nb, NbTa, and NbW alloys, the VIM in high-density isostatic graphite crucibles resulted in a reproducible increase in C content due to graphite–melt interactions by (0.42 ± 0.06) at.% at 1650 °C. In order to achieve the designed chemical composition of the alloys, the content of the alloying TiC particles in the charge was adjusted to this increase in C content resulting from the graphite crucible-melt exothermic reactions.
The studied alloys show a reproducible content of Al. As has been shown in recent works [44,46], the low vacuum pressure of 6.8 Pa (medium vacuum) leads to intensive evaporation of Al during VIM of TiAl-based alloys. Vacuum pressure ranging from 1 × 103 to 1 × 104 Pa (low vacuum) has been reported to be sufficiently high to prevent Al evaporation even in Nb-rich TiAl-based alloys [36,37,43,60].

4.2. Microstructure Formation

Numerous irregularly shaped (Ti,Nb,X)2AlC particles contain central regions enriched by C and lean in Al, as shown in Figure 1 and Table 2. However, their size, and especially their chemical composition, differ significantly from a size of 3 μm and chemistry of the pure TiC powder used for the alloying. Hence, during the melting of the charge in graphite crucibles, the TiC powder decomposed, and thermodynamically more stable and chemically more complex nonstoichiometric (Ti,Nb,X)C1-x particles were formed. In the ternary Ti-Al-C system, the single liquid (L) phase region can be achieved only at temperatures significantly higher than the melt temperature used in the present work. Witusiewicz et al. [62] have calculated that, below full melting of Ti-Al-C alloys, a two-phase region composed of the liquid and nonstoichiometric solid TiC1-x phase is thermodynamically stable. Taking into account the reported thermodynamic calculations and microstructural observations shown in Figure 1 and Figure 3, Figure 4 and Figure 5, the stabilisation temperature of the charge of 1650 °C corresponds to a phase equilibrium L + (Ti,Nb,X)C1−x in the studied alloys before casting. After the tilt casting into a cold graphite mould, the temperature of the melt decreases continuously and a new phase equilibrium is achieved according to transformation pathway L + (Ti,Nb,X)C1−x → L + (Ti,Nb,X)2AlC. This reaction/transformation starts by the growth of solid (Ti,Nb,X)2AlC layer at the L/(Ti,Nb,X)C1−x interface and leads to a partial or full transformation of (Ti,Nb,X)C1−x particles to the (Ti,Nb,X)2AlC phase. Since the remaining liquid is still supersaturated by carbon, the solidification continues by the formation of plate-like (Ti,Nb,X)2AlC particles that nucleate spontaneously from the melt, but their growth is restricted by other growing carbide particles. The carbide particles forming in the melt serve as nucleation sites for the formation of β phase, according to transformation pathway L + (Ti,Nb,X)2AlC → L + β + (Ti,Nb,X)2AlC → β + α + (Ti,Nb,X)2AlC. The microstructure formation is finalised by the solid-state phase transformations β + α + (Ti,Nb,X)2AlC → β + α + γ + (Ti,Nb,X)2AlC → α2 + γ + (Ti,Nb,X)2AlC. The presence of γ, β, and (Ti,Nb,X)C phase regions in the tilt-cast alloys result from Al segregation into the interdendritic region, the incomplete transformation of β to α phase, and insufficient time and temperature for the diffusion-controlled transformation of (Ti,Nb,X)C to (Ti,Nb,X)2AlC, respectively. As shown by Lapin and Klimová [44], the size of the primary carbide particles is strongly affected by the casting technology. Fast turbulent filling of the mould and fast cooling rates lead to the formation of finer primary carbide particles and smaller grain size during centrifugal casting compared to those formed during tilt casting.

4.3. Mechanical Properties

Although the Vickers hardness measurements do not show a clear effect of the alloying on room temperature mechanical properties, the nanohardness and microhardness values indicate some differences between the tilt-cast Nb, NbTa, and NbW alloys. Figure 8 shows the variations in nanohardness values of primary (Ti,Nb,X)2AlC particles, indicating softening and lower indentation elastic modulus of the NbTa alloy compared to those in the Nb and NbW alloys. However, this softening cannot be attributed only to the substitution of Ti with Nb and Ta in the (Ti,Nb,Ta)2AlC phase when compared with the chemistry of this phase in the Nb and NbW alloys, as seen in Table 2. Meng et al. [63] have shown that a significantly higher concentration of V (atomic radius of V corresponds to that of W) substituting for Ti than the measured variations in Nb, Ta, and W contents in Ti2AlC phase is required to quantify its solid solution strengthening effect using Vickers hardness measurements. The observed variations in the nanohardness and elastic modulus in the studied alloys can also be significantly affected by the crystallographic orientations of the nanoindentation planes of the (Ti,Nb,X)2AlC phase. Liu et al. [64] have calculated that the nanoindentation hardness depends on the nanoindentation crystallographic plane and its increase is also connected with an increase in the elastic modulus. Figure 8 qualitatively shows a very similar evolution of nanohardness with indentation elastic modulus of the (Ti,Nb,X)2AlC phase in the Nb, NbTa, and NbW alloys to that calculated by Liu et al. [64] for Ti2AlC crystals. Despite the statistical number of measurements performed on the carbide particles in each studied alloy (up to 70 measurements), the effect of nanoindentation planes on the variations of the measured nanohardness and indentation elastic modulus values cannot be fully excluded.
The aligned α2 lamellae are strong barriers to dislocation motion in the lamellar α2 + γ regions. Dislocations generated in the softer γ phase pile-up at α2/γ interfaces during deformation at room and high temperatures [65,66]. The decrease in the α2-α2 interlamellar spacing in the NbW compared to those of the Nb and NbTa alloys increases the number of strong α2/γ interfaces that hinder dislocation motion during plastic deformation below the indenter. The increased strain hardening of the lamellar α2 + γ matrix leads to an increase in Vickers microhardness values. As shown by several authors [66,67], the Vickers microhardness can be related to the α2-α2 interlamellar spacing according to a Hall–Petsch type relationship in a form HV ∝ λ−0.5 and also linearly to offset compression yield strength of TiAl-based alloys with lamellar microstructure [68].

5. Conclusions

The effect of Ta and W additions on microstructure and mechanical properties of tilt-cast Ti-45Al-5Nb-2C (at.%) alloy was studied. The following conclusions are reached:
  • The VIM in graphite crucibles followed by tilt casting into graphite moulds leads to the reproducible chemical compositions of the tilt-cast alloys designated as Nb, NbTa, and NbW that correspond well to the nominal ones.
  • The microstructures of the tilt-cast alloys consist of primary plate-like and irregularly shaped (Ti,Nb,X)2AlC particles, where X is Ta or W, that are relatively uniformly distributed in the intermetallic matrices composed of α2, γ, and β/B2 phases. A small amount of (Ti,Nb,X)C phase is identified in the cores of some coarse irregularly shaped particles.
  • The carbide particles serve as the nucleation sites for β dendrites during solidification and contribute effectively to the grain refinement of the studied tilt-cast Nb, NbTa, and NbW alloys. The alloying with W refines the lamellar α2 + γ microstructure and leads to the preservation of an untransformed irregularly shaped network of the β/B2 phase during solid-state phase transformations.
  • The alloying with Ta and W has no significant effect on Vickers hardness and indentation elastic modulus of the studied Nb, NbTa, and NbW alloys. The measured nanohardness and indentation modulus of the reinforcing (Ti,Nb,Ta)2AlC particles in the Ta containing NbTa alloy are significantly lower compared to those of (Ti,Nb)2AlC or (Ti,Nb,W)2AlC phases of the Nb and NbW alloys, respectively. The Vickers microhardness of lamellar α2 + γ regions measured in the NbW alloy is significantly higher than that of the Nb and NbTa alloys.

Author Contributions

Conceptualisation, J.L.; methodology, J.L. and K.K.; validation, J.L.; formal analysis, J.L. and K.K.; investigation, K.K. and J.L.; writing—original draft preparation, J.L.; writing—review and editing, J.L. and K.K.; visualisation, K.K. and J.L.; supervision, J.L. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Slovak Grant Agency for Science under the contract VEGA 2/0074/19 and the Slovak Research and Development Agency under the contracts APVV-20-0505 and APVV-15-0660.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

All the data generated during this study are included in this article.

Acknowledgments

The authors would like to thank Soňa Kružlíková for technical assistance and metallographic preparation of samples.

Conflicts of Interest

The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript; or in the decision to publish the results.

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Figure 1. BSE micrographs showing the microstructure of tilt-cast alloys: (a,b) Nb, (c,d) NbTa, and (e,f) NbW.
Figure 1. BSE micrographs showing the microstructure of tilt-cast alloys: (a,b) Nb, (c,d) NbTa, and (e,f) NbW.
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Figure 2. The typical XRD patterns of tilt-cast Nb, NbTa, and NbW alloys.
Figure 2. The typical XRD patterns of tilt-cast Nb, NbTa, and NbW alloys.
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Figure 3. SEM micrograph and corresponding EDS map analysis of the tilt-cast Nb alloy.
Figure 3. SEM micrograph and corresponding EDS map analysis of the tilt-cast Nb alloy.
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Figure 4. SEM micrograph and corresponding EDS map analysis of the tilt-cast NbTa alloy.
Figure 4. SEM micrograph and corresponding EDS map analysis of the tilt-cast NbTa alloy.
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Figure 5. SEM micrograph and corresponding EDS map analysis of the tilt-cast NbW alloy.
Figure 5. SEM micrograph and corresponding EDS map analysis of the tilt-cast NbW alloy.
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Figure 6. Measured grain size and interlamellar spacing in Nb, NbTa, and NbW alloys after tilt casting: (a) log-normal distribution curves of grain size, (b) log-normal distribution curves of α2-α2 interlamellar spacing. The types of the distribution curves are marked in the figure.
Figure 6. Measured grain size and interlamellar spacing in Nb, NbTa, and NbW alloys after tilt casting: (a) log-normal distribution curves of grain size, (b) log-normal distribution curves of α2-α2 interlamellar spacing. The types of the distribution curves are marked in the figure.
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Figure 7. Vickers hardness and elastic modulus of tilt-cast Nb, NbTa, and NbW alloys: (a) Vicker hardness HV30, (b) indentation elastic modulus. The studied alloys are marked in the figures.
Figure 7. Vickers hardness and elastic modulus of tilt-cast Nb, NbTa, and NbW alloys: (a) Vicker hardness HV30, (b) indentation elastic modulus. The studied alloys are marked in the figures.
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Figure 8. Nanohardness and elastic modulus of Ti2AlC particles in the tilt-cast Nb, NbTa, and NbW alloys: (a) nanohardness, (b) indentation elastic modulus. The studied alloys are marked in the figures.
Figure 8. Nanohardness and elastic modulus of Ti2AlC particles in the tilt-cast Nb, NbTa, and NbW alloys: (a) nanohardness, (b) indentation elastic modulus. The studied alloys are marked in the figures.
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Figure 9. Vickers microhardness of lamellar α2 + γ regions in the tilt-cast Nb, NbTa, and NbW alloys: (a) Vickers microhardness HV0.05, (b) dependence of Vickers microhardness HV0.05 on α2-α2 interlamellar spacing. The studied alloys are marked in the figures.
Figure 9. Vickers microhardness of lamellar α2 + γ regions in the tilt-cast Nb, NbTa, and NbW alloys: (a) Vickers microhardness HV0.05, (b) dependence of Vickers microhardness HV0.05 on α2-α2 interlamellar spacing. The studied alloys are marked in the figures.
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Table
Table 1. Nominal and measured chemical composition of studied tilt-cast alloys (at.%).
Table 1. Nominal and measured chemical composition of studied tilt-cast alloys (at.%).
Alloy TiAlNbTaWC
Nbnominal48455--2
measured48.1 ± 0.344.9 ± 0.35.0 ± 0.1--2.0 ± 0.1
NbTanominal464552-2
measured46.2 ± 0.244.9 ± 0.25.0 ± 0.12.0 ± 0.1-1.9 ± 0.1
NbWnominal46455-22
measured46.0 ± 0.345.0 ± 0.35.1 ± 0.1-2.0 ± 0.12.1 ± 0.1
Table
Table 2. Chemical composition of coexisting phases of the studied alloys (at.%).
Table 2. Chemical composition of coexisting phases of the studied alloys (at.%).
AlloyRegionPhaseTiAlNbTaWC
Nb1γ46.8 ± 0.248.1 ± 0.25.1 ± 0.1---
2α2 + γ48.5 ± 0.145.8 ± 0.25.7 ± 0.1---
3TiC55.6 ± 0.81.5 ± 0.42.8 ± 0.2--40.1 ± 0.7
4Ti2AlC44.2 ± 0.224.6 ± 0.14.3 ± 0.1--26.9 ± 0.4
NbTa1γ45.9 ± 0.247.6 ± 0.35.4 ± 0.11.1 ± 0.1--
2α2 + γ46.7 ± 0.146.3 ± 0.25.8 ± 0.11.2 ± 0.1--
3TiC44.4 ± 0.10.4 ± 0.23.1 ± 0.15.8 ± 0.1-46.3 ± 0.5
4Ti2AlC42.0 ± 0.323.6 ± 0.43.8 ± 0.13.9 ± 0.1 26.7 ± 0.5
NbW1γ43.6 ± 0.250.5 ± 0.35.0 ± 0.2-0.9 ± 0.1-
2α2 + γ45.3 ± 0.147.1 ± 0.26.2 ± 0.1-1.4 ± 0.1-
3TiC53.9 ± 0.92.1 ± 0.92.1 ± 0.4-0.2 ± 0.141.7 ± 0.4
4Ti2AlC45.2 ± 0.923.7 ± 1.23.6 ± 0.2-0.4 ± 0.127.1 ± 0.6
5β46.8 ± 0.138.6 ± 0.28.2 ± 0.1-6.4 ± 0.1-
Table
Table 3. Measured microstructural parameters of the tilt-cast Nb, NbTa, and NbW alloys.
Table 3. Measured microstructural parameters of the tilt-cast Nb, NbTa, and NbW alloys.
AlloyLmaj
(μm)		Lmin
(μm)		SFVC (vol.%)d
(μm)		λ
(μm)
Nb13.0 ± 0.43.8 ± 0.10.36 ± 0.057.5 ± 0.433.6 ± 0.60.638 ± 0.019
NbTa10.5 ± 0.42.3 ± 0.10.37 ± 0.057.1 ± 0.636.6 ± 0.70.678 ± 0.014
NbW11.0 ± 0.33.6 ± 0.10.43 ± 0.017.7± 0.542.5 ± 0.80.531 ± 0.009
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Lapin, J.; Kamyshnykova, K. Effect of Ta and W Additions on Microstructure and Mechanical Properties of Tilt-Cast Ti-45Al-5Nb-2C Alloy. Metals 2021, 11, 2052. https://doi.org/10.3390/met11122052

AMA Style

Lapin J, Kamyshnykova K. Effect of Ta and W Additions on Microstructure and Mechanical Properties of Tilt-Cast Ti-45Al-5Nb-2C Alloy. Metals. 2021; 11(12):2052. https://doi.org/10.3390/met11122052

Chicago/Turabian Style

Lapin, Juraj, and Kateryna Kamyshnykova. 2021. "Effect of Ta and W Additions on Microstructure and Mechanical Properties of Tilt-Cast Ti-45Al-5Nb-2C Alloy" Metals 11, no. 12: 2052. https://doi.org/10.3390/met11122052

APA Style

Lapin, J., & Kamyshnykova, K. (2021). Effect of Ta and W Additions on Microstructure and Mechanical Properties of Tilt-Cast Ti-45Al-5Nb-2C Alloy. Metals, 11(12), 2052. https://doi.org/10.3390/met11122052

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