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Article

Synthesis of a Vanadium-Substituted Fe–Ti-Based Ternary Alloy via Mechanical Alloying, Compacting, and Post-Annealing

by
Abhishek Kumar Patel
1,
Davide Violi
2,
Ivan Lorenzon
3,
Carlo Luetto
2,
Paola Rizzi
1 and
Marcello Baricco
1,*
1
Department of Chemistry, NIS and INSTM, University of Turin, Via Pietro Giuria 7, 10125 Turin, Italy
2
Methydor S.r.l., Via Antonio Cecchi 4/4, 16129 Genova, Italy
3
Pometon S.p.A., Via Circonvallazione 62, 30030 Maerne (VE), Italy
*
Author to whom correspondence should be addressed.
Metals 2025, 15(7), 723; https://doi.org/10.3390/met15070723
Submission received: 18 May 2025 / Revised: 18 June 2025 / Accepted: 26 June 2025 / Published: 28 June 2025
(This article belongs to the Special Issue Synthesis, Processing and Applications of New Forms of Metals)
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Abstract

In this study, we address the need for sustainable and scalable synthesis routes for hydrogen storage materials by developing a FeTi alloy in which vanadium (V) partially substitutes for titanium (Ti). The alloy was synthesized using mechanical alloying, compaction, and post-annealing, employing industrial-grade Fe and Ti powders and an alternative to pure vanadium, i.e., ferrovanadium (Fe–V). X-ray diffraction (XRD) analysis of the mechanically alloyed mixture revealed the partial formation of a Fe(V) solid solution, along with residual Ti. Subsequent compaction and annealing at 1000 °C led to the formation of the FeTi(V) phase, accompanied by two minor secondary phases, Fe2Ti and Fe2Ti4O. A maximum phase yield of 90% for FeTi was achieved after 48 h of annealing. The novelty of this work lies in the demonstration of a sustainable and economical synthesis approach for V-substituted FeTi alloys using industrial-grade raw materials, offering a potential reduction in the carbon footprint compared with conventional melting techniques.

1. Introduction

Metal hydrides are considered excellent candidates for solid-state hydrogen storage because of their high volumetric capacity and reasonable operating pressure and temperature conditions [1,2,3,4]. Among the metal hydrides, FeTi alloys show great potential for hydrogen storage in stationary applications, such as forklifts [5,6] and fuel-cell-powered submarines [7], because of their good volumetric capacity and low cost. However, FeTi alloys need to be exposed to high temperature and hydrogen pressure, a process usually called activation, in order to break the surface oxide layer (typically TiO2) and to enable hydrogen absorption [8,9,10]. Several studies have demonstrated that the addition or substitution of transition metals such as Zr, V, and Mn can significantly reduce the activation requirements of FeTi alloys [11,12,13,14,15]. These modifications may promote the formation of secondary phases, which generate grain boundaries, thereby facilitating enhanced hydrogen diffusion pathways. Also, researchers reported that severe plastic deformation, such as high-pressure torsion [16] and ball milling [17,18,19] can promote easier activation of FeTi alloys.
The most widely used synthesis processes for FeTi alloys are melting techniques, such as arc melting and induction melting [20,21]. While effective, these methods are energy-intensive and contribute significantly to carbon emissions due to their high-temperature operations and to the production of the alloy in pellet form, which further requires pulverization. In contrast, mechanical alloying (MA) of the starting materials is another method for their preparation, offering a more energy-efficient process and the advantage of directly producing a powder with a nanocrystalline microstructure [22,23,24,25].
Chu et al. synthesized amorphous FeTi through mechanical alloying for 25 h, observing a significantly lower hydrogen capacity after activation compared with crystalline FeTi [26]. Similarly, Eckert et al. reported amorphization during prolonged ball milling (60 h) [23]. Jung et al. compared melting techniques with mechanical alloying, finding that MA reduced the absorption plateau pressure for TiFe0.4Ni0.6 alloy, facilitating hydrogenation more effectively than arc melting [27]. Hotta et al. noted that FeTi crystallizes into a CsCl-type structure after 20 h of MA using Ti and Fe powders as starting materials [22]. Abe et al. explored synthesizing FeTi alloy using MA for shorter durations, followed by post-annealing [28]. Zadorozhnyy et al. synthesized nanocrystalline FeTi powder through MA, followed by post-compacting and annealing, yielding a bulk sample with high durability over 20 absorption–desorption cycles and an approximately 1.4 wt% hydrogen absorption capacity [29]. MA, followed by powder compaction and annealing, increases the pore density in the material compared with casting while providing better uniformity in the chemical composition and improved solid solubility. This method is more advantageous than melting techniques, which typically cause elemental segregation. The MA method reduced mechanical energy consumption, minimized contamination, and lessened powder adhesion, which are crucial for the successful synthesis of nanostructured FeTi alloys. However, the majority of these studies relied on high-purity raw materials and did not address the overall process yield or carbon footprint.
In response to growing environmental concerns, recent research studies have involved using industrial-grade starting materials for large-scale FeTi alloy synthesis [15,20,21]. Shang et al. achieved FeTi alloy synthesis via arc melting with industrial iron and titanium scraps, reducing the carbon footprint [30]. GKN Powder Metallurgy in Germany synthesized FeTi alloy on a large scale with Mn addition to facilitate hydrogenation [20]. Rasheed et al. reported direct FeTi alloy synthesis via metallothermic reduction of ilmenite sand (FeTiO3), eliminating the separate extraction and melting of starting materials [31].
Motivated by the need to develop environmentally sustainable hydrogen storage materials, this study investigates the synthesis of FeTi alloy using industrial-grade Fe and Ti powders through a short-duration mechanical alloying process, followed by compaction and post-annealing. In addition, vanadium was introduced via ferrovanadium (Fe–V) as a substitution for pure Ti to improve the hydrogen absorption characteristics. This novel approach not only reduces the dependence on high-purity precursors and energy-intensive processes but also supports carbon emission-reduction goals through a more sustainable synthesis route.

2. Materials and Methods

The alloy was synthesized by using mechanical alloying, post-compaction, and annealing processes. The starting materials were industrial-grade titanium and iron. Ferrovanadium (Fe–V) was used as a source of vanadium, with the aim of substituting the titanium with vanadium. The starting materials were in powder form and were provided by Pometon, Maerne (VE), Italy. The purity and particle size distribution of the starting materials were determined through structural, microstructural, and chemical composition analysis of the powders.
A mixture of Ti, Fe, and Fe–V powders with a stoichiometric composition of Ti0.9FeV0.06 were mechanically alloyed using a planetary ball mill (Fritsch Pulverisette 6, Fritsch GmbH, Idar-Oberstein, Germany). The mole fraction of iron was adjusted according to the determined fractions of iron and vanadium in the ferrovanadium. MA was carried out in an 80 mL stainless steel jar using 10 mm diameter stainless steel balls, with a ball-to-powder ratio of 10:1, at a rotation speed of 250 rpm under an argon (Ar) atmosphere. The milling was performed for short durations of 1 h, 3 h, and 5 h. In the case of 3 h and 5 h of milling, a pause time of 1 h and a reverse mode after each hour of milling were applied to cool the jar and to promote thorough mixing.
Subsequently, a 3-ton manual hydraulic press was used to compact the mechanically milled powders in a 13 mm diameter die for consolidation. The consolidated powders, wrapped in tantalum foil, were sealed in a quartz ampoule under a mild argon pressure (~mbar). The compacted powders were wrapped in a tantalum foil to capture any residual oxygen molecules present in the annealing atmosphere. Annealing was performed in a muffle furnace at 1000 °C for 48 h. The annealing temperature was selected based on the highest stable temperature of the FeTi phase, as indicated by the Fe–Ti binary phase diagram [32,33], while the annealing duration was chosen in accordance with conditions reported in previous studies [34]. X-ray diffraction (XRD) patterns of the mechanically milled and annealed powders were obtained using a Panalytical X’Pert X-ray Diffractometer (Malvern Panalytical, Malvern, UK) in Bragg–Brentano geometry with Cu Kα radiation. The diffraction patterns were qualitatively analyzed using X’Pert HighScore software v5.3a. Furthermore, Rietveld refinement of the XRD data was carried out with MAUD software v. 2.9997 [35] to determine the volume fractions and crystallographic parameters of the various phases. The microstructure and chemical composition of the starting materials and annealed powders were examined using a TESCAN VEGA scanning electron microscope (SEM) (Brno, Czech Republic) coupled with an Oxford Ultim-Max 100 energy-dispersive X-ray spectroscopy (EDX) system (Oxford Instruments, Abingdon, UK).

3. Results and Discussion

3.1. Starting Materials

Scanning electron microscope images of the starting materials Ti, Fe, and Fe–V are shown in Figure 1 at different magnifications. Particles with irregular shapes and agglomerated forms were observed for Ti and Fe–V, whereas in the case of Fe, a few globular-shaped particles were detected. The particle size distribution of the starting materials was estimated as follows: Ti particles were less than 50 µm, Fe particles ranged from 5 to 200 µm, and Fe–V particles ranged from 30 to 90 µm.
An elemental analysis of the starting materials was obtained through EDX using the point and square analysis method, as shown in Figure 2. The EDX analysis of the starting materials verified the presence of 99% pure Ti and of Fe, with minor oxide phases confined to the surface, as presented in Table 1. In the case of Fe–V, the average V content was found to be 84 wt%, with the presence of particles of aluminium oxide, which likely originated from the alumina crucible used during the synthesis of ferrovanadium.
The starting material powders were examined using powder XRD to determine the crystal structure of Ti, Fe, and Fe–V. Figure 3 presents the XRD patterns of Ti, Fe, and Fe–V. Ti powders show an HCP structure (a = b = 2.9504 Å, c = 4.6833 Å; space group P63/mmc). Additionally, a few minor oxide-related peaks were detected in the case of Ti. Fe exhibited a body-centered cubic (BCC) structure with a lattice parameter of a = 2.8684 Å and space group Im-3m. Fe–V powders show a BCC structure with space group Im-3m.
The exact concentration of iron and vanadium in Fe–V was determined through Rietveld refinement of the XRD pattern. The fitted pattern, along with the goodness-of-fit and residual value, is presented in Figure 4. The experimental data (black crosses) closely match the calculated profile (red line), indicating a high-quality fit. The difference plot (black line at the bottom) shows minimal deviation, confirming the reliability of the refinement. The refined lattice parameter of the BCC phase was evaluated as 2.9932 Å. Caraveo et al. investigated the variation in the BCC lattice parameter as a function of the vanadium content in Fe1−xVx solid solutions [36]. Based on the refined lattice parameter, the vanadium concentration was estimated to be 84 at%, which is in good agreement with the EDX analysis of the Fe–V powder.

3.2. Mechanical Alloying

The XRD patterns of mechanically alloyed Ti0.9FeV0.06 powders at different milling times, along with those of hand-mixed powders, are shown in Figure 5a. Distinct peaks corresponding to Fe (BCC phase) and Ti (hexagonal close packed (HCP) phase) are visible. The Fe–V XRD peaks were not clearly distinguishable due to the extremely low mole fraction of V in the stoichiometry of the alloy and the nearly identical lattice parameter of the BCC structure to that of iron. The presence of sharp peaks suggests a coarse grain microstructure, with no mixing, in the hand-mixed powders. In the early stage of milling (1 h and 3 h), peak broadening begins, indicating a reduction in the crystallite size and an increase in the lattice strain. Additionally, some peak intensities diminish, suggesting partial dissolution of certain phases, as shown in the enlargement reported in Figure 5b. High-energy milling induces plastic deformation and generates defects, such as dislocation and vacancies, which are attributed to the peaks broadening and the increased lattice strain [37,38]. Extended milling times (e.g., 5 h) result in further broadening of the XRD peaks. As the mechanical alloying proceeds, there is an indication of the formation of a solid solution, as suggested by the disappearance of some initial peaks and the appearance of new ones, implying alloying and interdiffusion between the Fe, Ti, and V elements.
The refined XRD patterns for the hand-mixed and mechanically alloyed powders are shown in Figure 6, Figure 7, Figure 8 and Figure 9. Table 2 provides a detailed summary of the refined crystallographic parameters, phase fractions, crystallite sizes, and microstrain values for the hand-mixed and mechanically alloyed powders at various milling times. The hand-mixed powders primarily contain Ti (74 wt%) and Fe (24 wt%), with a minor Fe–V phase (2 wt%). After 1 h of milling, the phase fractions remain largely unchanged, with Ti still dominant (76 wt%). However, slight variations in the lattice parameters of Fe and Fe–V indicate an initial alloying and increased microstrain. After the 3 h of milling, significant changes are observed, with Fe increasing to 67 wt% and Ti decreasing to 30 wt%, suggesting enhanced interdiffusion and the potential formation of a solid solution. After 5 h of milling, Fe remains the predominant phase (60 wt%), and Ti decreases further to 38 wt%, indicating the ongoing alloying and partial dissolution of Ti into Fe. The Fe–V phase remains relatively constant at approximately 2 wt%.
During the early stage of MA (1 h), the increase in Ti lattice parameters reflects a strain buildup, while the subsequent decrease indicates atomic rearrangement and interdiffusion, similar to observations in other Ti-based systems [39]. The observed decrease in Fe and Ti lattice parameters with increased milling time suggests progressive alloying, which is consistent with previous studies on mechanical alloying and phase evolution in the same system [40,41]. Conversely, the Fe–V phase remains largely unchanged, suggesting minimal interaction with the Fe–Ti system during mechanical milling.
The crystallite size of Fe decreases from 110 nm in the hand-mixed state to 56 nm after 5 h of mechanical milling, confirming grain refinement due to high-energy MA, which is consistent with findings by Benjamin et al. [42] and El-Eskandarany et al. [43]. Similarly, the crystallite size of Ti significantly decreases from 263 nm to 86 nm, indicating plastic deformation and grain fragmentation. Additionally, the increasing microstrain values with milling time suggest significant lattice distortion and internal strain buildup.
These findings indicate that partial alloying between Fe and Ti occurs after 3 h of mechanical alloying, which is evidenced by the phase shift and by the reduction in crystallite size. Furthermore, the microstrain increases with milling time, suggesting a higher concentration of defects and lattice distortions. So, this MA treatment was selected for further steps in the synthesis process.

3.3. Post-Compacting and Annealing Treatment

The powders that had been milled for 3 h were compacted and annealed at 1000 °C in order to promote the formation of the FeTi phase. Figure 10 shows the XRD pattern of the annealed sample after compaction, which was analyzed by Rietveld refinement. The diffraction pattern revealed peaks corresponding to the FeTi phase, along with minor contributions from the Fe2Ti and Fe2Ti4O phases. Similar phases were also evidenced by Barale et al. [20] and Jung et al. [11] in Fe–Ti–Mn and Fe–Ti–V ternary alloys, respectively.
After annealing of mechanically alloyed and compacted FeTi–V alloy, around 90 wt% of the FeTi phase was achieved, with minor phases having a fraction of 4 wt% Fe2Ti and 6 wt% Fe2Ti4O, as shown in Table 3. The lattice parameter of the FeTi phase showed a slight decrease from that of the binary intermetallic compound, which has a B2 crystal structure with a lattice parameter of a = 2.9760 Å [44]. Since vanadium has a smaller atomic radius than titanium (Ti: 1.47 Å, V: 1.34 Å), its substitution in the FeTi phase leads to a decrease in the lattice parameter, consistent with observations reported in previous studies [11,44]. The FeTi phase exhibited a crystallite size of 118 nm, indicating its coarse microstructure, as expected after a high-temperature annealing. The microstrain remained nearly constant across all the phases. The Fe2Ti4O at the surface can play a crucial role during activation by, for example, enhancing the catalytic effect of Fe, which promotes H2 dissociation. This, in turn, facilitates hydrogenation of the matrix due to the interaction at the oxide–matrix interface [20].
The EDX analysis of the annealed sample, using the square method, is shown in Figure 11, and the composition of different phases is presented in Table 4. Elemental analysis revealed that the FeTi phase (spectrum 6) consists of approximately 50.7 at% Ti and 49.3 at% Fe. The Fe2Ti phase (spectrum 5) contains 35.3 at% Ti and 64.7 at% Fe. Meanwhile, the Fe2Ti4O phase (spectrum 4) shows a composition of 57.3 at% Ti, 32.3 at% Fe, 1.2 at% V, and 9.2 at% O, indicating the incorporation of vanadium in the oxide phase. However, in this study, the EDS results were used primarily for qualitative and semi-quantitative analyses to indicate the presence of oxygen and to support the phase identification discussed in the XRD results. A discrepancy was observed when comparing the present results with those reported by Jung et al. [11], who synthesized FeTi-based ternary alloys partially substituted with a vanadium element of similar composition via arc melting using high-purity starting materials. In the present study, vanadium was found to substitute only within the Ti-rich oxide phase, with no detectable incorporation in the primary FeTi phase. In contrast, Jung et al. [11] reported the presence of approximately 2 at% vanadium within the FeTi phase. However, they also noted that the measured composition may not accurately reflect the overall FeTi stoichiometry due to local variations in the vanadium distribution.

4. Conclusions

In this work, the synthesis of vanadium-substituted FeTi-based ternary alloys was successfully achieved through mechanical alloying followed by compaction and annealing. Rietveld refinement of the XRD pattern for annealed samples—prepared from compacted powders milled for 3 h—indicated an estimated FeTi phase yield of approximately 90%. This work was focused on the critical role of vanadium substitution in Fe–Ti alloys, aimed at tailoring the equilibrium hydrogen desorption pressure, particularly above 5 bar, which is relevant for proton-exchange membrane (PEM) fuel-cell applications [11,20]. To ensure feasibility for large-scale production, industrial-grade powders were employed as starting materials. Although the synthesis method focused on practical applicability, the main goal of this study was to achieve around 90% yield of the FeTi phase, with small amounts of secondary phases like Fe2Ti and Fe2Ti4O. The Fe2Ti4O phase plays a critical role in enhancing the hydrogen storage properties of Fe–Ti alloys [44,45]. This oxide phase typically forms during mechanical milling, heat treatments, or arc melting and significantly promotes the activation behavior of the FeTi intermetallic compound [46,47]. Firstly, the presence of Fe2Ti4O contributes to surface modification by introducing defects and grain boundaries, which serve as additional diffusion pathways for hydrogen atoms, thereby facilitating improved hydrogen penetration. Secondly, this oxide phase may exhibit a catalytic effect by lowering the activation energy required for hydrogen dissociation and recombination, thus enhancing the kinetics of hydrogen absorption and desorption. Furthermore, while pure FeTi is prone to forming a dense and passive TiO2 layer that hinders hydrogen uptake [8,9], the formation of Fe2Ti4O disrupts this layer, promoting more efficient hydrogen activation and storage performance. The investigated synthesis route demonstrates the feasibility of producing FeTi alloys on a large scale using industrial-grade starting materials, making it a promising approach for hydrogen storage applications. The current study was primarily focused on the synthesis of FeTi alloy using industrial-grade raw materials. Future research will focus on investigating the hydrogen storage properties of this alloy.

Author Contributions

Conceptualization, M.B., P.R. and A.K.P.; methodology, A.K.P., D.V. and I.L.; software, A.K.P.; validation, M.B., C.L. and P.R.; formal analysis, A.K.P. and D.V.; investigation, A.K.P. and D.V.; resources, M.B., P.R. and I.L.; data curation, A.K.P. and D.V.; writing—original draft preparation, A.K.P.; writing—review and editing, C.L., P.R. and M.B.; visualization, M.B., P.R. and A.K.P.; supervision, M.B., C.L. and P.R.; project administration, M.B.; funding acquisition, M.B. and P.R. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by a REMEDHYS project that received funding from the European Union’s Horizon Europe research and innovation program and that was supported by the Clean Hydrogen Partnership and its members under Grant Agreement No. 10119203. The authors also acknowledge financial support under the National Recovery and Resilience Plan (NRRP), Mission 4, Component 2, Investment 1.1, Call for tender No. 1409, published on 14 September 2022 by the Italian Ministry of University and Research (MUR), funded by the European Union—NextGenerationEU—Project NEREHYDES—CUP B53D23006590006—Grant Assignment Decree No. 961 adopted on 30 June 2023 by the Italian Ministry of Ministry of University and Research (MUR). The research was supported by Project CH4.0 under the MUR program “Dipartimenti di Eccellenza 2023–2027” (CUP: D13C22003520001) and a part of the project NODES, which received funding from the MUR and M4C2 1.5 of PNRR under Grant Agreement No. ECS00000036.

Data Availability Statement

The raw data supporting the conclusions of this article will be made available by the authors on request.

Acknowledgments

Muhammad Owais is kindly acknowledged for his technical support.

Conflicts of Interest

Authors Davide Violi and Carlo Luetto were employed by the company Methydor S.r.l. Author Ivan Lorenzon was employed by the company Pometon S.p.A. The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

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Metals 15 00723 g001
Figure 1. SEM images of the starting materials: (a) Ti, (b) Fe, and (c) Fe–V.
Figure 1. SEM images of the starting materials: (a) Ti, (b) Fe, and (c) Fe–V.
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Figure 2. EDX analysis of the starting materials: (a) Ti, (b) Fe, and (c) Fe–V.
Figure 2. EDX analysis of the starting materials: (a) Ti, (b) Fe, and (c) Fe–V.
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Figure 3. X-ray diffraction patterns of the starting materials Ti (black), Fe (red), and Fe–V (blue).
Figure 3. X-ray diffraction patterns of the starting materials Ti (black), Fe (red), and Fe–V (blue).
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Figure 4. Rietveld refinements for XRD patterns of the Fe–V powders.
Figure 4. Rietveld refinements for XRD patterns of the Fe–V powders.
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Figure 5. X-ray diffraction patterns of MA Ti0.9Fe0.06 (a) that have been hand-mixed and mechanically alloyed for 1, 3, and 5 h and (b) enlargement of the low-angle XRD peaks.
Figure 5. X-ray diffraction patterns of MA Ti0.9Fe0.06 (a) that have been hand-mixed and mechanically alloyed for 1, 3, and 5 h and (b) enlargement of the low-angle XRD peaks.
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Figure 6. Rietveld refinements for XRD patterns of the hand-mixed Ti0.9FeV0.06 powders (GOF = 1.13, Rwp = 6.76%).
Figure 6. Rietveld refinements for XRD patterns of the hand-mixed Ti0.9FeV0.06 powders (GOF = 1.13, Rwp = 6.76%).
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Figure 7. Rietveld refinements for XRD patterns of the 1 h MA Ti0.9FeV0.06 powders (GOF = 1.31, Rwp = 6.39%).
Figure 7. Rietveld refinements for XRD patterns of the 1 h MA Ti0.9FeV0.06 powders (GOF = 1.31, Rwp = 6.39%).
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Figure 8. Rietveld refinements for XRD patterns of the 3 h MA Ti0.9FeV0.06 powders (GOF = 1.19, Rwp = 6.73%).
Figure 8. Rietveld refinements for XRD patterns of the 3 h MA Ti0.9FeV0.06 powders (GOF = 1.19, Rwp = 6.73%).
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Figure 9. Rietveld refinements for XRD patterns of the 5 h MA Ti0.9FeV0.06 powders (GOF = 1.52, Rwp = 6.02%).
Figure 9. Rietveld refinements for XRD patterns of the 5 h MA Ti0.9FeV0.06 powders (GOF = 1.52, Rwp = 6.02%).
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Figure 10. Rietveld refinement of the XRD pattern of compacted and annealed Ti0.9FeV0.06 powders (GOF = 1.13, Rwp = 6.76%).
Figure 10. Rietveld refinement of the XRD pattern of compacted and annealed Ti0.9FeV0.06 powders (GOF = 1.13, Rwp = 6.76%).
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Figure 11. EDX analysis using the square method for annealed consolidated Ti0.9FeV0.06 powders.
Figure 11. EDX analysis using the square method for annealed consolidated Ti0.9FeV0.06 powders.
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Table 1. Elemental analysis of the starting materials Ti, Fe, and Fe–V (wt%). The data refer to different spectra, as detailed in Figure 2.
Table 1. Elemental analysis of the starting materials Ti, Fe, and Fe–V (wt%). The data refer to different spectra, as detailed in Figure 2.
Fe12345678910
O1.8----2.92.11.11.11.3
Fe98.210010010010097.197.998.998.998.7
Ti11121314151617181920
O 14.9----12.2--12.9
Ti10085.110010010010087.810010087.1
Fe–V111213141516171819-
O----43.23-----
Al----56.77-----
Fe9.817.816.715.7-17.216.313.923.5-
V90.282.283.384.3-82.883.786.176.5-
Table 2. Refined crystallographic and microstructural parameters, including the phase fraction of each phase present, in hand-mixed and mechanically alloyed Ti0.9FeV0.06 powders after 1 h, 3 h, and 5 h.
Table 2. Refined crystallographic and microstructural parameters, including the phase fraction of each phase present, in hand-mixed and mechanically alloyed Ti0.9FeV0.06 powders after 1 h, 3 h, and 5 h.
Hand-Mixed (No Milling)
PhasePhase Fractions (wt%)Lattice Parameter (Å)Crystallite Size
(nm)		Microstrain × 10−3 (%)
Fe24(1)a = 2.8672(4)110(1)0.2(1)
Ti74(1)a = 2.9517(2)
c = 4.6871(1)		263(1)0.3(1)
Fe–V2(1)a = 2.9834(4)156(2)0.2(1)
1 h MA
Fe22(1)a = 2.8695(2)118(1)2.0(1)
Ti76(1)a = 2.9562(3)
c = 4.6908(1)		110(2)3.0(1)
Fe–V2(1)a = 2.9835(4)154(1)0.2(1)
3 h MA
Fe67(1)a = 2.8649(2)60(1)1(1)
Ti30(1)a = 2.9491(3)
c = 4.6825(1)		85(1)3(1)
Fe–V3(1)a = 2.9861(2)100 (2)4(1)
5 h MA
Fe60(1)a = 2.8613(2)56(1)2(1)
Ti38(1)a = 2.9456(3)
c = 4.6737(1)		86(1)4(1)
Fe–V2(1)a = 2.9737(2)99 (2)5(2)
Table 3. Phase fractions and refined crystallographic parameters of the phases present in the compacted and annealed sample of Ti0.9FeV0.06 alloy.
Table 3. Phase fractions and refined crystallographic parameters of the phases present in the compacted and annealed sample of Ti0.9FeV0.06 alloy.
PhaseFraction (wt%)Lattice Parameter (Å)Crystallite Size (nm)Microstrain × 10−3 (%)
FeTi90(1)a = 2.9767(2)118(1)1(1)
Fe2Ti4(1)a = 4.7478(2)
c = 8.0430(1)		99(1)3(4)
Fe2Ti4O6(1)a = 11.2725(1)123(1)1(1)
Table 4. Chemical composition (at%) of different phases using EDX analysis.
Table 4. Chemical composition (at%) of different phases using EDX analysis.
PhaseTi (at%)Fe (at%)V (at%)O (at%)
FeTi (spectrum 6)50.749.3--
Fe2Ti (spectrum 5)35.364.7--
Fe2Ti4O (spectrum 4)57.332.31.29.2
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MDPI and ACS Style

Patel, A.K.; Violi, D.; Lorenzon, I.; Luetto, C.; Rizzi, P.; Baricco, M. Synthesis of a Vanadium-Substituted Fe–Ti-Based Ternary Alloy via Mechanical Alloying, Compacting, and Post-Annealing. Metals 2025, 15, 723. https://doi.org/10.3390/met15070723

AMA Style

Patel AK, Violi D, Lorenzon I, Luetto C, Rizzi P, Baricco M. Synthesis of a Vanadium-Substituted Fe–Ti-Based Ternary Alloy via Mechanical Alloying, Compacting, and Post-Annealing. Metals. 2025; 15(7):723. https://doi.org/10.3390/met15070723

Chicago/Turabian Style

Patel, Abhishek Kumar, Davide Violi, Ivan Lorenzon, Carlo Luetto, Paola Rizzi, and Marcello Baricco. 2025. "Synthesis of a Vanadium-Substituted Fe–Ti-Based Ternary Alloy via Mechanical Alloying, Compacting, and Post-Annealing" Metals 15, no. 7: 723. https://doi.org/10.3390/met15070723

APA Style

Patel, A. K., Violi, D., Lorenzon, I., Luetto, C., Rizzi, P., & Baricco, M. (2025). Synthesis of a Vanadium-Substituted Fe–Ti-Based Ternary Alloy via Mechanical Alloying, Compacting, and Post-Annealing. Metals, 15(7), 723. https://doi.org/10.3390/met15070723

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