Instrumented indentation tests were performed to determine mechanical properties of the studied materials. Indentation force-penetration depth curves were evaluated by Oliver-Pharr method [20,21] and the hardness and Young's modulus were calculated according to standard ISO 14,577 [22] on NHT2 nanoindentation tester (Anton Paar, Graz, Austria). At least 15 particles from every powder mixture were tested to obtain statistically representative value of hardness and Young's modulus. As the hard powder particles are relatively small and embedded in the matrix with much lower hardness and Young's modulus, nanoindentation tests with partial unloadings at 1, 2, 5 and 10 mN were performed to determine the effect of the surrounding matrix on the measured values of hardness and Young's modulus of powder particles. It was found [23] that for the size of particles in the order of tens of micrometers, the values of hardness measured at the maximum load 2 mN (corresponding to the penetration depth of approximately 80 nm) were not affected by the surrounding resin and the scatter of the data was reasonably small. On the other hand, the measurements of Young's modulus of hard particles in such softer surrounding matrix were affected even for the penetration depths lower than 0.25% of the size of the particles. This means that in the case of the studied materials, penetration depths should be less than 50 nm, which is not possible because it will not meet the conditions of the ISO standard [22]. As a consequence, measured values of Young's modulus for powder particles were underestimated and they are not presented. Moreover, to avoid the effect of the particle size, which is changing with milling time [17], only the particles larger than 10 μm were tested to avoid the effect of the particle size in soft surrounding matrix on hardness measurements.

Based on this research, sintered samples were also tested at the maximum load of 2 mN in order to obtain the data for the same experimental conditions and to enable direct comparison of results.

To determine fracture toughness of the sintered materials, Vickers indentations with maximum load of 5 N were performed on MHT microindentation tester (Anton Paar, Graz, Austria) and the dimensions of radial cracks were measured (Figure 2). To evaluate the fracture toughness, model of Niihara [24] for Palmqvist cracks was employed. This model is based on the equation.

$$K\_{I\mathbb{C}} = 0.035 \left(\frac{l}{a}\right)^{-\frac{1}{2}} H a^{\frac{1}{2}} \frac{1}{\phi} \left(\frac{E\phi}{H}\right)^{\frac{2}{5}},\tag{1}$$

where *l* is the crack length, *a* is a half-diagonal of the indent, *H* is hardness, *E* stands for Young's modulus and φ is the constraint factor (ratio of hardness and yield stress reaching the value of approximately 3).

**Figure 2.** Geometry of the cracks created by the Vickers indentation [24].

#### **3. Results**

#### *3.1. Powders*

Typical evolution of the microstructure of the powders is presented in Figure 3. At the beginning of the milling process, coarse lamellas of the initial powders in the individual powder particles and the unchanged particles of the initial powders were present. With increasing milling time (2 h), the lamellas got thinner, the occurrence of particles unaffected by the milling process was still rarer and first homogeneous particles with required chemical composition were observed. After four hours of milling (Figure 4), the homogeneous particles were present for all four powder mixtures. For the mixture of pure elements Fe\_Al\_Si, the homogeneous particles completely replaced heterogeneous particles. For the mixtures prepared from pre-alloyed powders, particles with lamellar structure could still be observed. Complete homogenization of Fe\_AlSi and Fe\_SiAl powders was observed after 6 h of milling. Only the powder Fe\_AlSi was not completely homogenous and the particles with

lamellar structure were observed. The complete homogenization of FeAl\_Si powder occurred after 8 h of milling.

**Figure 3.** Typical example of the microstructure evolution with increasing time of mechanical alloying (Fe\_AlSi powder mixture, SEM BSE micrographs).

From the low magnification TEM micrographs of the Fe\_Al\_Si powder after MA for 8 h (not presented here) it was obvious, that only particles under 5 μm in size had stuck on the carbon film. The chemical composition of these small particles obtained from TEM-EDS was in the range (wt.%) 55% to 66% of Fe, 10% to 15% of Al, 13% to 16% of Si, and 0.7% to 4% of O. Two typical Fe-rich particles, agglomerates of very small grains, 20 to 100 nm in size, are presented in Figure 5a. According to simulations performed using JEMS software (version 4.7830U2019b11) by Stadelmann [25], the diameters of the rings in the related electron diffraction pattern (inset in Figure 5a) correspond to the Fe3Si phase (*Fm*3*m*, space group 225, a = 0.5655 nm).

**Figure 4.** SEM BSE micrographs of powders after 4 h of mechanical alloying: (**a**) Fe\_Al\_Si, (**b**) FeAl\_Si, (**c**) Fe\_AlSi, (**d**) FeSi\_Al.

Characteristic change of phase composition during mechanical alloying is shown in Figure 6 in the example of the Fe\_AlSi powder mixture. At the beginning of the alloying process, the initial powders of the respective mixtures (i.e., Fe, Al, Si, and phases FeAl, Fe3Al, FeSi, Fe3Si in pre-alloyed powders) could be identified. With increasing milling time, the peaks of pure elements (Fe, Al and Si) disappeared and the peaks of the phases FeSi, Fe3Si and FeAl occurred and started to prevail for higher milling times. The change rate was different for different powder mixtures (Figure 7). At the end of the mechanical alloying process, only the phases FeSi, Fe3Si (supersaturated by Al) and small fraction of FeAl were present. The supersaturation of the phases FeSi and Fe3Si by Al atoms caused the increase of lattice parameters and shift of the peaks to the lower diffraction angles.

**Figure 5.** *Cont*.

**Figure 5.** (**a**–**d**) Bright field TEM micrographs of: (**a**) nanocrystalline Fe3Si particles in the loose powder mixture from pure elements (Fe\_Al\_Si) after 8 h of mechanical alloying (related ring diffraction pattern is in the inset), (**b**–**d**) corresponding spark plasma sintering (SPS) compact crushed in agate mortar: (**b**) refined crystallites of the Fe3Si phase, (**c**) another (unidentified) phase with nanometric grains (**d**) amorphous phase, (**e**) STEM HAADF micrograph of a focused ion beam (FIB) lamella prepared from the same SPS compact, (**f**) STEM EDS line analysis across dark alumina particles in the micrograph (**e**).

For Fe\_Al\_Si powder mixture from pure elements, the phases FeSi, Fe3Si and FeAl occurred after 2 h of milling and the peaks of original pure powders vanished after 4 h when the homogenization process was completed. The phases FeSi and Fe3Si, which formed the final powder mixture, started to occur after 4 h of milling for the mixture FeAl\_Si. The phases of the initial powders (i.e., FeAl, Fe3Al and Si) were present even after 6 h of mechanical alloying and only after 8 h were not observed. The final phases (FeSi, Fe3Si and FeAl) for the mixture Fe\_AlSi were firstly observed after 2 h of the mechanical alloying and the pure elements were lastly observed after 4 h of milling. The phases FeSi and Fe3Si in powder mixture FeSi\_Al were present already from the beginning of milling because they formed the pre-alloyed powder; other phases disappeared after 4 h of milling.

**Figure 6.** Typical example of the phase composition evolution with increasing time of mechanical alloying (Fe\_AlSi powder mixture).

**Figure 7.** *Cont*.

**Figure 7.** Phase composition of the powder mixtures in different stages of mechanical alloying.

Hardness values evaluated for maximum load 2 mN are presented in Figure 8. Values presented for 0 h correspond to initial pre-alloyed powders (i.e., FeAl25, AlSi30, FeSi25). Except FeSi\_Al powder mixture, hardness increased at the beginning of the milling process. For Fe\_Al\_Si mixture, the significant increase ended after 2 h and the hardness stayed constant. For FeAl-Si mixture the increase was slower and more progressive up to the end of mechanical alloying at 8 h. Hardness of Fe\_AlSi mixture started from the lowest values, but it increased very rapidly and it reaches the stable value after 2 h. On the other hand, mixture FeSi\_Al started at very high hardness values (hard pre-alloyed powder FeSi25), hardness decreased to the minimum after 2 h after which it increased and it reached the constant value after 4 h. All the powder mixtures had hardness of approximately 15 GPa at the end of the mechanical alloying.

**Figure 8.** Hardness of the powder mixtures in different stages of mechanical alloying (solid horizontal line shows the value of the compacted sample measured at the maximum load of 2 mN, dotted lines represent the error bars).

#### *3.2. Compacted Materials*

An example of the microstructure of the sintered sample by the SPS method is shown in Figure 9. It was homogeneous with a very low porosity (less than 0.1% for Fe\_Al\_Si sample [19,26], relative density nearly 100%, and fine microstructure (the grain size was around 1 μm for all four sintered samples—Table 1). There were not observed any significant differences between different initial powder mixtures as the powder mixtures used for SPS were of the same chemical and phase composition after MA.

**Figure 9.** Microstructure of the compacted powder Fe\_AlSi.

**Table 1.** Average grain size and fracture toughness of the samples prepared by SPS (measured at the maximum load of 5 N).

From TEM analysis of the corresponding powder from the crushed Fe\_Al\_Si SPS compact it follows that SPS led to refining of the crystallites of the Fe3Si phase to nanometric size (Figure 5b), probably due to recrystallization (micrographs and diffraction patterns in Figure 5a,b are to scale). Isolated particles of other phases were also observed (Figure 5c). However, the crystallographic identification of these phases was not successful. Furthermore, amorphous nanoparticles were found (Figure 5d), formed possibly due to the local liquid phase sintering and rapid cooling of the compact. These amorphous phases may be at the origin of the high brittleness of the compact. According to EDS analysis, the chemical composition of the particles from the crushed compact was practically the same as of the loose MA powder.

From the scanning transmission electron microscopy (STEM) imaging of the FIB lamella prepared also from the Fe\_Al\_Si SPS compact using high-angle annular dark field detector (HAADF) it follows, that the grain size ranges from 100 nm up to 3 μm (Figure 5e). Many small particles were observed inside the grains and also on grain boundaries. The STEM HAADF signal is proportional to the scattering power (Z-contrast, where Z is the proton number); hence heavy elements appear bright, and holes and light particles appear dark. Therefore, the differences in the gray level of the grains in Figure 5e indicate a quite heterogeneous Fe concentration—grains appearing bright contain more iron, grains appearing dark contain more Al and Si. It can also be seen that the SPS compact includes a very high density of dark oxide particles, mostly 5 to 100 nm in size. These oxides are distributed inside of the majority of the grains and also on grain boundaries (Figure 5e). Figure 5f presents results of line

STEM-EDS analysis across the particles, the red line in the inset indicating the path of the electron beam. From the plot it is obvious that majority of particles are oxides of aluminum. Therefore, argon flushing atmosphere during mechanical alloying did not prevent alumina formation. TEM analyses of the powders and compacts of the other three materials (FeAl\_Si, Fe\_AlSi, and FeSi\_Al gave almost the same results, so they are not presented here).

XRD patterns are in Figure 10 and they show only small changes compared to the mechanically alloyed powders. The XRD revealed that the sintered materials were predominantly composed of FeSi and Fe3Si phases (supersaturated by Al). Contrary to the MA powders the peaks of ternary phase Fe3Al2Si3 were identified in the sintered samples (compare Figures 7 and 10).

Hardness measured for maximum load of 2 mN is represented in Figure 8 by solid horizontal line (dotted horizontal lines show the scatter of the data). The measured hardness was about 15–16 GPa and it was in good agreement with the values reached by the mechanically alloyed powders after 8 h of milling.

The fracture toughness was determined by measuring hardness, Young's modulus and crack dimensions formed by the Vickers indentation to the maximum load of 5 N. The results are summarized in Table 1 and observed crack systems are shown in Figure 11. The measured hardness was about 13 GPa and Young's modulus 225 GPa. Fracture toughness calculated using Equation (1) from measured radial (Palmqvist) cracks was about 3 MPa·m1/2. It was lowest for the Fe\_Al\_Si sample reaching only 2.57 MPa·m1/2. The highest fracture toughness was found for the sample FeAl\_Si (3.61 MPa·m1/2). The SEM images confirmed that the cracks were not connected under the indent, which means that they were radial (not median) cracks. The system of the radial cracks was well developed for the samples Fe\_Al\_Si and Fe\_AlSi. For the samples FeSi\_Al and FeAl\_Si, the systems with only three cracks were also sometimes observed. These indents were excluded from the analysis and only the systems with four cracks were analyzed. On the other hand, lateral cracks were occasionally observed for the samples FeSi\_Al and FeAl\_Si with higher fracture toughness, which could be the reason for the increase in *KIC* (change of the mechanism of cracking associated with plastic deformation).

**Figure 10.** Phase composition of the sintered samples.

**Figure 11.** Cracks propagating from the Vickers indentation: (**a**) Fe\_Al\_Si, (**b**) FeAl\_Si.

#### **4. Discussion**

The results presented in the previous section show that the kinetics of the mechanical alloying process depend on the initial powders which were mixed to obtain final homogeneous product. It can be stated that the microstructure, phase composition, hardness and Young's modulus at the end of the mechanical alloying and after the spark plasma sintering are very similar. However, the kinetics of the formation of the homogeneous powder is different. All the measurements show that the fastest homogenization is for the Fe\_Al\_Si mixture made of pure elements. Full homogenization is reached after 4 h of mechanical alloying. 6 h are sufficient to homogenize the powders Fe\_AlSi and FeSi\_Al. The longest time (8 h) is required for the preparation of FeAl\_Si mixture.

These differences are caused by the properties of the initial powders. For fast homogenization, it is advantageous to combine hard brittle and soft ductile powders. Ductile particles are easily plastically deformed and create cold welds. Brittle particles fracture, they get smaller and they can easily get inside soft plastic particles [11]. Only a small amount of the brittle hard particles causes fast increase in the hardness and accelerates the homogenization. This is the case for Fe\_Al\_Si and Fe\_AlSi mixtures, in which Si has the role of the hard and brittle material in soft ductile Fe, Al, or AlSi30 pre-alloyed powder, respectively. Very fast increase in hardness of Fe\_AlSi alloy during the first two hours of mechanical alloying can be caused by Si fine dispersion in the Al-Si eutectic. Increased temperature during mechanical alloying together with the fine Si eutectic particles in Al matrix accelerates the reactions with Fe and the formation of solid solution at the beginning of mechanical alloying. At the later stages of mechanical alloying, the primary Si acts similarly as the pure Si powder and the kinetics of the MA is equivalent with the case of the mixture of pure elements.

For the mixture FeSi\_Al, the dominant component is hard FeSi25 intermetallic alloy. During the first two hours, the predominant hard particles are mixed with soft aluminum which causes the decrease in the hardness. Hardening processes are not fast enough to compensate this decrease caused by initial mixing. After 2 h of milling, homogenization occurs, hardness increases and the process continues in the same way as for Fe\_Al\_Si and Fe\_AlSi mixtures.

The case of FeAl\_Si mixture is fundamentally different because two hard brittle components are mixed. The brittle particles are preferentially crushed during the collisions and the formation of cold welds is more complicated. The diffusion between the components and the plastic deformation is lower and the homogenization takes a longer time.

As it was mentioned above, the particle size is slightly changing with milling time (from 30 to about 10 μm on Fe\_Al\_Si [17]). Only the particles larger than 10 μm were tested to avoid the effect of the particle size in soft surrounding matrix on hardness measurements. No substantial effect of particle size on hardness values was however noted, which is in agreement with the values measured on sintered samples (where the whole particle size distribution intervenes). Thus, the particles size distribution should not have a significant effect on the presented values of the hardness. As a result, the changes in hardness should correspond only to the changes in microstructure and homogenization of the powder particles caused by mechanical alloying.

Increase in hardness and the saturation on the constant value during mechanical alloying is generally observed [27,28]. The hardening rate depends on the kinetic energy of the collisions which can be driven by the milling parameters such as rotational speed, powder to ball ratio (material of the balls), etc. The lower rotational speed or lighter balls cause slower increase in hardness and require more time for the completion of the MA process [27]. The increase in hardness of the MA powders can be caused by different mechanisms. The first is work hardening when the particles are plastically deformed during the collisions with balls causing the increase of dislocation density [29]. The second mechanism is the grain refinement (Hall-Petch hardening) [28,30]. The next strengthening mechanism is dispersion hardening where the nanoparticles blocking the dislocation movement are homogenously distributed in the matrix. This type of hardening is typical for oxide dispersion strengthened steels [12]. The last type is the solid solution hardening where the atoms from the crystal structure are replaced by other atoms with not exactly the same atomic radius, causing the distortion of the lattice (the case of Al and Si atoms in the studied material). All these mechanisms can cause the increase in hardness observed in this study. When the equilibrium with the recovery mechanisms is established, the hardness does not increase more and saturates at the final level.

As it was discussed in previous paragraphs, the kinetics of the mechanical alloying and thus the strengthening mechanisms depend on the initial combination of the powders. If at least one component is ductile, the hardening mechanisms based on the plastic deformation and dislocation movement are easier and the hardening is faster. Increasing the amount of hard brittle components causes cracking and the particles cannot be easily plastically deformed. Thus, the hardening and the final homogenization is more time-demanding.

After the MA, FeSi and Fe3Si are predominant phases in homogeneous powders instead of the stable ternary phases [31]. It is caused by the easy contact of two particles during the milling and increased solubility of the elements caused by mechanical alloying [32]. Other phases found for the alloys with different Al/Si ratio (e.g., Fe2Al5 [17]) were not identified. Spark plasma sintering was chosen for the consolidation of the mechanically alloyed powders because it is a fast method which prevents significant changes in microstructure, grain growth and mechanical properties. Nevertheless, the presence of stable ternary phase Fe3Al2Si3 was observed after the sintering in all the samples. Exact phase ratio for Fe\_Al\_Si mixture was determined in the previous study [19] (33.0 wt.% of Fe3Si, 37.5 wt.% of FeSi, 29.5 wt.% of Fe3Al2Si3) and according to the diffractograms in Figure 10, there is no significant difference between various mixtures. This ternary phase usually complicates the processing of this type of materials by conventional methods because of its low plasticity. It would therefore be desirable to prevent this phase as well as the amorphous brittle phases before final processing. Compared to other initial compositions or processing ways (self-propagating high-temperature synthesis), the presence of the phases Fe2Al5 [8] or FeAl2Si [9], respectively, was not observed.

As can be seen from Figure 8, hardness after mechanical alloying was not affected by the sintering process and it stays very high (15 GPa). Hardness of compacted samples measured for maximum load of 5 N is slightly lower than the value measured for the load of 2 mN. This can be caused by several reasons which could include porosity of the sample (even if it was very low), cracking at higher loads or indentation size effect that could already be present at used low loads [33].

A wide range of formulas for different crack systems and materials exists for the determination of the fracture toughness by indentation methods (e.g., [34–39]). It was confirmed by SEM observations that the system of cracks in the studied material is radial (Palmqvist cracks). For this reason, Equation (1) which is valid in the range of crack lengths 0.25 ≤ *l*/*a* ≤ 2.5 was employed. The fracture toughness is quite low, in the range of ceramics [40], which could be caused by the high amount of stored plastic deformation and/or the presence of brittle amorphous phases. On the other hand, the manufacturing process performed in this study (i.e., mechanical alloying and spark plasma sintering) leads to significant fracture toughness improvement compared to standard casting of these materials (about 10 times [19]). The lowest fracture toughness was found for the mixture Fe\_Al\_Si which was homogenized in the shortest time. The highest fracture toughness was measured for FeAl\_Si mixture, for which the homogenization took the longest time of 8 h. This means that the fast homogenization leads to a decrease of the fracture toughness. It can be caused by the stored plastic deformation in the powder particles. When the homogenization process is finished soon (e.g., after 4 h), no significant changes happen during further milling and only plastic deformation is stored in the particles. This means that from the point of the fracture toughness, the reduction of the milling time to the shortest time necessary for the homogenization of the powder would be optimal.

Other reasons for the differences in fracture toughness could be different levels of supersaturation of Fe3Si and/or FeSi phases by aluminum or different volume fraction of amorphous phase. All aluminum in solid solution in FeAl pre-alloyed powders could suppress the local liquid phase sintering. However, the TEM observation does not allow the exact quantification of the amorphous phase content in different pre-alloyed powders. On the other hand, no correlation between fracture toughness and grain size could be found as the grain size is practically the same for all pre-alloyed powders, which is probably closely connected with the saturation of the plastic deformation during milling.

#### **5. Conclusions**

The presented results show the differences in the kinetics of the mechanical alloying process depending on the initial feedstock materials used for the production of FeAl20Si20 alloy. The results can be summarized as follows:

**Author Contributions:** Conceptualization, J.C. and P.H.; methodology, J. ˇ C., P.H., F.P. and P.N.; formal analysis, ˇ J.C., M.K. and P.H.; investigation, J. ˇ C., P.H., M.K., V.B. and K.N.; resources, K.N., J.K.; data curation, J. ˇ C.; ˇ writing–original draft preparation, J.C.; writing–review and editing, P.H.; funding acquisition, F.P. ˇ

**Funding:** The funding by Czech Science Foundation (grant No. 17-07559S) and European Regional Development Fund in the frame of the project Centre of Advanced Applied Sciences (No. CZ.02.1.01/0.0/0.0/16\_019/0000778) is gratefully acknowledged.

**Acknowledgments:** This research was financially supported by Czech Science Foundation (grant No. 17-07559S) and European Regional Development Fund in the frame of the project Centre of Advanced Applied Sciences (No. CZ.02.1.01/0.0/0.0/16\_019/0000778).

**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.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Mechanism of the Intermediary Phase Formation in Ti-20 wt. % Al Mixture during Pressureless Reactive Sintering**

#### **Andrea Školáková 1,\*, Pavel Salvetr 1, Pavel Novák 1, Jindˇrich Leitner 2 and Davy Deduytsche 3**

Received: 7 June 2019; Accepted: 4 July 2019; Published: 6 July 2019

**Abstract:** This work aims to describe the mechanism of intermediary phases formation in TiAl20 (wt. %) alloy composition during reactive sintering. The reaction between titanium and aluminum powders was studied by in situ diffraction and the results were confirmed by annealing at various temperatures. It was found that the Ti2Al5 phase formed preferentially and its formation was detected at 400 ◦C. So far, this phase has never been found in this alloy composition during reactive sintering processes. Subsequently, the Ti2Al5 phase reacted with the titanium, and the formation of the major phase, Ti3Al, was accompanied by the minor phase, TiAl. Equations of the proposed reactions are presented in this paper and their thermodynamic and kinetic feasibility are supported by Gibbs energies of reaction and reaction enthalpies.

**Keywords:** in situ diffraction; aluminides; reactive sintering; mechanism; powder metallurgy

#### **1. Introduction**

A Ti-Al system consists of five important phases, including a Ti3Al compound with a hexagonal close-packed superlattice (space group P63/mmc), an equiatomic TiAl compound with a tetragonal structure (space group P4/mmm), and aluminum-rich intermetallic compounds, namely TiAl2 (space group I41/amd), Ti2Al5 (space group P4/mmm), and TiAl3 (space group I4/mmm), also with a tetragonal structure. All intermetallic compounds are collectively called titanium aluminides. Titanium aluminides belong to the group of innovative materials that gradually replace nickel-based superalloys in highly demanding applications [1]. They possess a great combination of stable mechanical properties at high temperatures (500–900 ◦C), low density, and good oxidation resistance. For this reason, they are suitable candidates as structural materials for the aerospace and automotive industries. So far, they have been used in turbocharger wheels and turbine blades. The current research is mainly focused on the development of alloys with microstructures containing TiAl and Ti3Al phases, which should ensure great creep resistance [1–5]. Despite this advantage, the application range of titanium aluminides is still limited because they suffer from room-temperature brittleness and have poor melt-metallurgic properties [5]. Moreover, the extreme reactivity of molten titanium usually causes contamination of the obtained products [6]. These compounds are produced by a melt-metallurgy process comprising vacuum induction melting (VIM), vacuum arc remelting (VAR), centrifugal casting, conventional melting, and hot isostatic pressing (HIP) [1]. Powder metallurgy techniques are much more suitable because melting is avoided. Thus, molten titanium does not occur, and a high purity of the products can be obtained [1]. One of the currently studied processes of these compounds is Self-propagating

High-temperature Synthesis (SHS) which uses solid state diffusion to obtain a mostly local, usually eutectic, composition and a subsequent formation of the liquid phase at temperatures lower than the melting points of the initial materials. This mechanism is, thus, different from reactive sintering using only solid-state diffusion. However, SHS is taken as a specific type of reactive sintering.

The SHS process is a modern method of sintering because time and energy are saved [7,8]. SHS lies in the heating of initial reactants, which are usually powders [9]. Highly exothermic reactions occur during heating [10]. Heat released by these reactions is supported other exothermic reactions, and, thus, the SHS reaction can propagate through the whole heated powder mixture spontaneously [9,11]. For this reason, this reaction is often called as "self-sustaining". Reactions are usually thermally activated. SHS reaction is divided into two modes, which differ in their heating initiation. A reaction can be initiated by heating of the whole sample, and then it is called Thermal Explosion mode or by heating of only one side of sample and this case of heating is called Ignition mode. The thermal explosion mode usually involves homogeneous heating, while the Ignition mode uses extremely fast heating in most cases [12]. Both modes have their own disadvantages. For the former mode, pre-combustion is necessary and can modify the reaction mechanism. On the other hand, this mode is more accurate in the determination of reaction temperatures. The latter mode is faster, but the results have poor reliability [12].

TiAl and TiAl3 aluminides prepared by SHS reaction are the most studied ones [10,13–16]. Ti3Al is not a commonly studied aluminide from the viewpoint of the SHS process, and, hence, its mechanism of phase formation has not yet been described. On the other hand, published research shows contradictory or incomplete results, and the description of the reaction mechanisms of phase formation is not united for TiAl and TiAl3 alloy composition. It is known that the reaction between titanium and aluminum powders is initiated when the temperature reaches the melting point of aluminum during heating [17]. The TiAl3 phase is concluded to be the first phase during the reaction between liquid aluminum and solid titanium, followed by the formation of the TiAl and Ti3Al phases [13,16–19]. It is believed that the TiAl3 phase forms preferentially due to its thermodynamic and kinetic system preferences [20]. The TiAl2 phase should form after the formation of the TiAl and Ti3Al phases by interdiffusion between the TiAl3 phase and titanium [2]. Other works showed that the TiAl phase formed first after a reaction between the melted aluminum and titanium [21]. Further, the TiAl phase is considered to be the starting phase before the formation of the Ti2Al5 and TiAl2 phases [20], followed by the formation of the TiAl3 phase. The explanation for the relationship between phases can be found in their low free energies [19]. However, other works showed that a very thin layer of the Ti3Al phase formed first on the titanium side surrounded by the TiAl phase and TiAl3 formed as the last one [22]. This result shows a completely different phase formation sequence in comparison to the results mentioned before. For this reason, there are holes in the description of reaction mechanisms, and it is necessary to determine the accurate reaction conditions of phase formation.

Ti and Al foils, instead of powders, have also been intensively studied, mainly because of their better mechanical properties. The aim is to synthesize the composite as a sandwich. For this reason, foils are most often produced by electron beam deposition, which affects the phase composition of the interface between aluminum and titanium, and, subsequently, these foils are heated. It was found that diffusion welding formed. Authors usually observe the growth of the intermetallic layer while both reactants are in the solid state and obtained foundational results that they could use determine kinetics. However, they could study only the interface between aluminum and titanium whose chemical composition corresponds to 50:50 in at. % of Ti:Al [23–27]. For this reason, the reaction between solid aluminum and solid titanium is better to simulate the annealing of a compressed powder mixture at temperatures lower than the melting point of aluminum (used in this work). Advantage is possible to mix a different ratio of aluminum to titanium and study Ti3Al, TiAl, or TiAl3 phases. Moreover, the powder mixture is only compressed, and, thus, the interface between aluminum and titanium is not affected by pre-heating.

In this work, a Ti-20 wt. % Al powder mixture (corresponding to Ti3Al compound) was studied. This chemical composition was deliberately chosen because this work is part of project in which the sintering of Ti3Al, TiAl, and TiAl3 compounds has been studied. The amount of aluminum (20 wt. %) lies in the center of the stability range of the Ti3Al phase, according to diagram Ti-Al [28]. This mixture was subjected to the Thermal Explosion mode (TE) SHS, and the mechanism of phase formation was described. In situ diffraction was applied to determine the sequence of phase formation, and their formation was observed at the highest heating rate that the in situ XRD (X-ray diffraction) device allows. To determine whether the phase formed below the melting point of aluminum, powder mixtures were annealed at various temperatures and times. This enabled us to show and confirm the phases' formation sequence. This work offers the first detailed descriptions of the phase formation of this alloy composition.

#### **2. Materials and Methods**

All tested samples were prepared from blended powders of titanium (with a purity of 99.5% and a particle size of 44 μm, STREM CHEMICALS, Newburyport, MA, USA) and aluminum (99.62%, 44 μm, STREM CHEMICALS, Newburyport, MA, USA). Powder blends corresponding to the Ti3Al compound, i.e., TiAl20 (in wt. %), with a weight of 3 g, were prepared. Subsequently, powder mixture was uniaxially cold pressed at an ambient temperature to cylindrical green bodies of 10 mm in diameter by a pressure of 450 MPa for 5 min using the LabTest 5.260SP1-VM universal loading machine (Labortech, Opava, Czech Republic). Reaction kinetics, especially the initiation of the reactions, strongly depends on the green density of the compacts. We have studied intermetallics for several years and we have experienced that reaction kinetics mainly depend on particle size and compaction pressure. For this reason, we used the finest powders of titanium and aluminum and a compaction pressure of 450 MPa to obtain the powder mixture with the best contacts of particles and the lowest possible porosity.

Afterwards, the sample was inserted into the induction furnace and heated under a protective Ar atmosphere. Heating was recorded by an optical pyrometer Optris OPTP20-2M (Optris, Portsmouth, NH, USA) to observe the emerging exothermic reaction in the Thermal Explosion (TE-SHS) mode. The heating rate was determined, according to the slope of the obtained curve, as 109 ◦C·min−1. In order to describe the phases' formation sequence during the SHS process, in situ XRD analysis (Department of Solid State Sciences, Ghent University, Ghent, Belgium) of the compressed powder mixture was performed under an He atmosphere during heating to 900 ◦C, with a heating rate of 60 ◦C·min−1. XRD source and linear Vantec detector were fixed at the positions of 21◦ (source) and 42◦ (detector).

Further, the formation of phases was observed at temperatures lower than the melting point of aluminium −400, 450, 500, and 600 ◦C. Samples were evacuated in silica ampoules and exposed at these temperatures for 8, 24, and 48 h.

The obtained samples were ground by sandpaper P80–P4000 with SiC abrasive particles (Hermes Schleifmittel GmbH, Hamburg, Germany), polished by suspension of colloidal silica Eposil F (ATM GmbH, Mammelzen, Germany) with hydrogen peroxide (volume 1:6), and etched by Kroll´s reagent (5 mL HNO3, 10 mL HF, 85 mL H2O). The microstructure was examined by a scanning electron microscope TESCAN VEGA 3 LMU (Tescan, Brno, Czech Republic) equipped with an OXFORD Instruments X-max EDS SDD 20 mm2 detector (Oxford Instruments, High Wycombe, UK) for identification of the chemical composition of individual phases (SEM-EDS). Phase composition was determined by X-ray diffraction using the PANalytical X´Pert Pro diffractometer with CuKα radiation (PANalytical, Almeo, The Netherlands).

#### **3. Results and Discussion**

Figure 1 shows the heating curve obtained during the reactive sintering of the Ti-20 wt. % Al powder mixture in an induction furnace. The exothermic peak is associated with the SHS reaction and formation of phases. The reaction started above the temperature of the melting point of aluminum because Tonset was 716 ◦C, indicating that the presence of a liquid phase triggered the reaction. This means that the reaction was not initiated at the melting point of aluminum at 660 ◦C, as is typical for aluminides [16]. The liquid phase was present, accelerated the sintering at higher temperatures, and was gradually consumed during the process [2]. The maximum SHS reaction, called the combustion temperature, was 760 ◦C. The microstructure was composed of unreacted titanium, which was surrounded by large areas of the Ti3Al phase (Figure 2). TiAl was found between the layers of the Ti3Al phase. These phases were also detected by XRD analysis (Figure 3), and, moreover, the Ti2Al5 phase was found. This phase is very fine and is probably dispersed in the matrix uniformly. According to the area fraction, it can be assumed that the Ti3Al phase formed during an exothermic SHS reaction, which was observed by an optical pyrometer (Figure 1). This phase also formed as a major phase in Ti:Al = 1:1 system [16], and it was stated that its formation is associated with diffusion. On the other hand, the Ti2Al5 phase was not detected in this aluminum enriched system [16]. Results from the EDS analysis are shown in Table 1.

**Figure 1.** Heating curve obtained by the optical pyrometer at a heating rate of 109 ◦C·min−1.

**Figure 2.** SEM image of microstructure obtained after induction at heating rate of 109 ◦C·min−1.

**Figure 3.** XRD patterns obtained after induction at a heating rate of 109 ◦C·min−1. **Table 1.** SEM + EDS point analysis results of the reactively sintered TiAl20.

The in situ XRD analysis shows the continuous formation of phases from the powder mixture during heating (Figure 4). This observation of phase composition during heating helps to describe the mechanisms of phase formation. It can be seen that titanium and aluminum were present up to approximately 500 ◦C. At this temperature, the Ti2Al5 phase started to form. Titanium and aluminum diffraction lines disappeared at approximately 700 ◦C. This disappearance is associated with the formation of other phases—Ti3Al and TiAl. These phases were accompanied by the Ti2Al5 phase, which had formed earlier. It was impossible to determine which of the phases (Ti3Al or TiAl) formed preferentially. All their diffraction lines appeared at the same temperature, i.e., approximately 700 ◦C. The intensity of the lines of the Ti2Al5 phase was distinctive from 500 ◦C to 700 ◦C. This result implies that the SHS reaction between the titanium and aluminum powders is initiated by the formation of an intermediary Ti2Al5 phase. Many works [2,16–20] claimed that the TiAl3 phase formed before the formation of the TiAl, Ti3Al, or TiAl2 phases already at 580 ◦C. Their explanation as to why the TiAl3 phase arose preferentially lies in its minimum free energy, suggesting that its formation is thermodynamically and kinetically favored over the formation of other aluminides. Only the presence of liquid aluminum is sufficient for the formation of the TiAl3 phase [20]. On the other hand, in situ XRD analysis, and our other results, showed that the Ti2Al5 phase is formed preferentially in a TiAl20 powder blend. The reaction temperature is an important parameter that can support or induce synthesis of aluminides, as described in work [29]. Based on the presented results, the other part of this study focused on individual temperatures to clarify aluminide formation.

From the obtained results, it can be seen that the Ti2Al5 phase formed below the melting point of aluminum. For this reason, mixtures of powders were annealed at various temperatures, and times, microstructures, and phase compositions were observed. Figure 5a–c shows microstructures obtained after annealing at 400 ◦C. Unreacted particles of aluminum and titanium were found. However, XRD analysis (Figure 6) revealed that Ti2Al5 formed at this temperature, and this phase was observed at the interface of titanium and aluminum. In situ XRD analysis detected its formation at 500 ◦C, but analysis was performed at a heating rate of 60 ◦C·min−1 (the maximum heating rate that can be set), and it is

known that with a decreasing heating rate, the initiation temperatures also decrease [30]. Moreover, the heating during the in situ XRD experiment was rather quick and continuous. Therefore, the amount in the Ti2Al5 phase, which is expected to form by a diffusion mechanism, is very low during continuous heating, and, therefore, it may not be detected. During annealing, the detectable amount of the Ti2Al5 phase was formed after 8 h at 400 ◦C (Figure 6). The EDS results are listed in Table 2. It can be seen that chemical composition does not correspond to Ti2Al5 but the analyzed area is small and, thus, strongly affected by titanium.

**Figure 4.** In situ diffraction obtained during heating (heating rate of 60 ◦C·min−1).

**Figure 5.** SEM images of a microstructure of a TiAl20 powder mixture annealed at 400 ◦C: (**a**) 8 h; (**b**) 24 h; (**c**) 48 h.

**Figure 6.** XRD patterns of TiAl20 powder mixture annealed at 400 ◦C.

**Table 2.** SEM + EDS point analysis results of TiAl20 annealed at 400 ◦C.

Annealing at a higher temperature (450 ◦C) for 8 h allowed the Ti2Al5 phase to be observed in the microstructure more markedly (Figure 7a). The microstructure was, again, composed of unreacted particles. One day annealing affected the microstructure significantly. The Ti2Al5 phase was detected by EDS analysis (see Table 3) on the interface between the titanium and aluminum powders (Figure 7b). With prolongation of annealing time, the areas of occurrence of the Ti2Al5 phase were larger (Figure 7c). XRD analysis confirmed the presence of hexagonal titanium, aluminum, and the Ti2Al5 phase (Figure 8).

**Figure 7.** SEM images of the microstructure of the TiAl20 powder mixture annealed at 450 ◦C: (**a**) 8 h; (**b**) 24 h; (**c**) 48 h.

**Table 3.** SEM + EDS point analysis results of TiAl20 annealed at 450 ◦C.

**Figure 8.** XRD patterns of TiAl20 powder mixture annealed at 450 ◦C.

The area of the Ti2Al5 phase was found during EDS analysis of the microstructure after annealing at 500 ◦C for 8 h (Figure 9a), but prolongation of times clearly revealed its location. The Ti2Al5 phase formed on the interfaces between aluminum and titanium particles, as well as at lower temperatures (Figure 9b,c). EDS analysis also confirmed its presence (Table 4). The phase composition was confirmed by XRD analysis (Figure 10).

**Table 4.** SEM + EDS point analysis results of TiAl20 annealed at 500 ◦C.

A temperature of 600 ◦C changed the microstructure significantly. Unreacted aluminum was not detected after all annealing times (Figure 11a–c). The Ti3Al and TiAl phases were determined by a combination of EDS (Table 5) and XRD analysis (Figures 11a–c and 12). Ti3Al replaced the Ti2Al5 phase and formed an interface between the original titanium particles. The TiAl phase was found only in some interfaces, and these interfaces were always darker (Figure 11a–c). In situ diffraction (Figure 4) and the results from microstructure observation showed that the Ti3Al and TiAl phases formed simultaneously, but the Ti3Al phase is the major phase. XRD analysis detected the Ti2Al5 phase, which was still present and did not disappear after formation of phases enriched by titanium (Figure 12). Ti2Al5 occurring in the microstructure was not detected by EDS analysis, likely due to its small dimensions and low volume fraction. This means that the Ti3Al phase really formed as the main phase during the SHS reaction. Its formation was accompanied by the formation of the minor phase, *Materials* **2019**, *12*, 2171

TiAl. The Ti2Al5 phase formed preferentially as a metastable reaction intermediate. The Ti2Al5 and TiAl phases both have a tetragonal structure [2], and this could be one of the reasons why the TiAl phase could be stabilized as a minor phase.

**Figure 9.** SEM images of the microstructure of TiAl20 powder mixture annealed at 500 ◦C: (**a**) 8 h; (**b**) 24 h; (**c**) 48 h.

**Figure 10.** XRD patterns of TiAl20 powder mixture annealed at 500 ◦C.

**Figure 11.** SEM images of the microstructure of TiAl20 powder mixture annealed at 600 ◦C: (**a**) 8 h; (**b**) 24 h; (**c**) 48 h.

**Table 5.** SEM + EDS point analysis results of TiAl20 annealed at 600 ◦C.

**Figure 12.** XRD patterns of TiAl20 powder mixture annealed at 600 ◦C.

The probable reaction scheme, with the Gibbs energies of reaction and reaction enthalpies, is as follows (Equations (1)–(3)):

$$2\text{ Ti} + 5\text{ Al} \rightarrow \text{Ti}\_2\text{Al}\_5, \qquad \Delta\text{H} \text{ (400 °C)} = -271.06\text{ kJ}, \quad \Delta\text{G} \text{ (400 °C)} = -226.29\text{ kJ} \tag{1}$$

Ti2Al5 + 13 Ti → 5 Ti3Al, ΔH (600 ◦C) = −293.81 kJ, ΔG (600 ◦C) = −242.84 kJ (2)

$$\text{Ti}\_2\text{Al}\_5 + 3\text{ Ti} \to 5\text{ TiAl}, \quad \Delta\text{H} \text{ (600 °C)} = -72.38 \text{ kJ}, \quad \Delta\text{G} \text{ (600 °C)} = 15.75 \text{ kJ} \tag{3}$$

Our calculations are derived from the thermodynamic data published in [31,32].

#### **4. Conclusions**

According to the obtained results, the mechanism of phase formation in Ti-Al systems can be described. The SHS reaction started at temperatures higher than the melting point of aluminum. It was found that Ti2Al5 phase formed preferentially, even though a titanium-rich powder mixture was investigated. The Ti2Al5 phase formed at 400 ◦C by a diffusion-controlled reaction. This phase subsequently reacted with the titanium, and the Ti3Al and TiAl phases formed simultaneously. These two phases could form only above 600 ◦C. The Ti3Al phase was a major phase. The results were compared between the data obtained by in situ diffraction and the annealed samples, and the described mechanism was confirmed.

**Author Contributions:** A.Š. was responsible for the theoretical background, writing the article and data analysis. P.S. prepared samples for metallographic observations. P.N. was responsible for final correction. J.L. was responsible for calculations. D.D. provided in situ measurements.

**Funding:** This research was financially supported by specific university research (MSMT No 21-SVV/2019).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Properties Comparison of Ti-Al-Si Alloys Produced by Various Metallurgy Methods**

#### **Anna Knaislová 1,\*, Pavel Novák 1, Jaromír Kopeˇcek 2 and Filip Pr ˚uša 1**

Received: 9 September 2019; Accepted: 19 September 2019; Published: 21 September 2019

**Abstract:** Melting metallurgy is still the most frequently used and simplest method for the processing of metallic materials. Some of the materials (especially intermetallics) are very difficult to prepare by this method due to the high melting points, poor fluidity, or formation of cracks and pores after casting. This article describes the processing of Ti-Al-Si alloys by arc melting, and shows the microstructure, phase composition, hardness, fracture toughness, and compression tests of these alloys. These results are compared with the same alloys prepared by powder metallurgy by the means of a combination of mechanical alloying and spark plasma sintering. Ti-Al-Si alloys processed by melting metallurgy are characterized by a very coarse structure with central porosity. The phase composition is formed by titanium aluminides and titanium silicides, which are full of cracks. Ti-Al-Si alloys processed by the powder metallurgy route have a relatively homogeneous fine-grained structure with higher hardness. However, these alloys are very brittle. On the other hand, the fracture toughness of arc-melted samples is immeasurable using Palmqvist's method because the crack is stopped by a large area of titanium aluminide matrix.

**Keywords:** titanium aluminides and silicides; casting; powder metallurgy

#### **1. Introduction**

Titanium intermetallics with other light elements (for example aluminium or silicon) are prospective high-temperature alloys, applicable especially for construction components working at high temperatures under static loads [1–5]. They are considered for high-temperature service because they offer a balance of oxidation resistance and mechanical properties at higher temperatures superior to conventional titanium alloys [6,7]. Ti-Al(Si) alloys are characterized by low density, good resistance against oxidation at 600–800 ◦C [8,9], good thermal stability, high specific strength at high temperatures [10,11], and a favourable ratio of mechanical properties to density [12]. Alloys based on α2-Ti3Al and γ-TiAl can be used as high-temperature construction materials for aviation applications [13–15], but low room-temperature ductility, as well as insufficient high-temperature strength and toughness, are problems which limit their broader use [16].

The production of Ti-Al alloys (and generally intermetallic compounds) is focused on conventional casting techniques, for example, on melting in an electric arc furnace under an argon atmosphere or electroshock remelting in an inert atmosphere [17]. Despite the enormous efforts made over the last 40 years, no castings have been produced to meet the reliability and cost requirements of the aviation industry [18]. The problems are caused especially by the high reactivity of molten titanium alloys with the ceramic crucible, and hence it is needed to use cold wall crucibles. However, cold wall crucibles enable low superheating of only about 60 ◦C. Slow cooling rates after centrifugal casting lead to the production of castings with a coarse-grained structure and large porosity. In order to remove the internal defects, the castings are processed by the HIP (Hot Isostatic Pressing) method. This technology is used by GE for engines for Boeing, but it is very expensive [19].

The melting of Ti-Al alloys is a multi-step process. The process involves the melting of an alloy, casting into a mould, extended isothermal annealing of ingot and controlled rapid crystallization of the melted metal [17]. Casting is often followed by the isostatic pressing of hot castings, which has a positive effect on homogeneity, but the production is more expensive [12].

Ti-Al alloys for turbine wheels of turbochargers are now manufactured by lost-wax casting [20]. The model of the required object is made of the wax and coated by the ceramics, therefore the massive mould is made. Then the wax is melted and poured out the mould. The metallic melt is filled into the mould and the casting solidifies. The process finishes with the removal of the ceramic mould after the solidification of the alloy. Another technology is induction skull melting (ISM). This process combines the advantages of induction melting and cold crucible melting. ISM is used for melting highly reactive material (for example TiAl) with high purity. According to this method, γ-TiAl alloys are melted in the vacuum induction furnace in a water-cooled crucible. The contamination from the crucible into the alloy is minimal. The big advantage is the continuous stirring, which guarantees the composition's uniformity [21,22]. This process is very energy-intensive and highly expensive because of the high cost of melting furnace and its non-ecological operation, high consumption of cooling water and inert gas and low achievable melt overheating [20]. The problems of cast Ti-Al alloys also include the reactions of titanium and aluminium with the atmosphere of the melting facility, the reaction of titanium with melting crucibles or ceramic moulds, or the evaporation of aluminium. The γ-TiAl alloys prepared by this process have a polycrystalline lamellar microstructure. If the sample is heated to a temperature above 1150 ◦C, the polycrystalline lamellar structure is replaced by a fine homogeneous duplex one with uniform grains. The lamellar structure is less ductile than the duplex one, but it has better fracture toughness, fatigue resistance, and creep strength at high temperatures [20]. However, it is also very difficult to forge the Ti-Al alloys due to inherent poor deformability even at temperatures above 1000 ◦C [23,24]. Chesnutt et al. wrote that it is possible to form α2-TiAl and γ-TiAl alloys (with difficulty). The α2-TiAl alloys are produced in ingots of 3200 kg [25]. Lapin wrote about extrusion and forging for producing compressor blades for engine testing. Compressor blades were produced by Thyssen, GfE, Leistritz and GKSS for Rolls-Royce using Ti-Al-Nb-(B,C) alloy, but the microstructure was very heterogeneous, and it can be expected that the segregation effects would be even larger for the preparation of bigger components. The boron addition enabled the refinement of the grain size of the as-cast titanium aluminides [19].

The production of Ti-Al-Si alloys by melting metallurgy is very difficult. The disadvantages include the high melting points of intermetallics (titanium silicide Ti5Si3 melts at the temperature of 2130 ◦C [12]) and high melt reactivity with the melting crucibles [8], so it is necessary to use Y2O3 and ZrO2 crucibles (which are more expensive than corundum or graphite crucibles). Other problems include the contamination of the melt from the reaction with the atmosphere in the furnace [12]. The Ti-Al-Si alloys have very poor fluidity, so the alloy has many casting defects after casting, for example, pores and cracks in the structure [26,27].

Hard and brittle sharp-edged titanium silicide (Ti5Si3) phases are formed by melting and casting of Ti-Al-Si alloy.  Hence, there is an effort to eliminate these big brittle phases of titanium silicides. In the previous study, the formation of in situ composites composed of elongated particles of Ti5Si3 phases in a tough matrix made of TiAl or Ti3Al aluminides has been investigated [29]. In the structure, the cracks perpendicular to the direction of solidification were found because of the different coefficient of thermal expansion of Ti5Si3 in different crystal directions. Therefore, tensile stresses and crack initiation occur [12]. These problems limit the possibility of processing of Ti-Al-Si materials by directional crystallisation [30].

Applicability of forming of Ti-Al-Si alloys is also limited due to the low fracture toughness and ductility of the material, which persists even at temperatures above 1000 ◦C [12]. However, the preparation of Ti-Al alloys with the addition of Si by ingot metallurgy is limited to eutectic and hypoeutectic alloys based on α2-Ti3Al, since hypereutectic alloys are extremely brittle because of the coarse primary silicides formed upon solidification [10].

In this work, Ti-Al-Si alloys were processed by melting metallurgy in an arc melting furnace. The results were compared with the same alloys prepared by powder metallurgy by the means of a combination of mechanical alloying and spark plasma sintering.

#### **2. Materials and Methods**

The Ti-Al-Si alloys (TiAl10Si20, TiAl10Si30, TiAl15Si15, TiAl20Si20 (wt. %)) were prepared by melting metallurgy in the arc melting furnace Bühler MAM-5 in the Institute of Physics of the Czech Academy of Sciences. These compositions have been chosen according to our previous research [26]. The pieces of individual elements (purity 99.95%) were four times re-melted under the argon atmosphere.

The reference samples were prepared by powder metallurgy. The optimal preparation route was chosen as a combination of mechanical alloying and spark plasma sintering (MA + SPS). The mixture of pure titanium (with a purity of 99.5% and a particle size of 44 μm, STREM CHEMICALS, Newburyport, MA, USA), aluminium (99.62%, 44 μm, STREM CHEMICALS, Newburyport, MA, USA) and silicon (99.5%, 44 μm, Alfa Aesar, Haverhill, MA, USA) powders were filled in a steel vial with steel milling balls (ball-to-powder weight ratio was 60:1) and milled under the argon atmosphere. For the mechanical alloying, a planetary ball mill (PM 100 CM, Retsch, Haan, Germany) was applied. Milling was performed for four hours with rotation velocity of 400 min−1. Mechanically alloyed powders were consolidated by spark plasma sintering (FCT Systeme GmbH, Rauenstein, Germany) under the pressure of 80 MPa and the temperature of 1100 ◦C for 15 min. The applied heating rate and cooling rates were 100 and 50 ◦C/min, respectively [31].

The phase composition was analyzed by the means of the X-ray diffraction analysis (XRD) using a X'Pert Pro (PANalytical, Almelo, Netherlands) X-ray diffractometer with CuKα radiation and a LynxEye XE detector (PANalytical, Almelo, The Netherlands). XRD patterns were evaluated qualitatively by X'Pert HighScore 3.0 software package (PANalytical, Almelo, Netherlands), exploiting PDF-2 2018 database. For the microstructure inspections, samples were ground by P80 to P4000 grinding papers (Hermes Schleifmittel GmbH, Hamburg, Germany) and polished by diamond paste with the particle size of 1–2 μm. The polished samples were etched by Kroll's reagent (it was prepared in our laboratory) (10 mL HF, 5 mL HNO3 and 50 mL H2O). The microstructure was investigated by the inverted optical microscope Olympus PME3 (Olympus, Prague, Czech Republic) and documented by the Carl Zeiss AxioCam ICc3 (Carl Zeiss, Jena, Germany) digital camera and AxioVision software package (version 4.8.2, Carl Zeiss, Jena, Germany). Porosity was evaluated by Lucia 4.8 image analyser (Laboratory Imaging, Prague, Czech Republic). TESCAN VEGA 3 LMU (TESCAN, Brno, Czech Republic) electron microscope with EDS analyser Oxford Instruments X-max 20 mm2 (Oxford Instruments, High Wycombe, UK) (SEM-EDS) was used for the deeper microstructure investigation and the identification of present phases on the micrographs.

Compressive strength tests were performed with the use of the universal testing machine LabTest 5.250SP1-VM (LaborTech, Opava, Czech Republic). Values of ultimate compressive strength were determined from the measured stress-strain curves. Vickers hardness with a load of 5 kg (HV 5), 100 grams (HV 0.1) and 50 grams (HV 0.05) was measured on the polished samples (10 measurements on each specimen). Fracture toughness was measured by Vickers indentation method with the load of 1 kg on microhardness tester Future-Tech FM-700 (Future-Tech, Kawasaki, Japan). Indentions were observed by the means of the above-mentioned Olympus PME3 microscope. Fracture toughness was calculated using Palmqvist's Equation (1):

$$\mathcal{K}c = 0.016 \cdot \left(\frac{E}{HV}\right)^{\frac{1}{2}} \cdot \left(\frac{F}{c^{\frac{3}{2}}}\right) \tag{1}$$

where E is modulus of elasticity (GPa), HV is Vickers hardness (GPa), F is load (N), and c is half of the crack length after indention (mm).

#### **3. Results**

The phase composition of as-cast Ti-Al-Si alloys is displayed in Figure 1. The TiAl10Si20 and TiAl15Si15 alloys consists of titanium silicide (Ti5Si3) and titanium aluminide (TiAl). Titanium silicide TiSi, Ti5Si4 and pure silicon were, as expected, found in the TiAl10Si30 alloy due to the higher amount of silicon in the alloy. The TiAl20Si20 alloy is composed of Ti5Si3 and Ti5Si4 silicides in TiAl3 aluminide matrix.

**Figure 1.** XRD patterns of Ti-Al-Si alloys processed by melting metallurgy.

Figure 2 shows the microstructure of Ti-Al-Si alloys prepared by arc melting. The as-cast alloys have coarse dendritic structures. As-cast TiAl10Si20 alloy is characterized by a very porous structure, as well as TiAl20Si20. The pores (black spots on the photos) are concentrated largely in the center of the sample and they are irregular in shape. Alloys TiAl10Si30 and TiAl15Si15 have much smaller pores. The particles of titanium silicides are very coarse in each alloy, and they have various local characters of morphology. The microstructure is characterized by primary silicides, silicides with fibrous and lamellar morphology and fine eutectic structure. The colonies of titanium silicides are differently oriented, depending on the local direction of heat transfer, which causes the elongation of silicides.

**Figure 2.** Microstructures of Ti-Al-Si alloys processed by melting metallurgy (light microscopy): (**a**) TiAl10Si20, (**b**) TiAl10Si30, (**c**) TiAl15Si15, (**d**) TiAl20Si20 (black areas are pores, lighter grey particles are titanium silicides, darker grey ones are titanium aluminides).

Figure 3 shows the porosity and average equivalent diameter of pores of Ti-Al-Si alloys prepared by melting metallurgy. The big differences between the values of porosity are given by preparation techniques. During the fast cooling, the gasses are closed in the middle of the sample and formed the middle porosity. At the edges of sample, the porosity is minimal.

**Figure 3.** Porosity vol. and average equivalent diameter of pores of Ti-Al-Si alloys processed by melting metallurgy. no subscript

Figure 4 presents the SEM micrographs of Ti-Al-Si alloys acquired in a backscattered electrons (BSE) regime. Except for the TiAl10Si30, all alloys have multiple cracks in the titanium silicides. Composition of each phase (TiAl, Ti5Si3) in TiAl10Si20 alloy corresponds well with the phase diagram. Only spectrum four contains a local lower amount of silicon, which was detected near the crack. TiAl10Si30 alloy has the finest particles of titanium silicides. At the same time, TiAl10Si30 alloy has the lowest amount of aluminide matrix (Figure 5). It is due to the highest amount of silicon in the basic chemical composition and the higher ratio between aluminium and silicon in the alloy. The TiAl10Si30 alloy processed by casting exhibited the presence of areas of aluminides, containing higher degree of substitution of aluminium by silicon (Figure 4b). In addition, TiAl10Si30 alloy are contaminated by iron, copper and oxygen from the preparation route. In TiAl15Si15 alloy, the chemical composition of TiAl and Ti5Si3 phase corresponds well with the phase diagram, i.e., the phases are minimally substituted by other elements. According to the EDS (Energy Dispersive Spectroscopy) analysis (Spectrum 11) and Ti-Al phase diagram, TiAl20Si20 should contain aluminide phase (TiAl3), which was confirmed by XRD diffraction.

The hardness of the Ti-Al-Si alloys (Figure 6) prepared by melting metallurgy varies between 416 and 549 HV 5. TiAl10Si30 alloy with the highest amount of silicon has a higher amount of hard silicides (Figure 5) and, hence, it reaches the highest values of hardness.

The microhardness of the Ti-Al-Si alloys (Figure 7) is in the range between 750 and 1066 HV 0.1 and TiAl10Si30 alloy achieves the highest hardness. It is caused by the fact that silicon in alloy is bonded in silicides (Ti5Si3, Ti5Si4 and TiSi), which are harder than the contained aluminides. The predicted hardness was calculated by Equation (2). The calculated hardness values of Ti-Al-Si alloys are compared with the measured ones in Figure 7. The values are comparable, only TiAl10Si30 alloy has lower values of calculated hardness than hardness, which was measured.

$$w\_{(\text{Ti}-\text{Si})} \cdot HV\_{(\text{Ti}-\text{Si})} + w\_{(\text{Ti}-\text{Al})} \cdot HV\_{(\text{Ti}-\text{Al})} = HV\_{(\text{Ti}-\text{Al}-\text{Si})} \tag{2}$$

where w(Ti-Si) is an area fraction of titanium silicides, HV(Ti-Si) is the Vickers hardness of titanium silicides, w(Ti-Al) is an area fraction of titanium aluminides, HV(Ti-Al) is Vickers hardness of titanium aluminides, and HV(Ti-Al-Si) is the hardness of the product.

**Figure 4.** Microstructure of Ti-Al-Si alloys processed by melting metallurgy (light microscopy): (**a**) TiAl10Si20, (**b**) TiAl10Si30, (**c**) TiAl15Si15, (**d**) TiAl20Si20 (lighter particles are titanium silicides, darker ones are titanium aluminides) (composition in atomic %).

**Figure 5.** Amount of silicides in Ti-Al-Si alloys processed by melting metallurgy.

**Figure 6.** Hardness of Ti-Al-Si processed by melting metallurgy.

Figure 8 shows the comparison of microhardness of silicide and aluminide phase. Silicide phase hardness varies between 1070 and 1175 HV 0.05, the hardness of the aluminide phase varies between 200 and 530 HV 0.05. The big variations in the hardness of the aluminide phase are given by the substitution of silicon in titanium aluminide and the formation of small particles of titanium silicides in Ti-Al matrix. Titanium aluminide phase TiAl3 present in TiAl20Si20 alloy has a higher hardness than TiAl phase. Values of the hardness of titanium silicides show only minor variations (Figure 8), so it is possible to say that silicides Ti5Si3 and Ti5Si4 have the same hardness, but titanium silicide TiSi present in TiAl10Si30 alloy has a slightly lower hardness.

**Figure 7.** Microhardness of the Ti-Al-Si alloys processed by melting metallurgy.

**Figure 8.** Comparison of microhardness of silicide and aluminide phase.

Intermetallics belong to the materials commonly referred as brittle. For these materials classes, the fracture toughness and compressive strength are the measures for potential use in the practice. Fracture toughness of Ti-Al-Si alloys processed by melting metallurgy was measured from the ten indentions into the silicide phase of each sample using Palmqvist's equation. Since titanium silicides are brittle, the cracks were formed during the indention. All of the alloys possess lower fracture toughness than the technical ceramics, such as alumina or silicon carbide having fracture toughness of around 4 MPa m1/2 [32]. Crack propagation was stopped by the present aluminide phase at the TiAl15Si15 alloy prepared by arc melting. On the other hand, the crack spreads through the sample in case of the same alloy prepared by powder metallurgy (Figure 9). Therefore, the fracture toughness of the titanium silicides was measured (Figure 10). It is possible to say that a coarser structure is conducive to an increase on the toughness of the material and the samples are less brittle. The signs of the plasticity of the as-cast alloys are also shown in the deformation curves in compression (Figure 11). The TiAl15Si15 alloy achieves the higher ultimate tensile strength in compression—1700 MPa. This value of UTS is comparable with the same alloy processed by the combination of mechanical alloying and spark plasma sintering [31].

**Figure 9.** Cracks after indentation for calculating the fracture toughness of TiAl15Si15 alloy processed by: (**a**) melting metallurgy, (**b**) powder metallurgy.

**Figure 10.** Fracture toughness of silicide phases in Ti-Al-Si alloys processed by melting metallurgy.

**Figure 11.** Compression tests of Ti-Al-Si alloys processed by melting metallurgy.

The reference samples of Ti-Al-Si alloys were prepared by powder metallurgy, by mechanical alloying and Spark Plasma Sintering (MA + SPS). SEM micrographs of the MA + SPS alloys are presented in Figure 12. The TiAl10Si20 alloy is characterized by a very homogeneous well-refined structure, and also the most fine-grained, with the Ti5Si3 strengthening phase in TiAl matrix. TiAl10Si30 is composed by Ti5Si3 and Ti5Si4 silicides in TiAl and TiAl2 matrix. The TiAl15Si15 alloy contains a small amount of Ti2Al phase. TiAl20Si20 alloy is composed of Ti5Si3, Ti5Si4, and TiSi2 silicides, while the higher ratio between aluminium and titanium unbound in silicides leads to the single-phase TiAl3 matrix. The sample also contains iron contamination originating from the milling steel vial.

Alloys processed by mechanical alloying and the spark plasma sintering method are characterized by lower porosity (Figure 13) than the alloys prepared by arc melting (Figure 3). The porosity of Ti-Al-Si alloys processed by powder metallurgy is lower than 1 vol. %. The pores are uniformly distributed (Figure 12).

The hardness of the Ti-Al-Si alloys (Figure 14) prepared by powder metallurgy varies between 800 and 1037 HV 5. These values of hardness are higher than values of Ti-Al-Si alloys processed by arc melting. The higher hardness corresponds to more homogeneous and fine-grained crackless microstructure. TiAl10Si30 and TiAl10Si20 reach the highest values of hardness.

**Figure 12.** Microstructure of Ti-Al-Si alloys processed by powder metallurgy using mechanical alloying and Spark Plasma Sintering (scanning electron microscopy): (**a**) TiAl10Si20, (**b**) TiAl10Si30, (**c**) TiAl15Si15, (**d**) TiAl20Si20.

The microhardness of the Ti-Al-Si alloys prepared by powder metallurgy (Figure 15) varies between 1246 and 1430 HV 0.1. The measured microhardness is higher than the microhardness of Ti-Al-Si materials processed by melting metallurgy.

**Figure 13.** Porosity vol. and average equivalent diameter of pores of Ti-Al-Si alloys processed by powder metallurgy.

**Figure 14.** Hardness of Ti-Al-Si processed by powder metallurgy.

The fracture toughness of Ti-Al-Si alloys prepared by powder metallurgy was measured by the indentation and was calculated by Palmqvist's Equation. All of the alloys have low values of fracture toughness, because the intermetallic phases present in these alloys are very brittle (Figure 16). The alloys processed by powder metallurgy have lower values of fracture toughness than the same alloys prepared by melting metallurgy.

**Figure 15.** Microhardness HV 0.1 of Ti-Al-Si processed by powder metallurgy (mechanical alloying and Spark Plasma Sintering).

**Figure 16.** Fracture toughness of silicide phases in Ti-Al-Si alloys processed by melting metallurgy.

The Ti-Al-Si alloys processed by powder metallurgy achieve higher ultimate tensile strength in comparison with the same alloys processed by melting metallurgy (Figure 17). The alloys have better values due to the more homogeneous and fine-grained structure.

**Figure 17.** Compression tests of Ti-Al-Si alloys processed by melting metallurgy.

#### **4. Discussion**

Table 1 shows the comparison of the phase composition of the Ti-Al-Si materials processed by melting and powder metallurgy. The phase composition of TiAl10Si20 alloy processed by both the arc melting and the MA + SPS is in good agreement with the equilibrium phase diagram of Ti-Al-Si ternary system [31,33]. The same results were obtained by reactive sintering followed by spark plasma sintering (RS + SPS) described in our previous article [26]. Another titanium silicide, TiSi, was found in the TiAl10Si30 alloy due to the higher amount of silicon in the alloy. Formation of TiSi corresponds to the equilibrium phase diagram [33] and this titanium silicide was also found in the TiAl10Si30 alloy prepared by RS + SPS [26]. As-cast TiAl15Si15 alloy has the same phase composition as in the equilibrium phase diagram, but the alloy prepared by MA + SPS contains a small amount of Ti2Al phase. The TiAl20Si20 alloy contains a combination of silicides Ti5Si3, Ti5Si4, and TiSi2 (by type of preparation) and a single-phase TiAl3 matrix caused by the higher ratio between aluminium and titanium unbound in silicides.

**Table 1.** Comparison of phase composition of Ti-Al-Si alloys (MA = mechanical alloying, RS = reactive sintering, SPS = Spark Plasma Sintering).

Mechanical alloying causes that the aluminium in titanium aluminide and silicon in titanium silicides are strongly substituted by silicon and aluminium, respectively. Authors described that Ti(Al,Si) supersaturated solid solutions are formed after at least 10 or 20 h of milling [34,35], but it is shown in this work that intermetallic phases are formed at the shorter time of mechanical alloying (after 4 h).

The Ti-Al-Si alloys prepared by melting metallurgy (Figure 2) have a very coarse structure with pores, which are concentrated in the centre of the sample. The coarse particles of titanium silicides are oriented depending on the local direction of heat transfer, which causes the elongation of silicides. These silicides are cracked due to the high cooling rate after melting [36]. The central porosity is given by the processing technology. Gases are trapped in the sample as a result of evaporation of the elements (mostly aluminium) during the exothermic formation of intermetallics and relatively rapid cooling. The highest cooling rate is on the surface of the samples, so the samples solidified from the surface to the core. The gases are trapped in the middle sample and they cannot escape from the sample because of the fast solidification of the surface. Therefore, the porosity of the samples is concentrated in the centre of the samples.

Powder metallurgy (MA + SPS) results in a very fine microstructure of the investigated alloys (Figure 12). Authors described that mechanical alloying is a very effective process in obtaining of nanosized grains [37,38]. Refining of Ti5Si3 crystallites is probably due to a strong deformation and consequent recrystallization. The distribution of titanium silicides is much more homogeneous because the brittle powder is crushed to very fine particles during the first step of mechanical alloying. The elements in the obtained compounds are substituted mutually, exceeding the equilibrium solubility limits [39]. Figure 8 compares the hardness of silicide and aluminide phase in as-cast Ti-Al-Si alloys. Titanium aluminide TiAl3 is harder than TiAl. This confirms the fact than TiAl3 has covalent bond and the TiAl phase has a metallic bond, which causes the higher plasticity [40,41]. The hardness of titanium aluminide TiAl3 is referred to as 450 HV [42], TiAl phase around 300 and 350 HV 5 [43]. The hardness of Ti5Si3 is referred to as 970 HV [44]. Calculated and measured results in Ti-Al-Si alloys are a little bit higher than the values reported in the literature. Hardness values of titanium silicides have a very small scattering of values, so it is possible to say that they have isotropic mechanical properties [36].

Powder metallurgy (MA + SPS) resulted in an enormous porosity reduction, which decreased from 26 vol. % in the case of as-cast TiAl10Si20 to under 1 vol. % by the same alloy (Figure 3, Figure 13). The other alloys recorded a similar decline. Also, the average equivalent diameter of pores decreases from the tens of micrometres to micrometre units. The applied the load, the sintering temperature and duration are the most important parameters, which influence the microstructure, porosity and pore size of the resulting compact sample. The better compaction of the powder and the smaller porosity and pore size are obtained from the higher temperature and applied load, but it may cause the grain coarsening [45]. In previous works, the optimal conditions of spark plasma sintering were studied so that the porosity and pore size were as low as possible [46].

The hardness, microhardness, and ultimate tensile strength (UTS) of the Ti-Al-Si alloys prepared by powder metallurgy (MA + SPS) are higher than the same alloys prepared by arc melting. The higher hardness and UTS corresponds to a more homogeneous and crackless microstructure, which is described above. The force during the mechanical alloying deforms the powder particles plastically, which leads to work hardening and fracture. Hardening, in combination with a fine-grained structure, increases the material hardness [45], but lowers the plasticity.

#### **5. Conclusions**

In this work, Ti-Al-Si were prepared by arc melting and powder metallurgy. Both technologies were compared. Ti-Al-Si alloys processed by arc melting are characterized by a porous structure with very coarse titanium silicide particles in a titanium aluminide matrix.  Reference samples processed by mechanical alloying and spark plasma sintering are characterized by a more homogeneous and finer structure with low porosity. These properties guarantee excellent mechanical properties. Vickers hardness of MA+SPS samples reaches double values, but Ti-Al-Si alloys processed by powder metallurgy are very brittle. The fracture toughness is worse than for technical ceramics. On the other hand, arc-melted samples have immeasurable fracture toughness using Palmqvist's method, the Vickers indentions do not cause the cracks, and crack propagation is stopped by the aluminide phase. For this reason, it is reasonable to say that TiAl15Si15 alloy prepared by melting metallurgy is the most promising alloy, as the formation of a coarse-grained structure improves the fracture toughness, the alloy is more ductile,

and comparable ultimate tensile strength (as MA+SPS) predetermines this alloy for use in applications where high hardness is not the most important. However, it is necessary to improve the melting process (to optimize the conditions of melting and casting), or to use the hot isostatic pressing of the castings, in order to decrease the porosity of the material.

**Author Contributions:** A.K. was responsible for the theoretical background, writing the article, preparing samples and data analysis. P.N. was responsible for final correction. J.K. prepared samples by melting metallurgy. F.P. provided sintering process.

**Funding:** This research was financially supported by specific university research MSMT No 21-SVV/2019.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Influence of Heat Treatment on Microstructure and Properties of NiTi46 Alloy Consolidated by Spark Plasma Sintering**

**Pavel Salvetr 1,\*, Jaromír Dlouhý 1, Andrea Školáková 2, Filip Pr ˚uša 2, Pavel Novák 2, Miroslav Karlík 3,4 and Petr Haušild 3**

Received: 12 November 2019; Accepted: 3 December 2019; Published: 6 December 2019

**Abstract:** Ni-Ti alloys are considered to be very important shape memory alloys with a wide application area including, e.g., biomaterials, actuators, couplings, and components in automotive, aerospace, and robotics industries. In this study, the NiTi46 (wt.%) alloy was prepared by a combination of self-propagating high-temperature synthesis, milling, and spark plasma sintering consolidation at three various temperatures. The compacted samples were subsequently heat-treated at temperatures between 400 ◦C and 900 ◦C with the following quenching in water or slow cooling in a closed furnace. The influence of the consolidation temperature and regime of heat treatment on the microstructure, mechanical properties, and temperatures of phase transformation was evaluated. The results demonstrate the brittle behaviour of the samples directly after spark plasma sintering at all temperatures by the compressive test and no transformation temperatures at differential scanning calorimetry curves. The biggest improvement of mechanical properties, which was mainly a ductility enhancement, was achieved by heat treatment at 700 ◦C. Slow cooling has to be recommended in order to obtain the shape memory properties.

**Keywords:** Ni-Ti alloy; self-propagating high-temperature synthesis; spark plasma sintering; aging; compressive test; hardness; shape memory

#### **1. Introduction**

The Ni-Ti alloys named NiTinol (derived from nickel, titanium, and laboratory of discovery-Naval Ordnance Laboratory) are well-known shape memory alloys with good mechanical properties and high corrosion resistance, which enables usage as implants, medical devices, and other applications as biomaterials [1]. The shape memory effects occur due to the reversible solid-state transformation between the high-temperature and low-temperature phases. The high-temperature phase is called austenite with a high-symmetry structure-ordered body-centered cubic phase B2 (CsCl). The low-temperature and low-symmetry phase is called martensite with monoclinic B19 lattice. Reversible strains at about 8% of the initial length are enabled due to the reversible phase transformation as well [1]. To achieve the desired shape memory effects-transformation temperatures, it is necessary to be careful about a chemical composition. The transformation temperatures are very sensitive for the nickel-titanium ratio and an increase in the content of nickel by 0.1 at a percentage causes a change of the transformation temperature AF (austenite finish) up to 10 ◦C [2,3]. This fact makes the production of the Ni-Ti alloys more difficult because titanium and, subsequently, NiTi alloys have high affinity to oxygen from atmosphere and carbon from the melting crucible during vacuum induction melting (VIM) [4,5].

Spark plasma sintering (SPS) is a modern consolidation process, which is suitable for various materials such as ceramics and metals including many intermetallic systems (e.g., Ni-Ti and Fe-Al alloys [6–9]). The high heating rate and shortness of the whole process enable the use of SPS for the consolidation of nanocrystalline materials as well [10–12]. The process is based on the sintering of powder under the simultaneous influence of high electric current (direct or pulsed) and uniaxial pressure. The Joule heat is generated by passing the current through the graphite punch and die and between the powder's particles. The high heating rate was described as the route to reduce the amount of undesirable secondary phase such as the Ti2Ni by the self-propagating high-temperature synthesis (SHS) [13]. Thus, the SPS process was examined as a heating source for initiating the SHS reaction between nickel and titanium elemental powders in the previous paper. However, the SPS process seems to be inapplicable for the initiation of the SHS reaction because the strongest increase of the temperature occurs on the surface of the particles. The formed intermetallic layers act further as diffusion barriers and separate unreacted nickel from titanium [14]. Therefore, the pre-alloyed Ni-Ti powder after mechanical alloying is usually sintered by SPS. This process can produce the highly dense NiTi materials [7] or the porous structure depending on the addition of the space holder (e.g., NH4HCO3) [15,16].

Heat treatment of the Ni-Ti alloys has a crucial effect on the properties of the samples. The parameters of heat treatment influence the microstructure, internal stresses, precipitation, shape memory, and mechanical properties [17]. The Ni-Ti alloys undergo homogenizing treatment at the temperature of about 800–1050 ◦C up to several hours of duration, which is followed by water quenching to get a homogeneous microstructure without the Ni-rich precipitates [17–22]. The second step of heat treatment represents the aging treatment, which usually occurs between 300 ◦C and 800 ◦C [23,24]. The grade of aging (density and size of the Ni4Ti3 precipitates) depends on the temperature and time. The precipitation process starts with the metastable Ni4Ti3 phase, which is transformed into the metastable Ni3Ti2 phase and the stable Ni3Ti [25]. The precipitation process is accompanied by the hardness changes during aging due to formation, coarsening, and decomposition of the Ni4Ti3 and Ni3Ti2 phases. The addition of Al enhances the microstructural and hardness stability of Ni-rich Ni-Ti alloys until 500–600 ◦C [17,26]. The results of phase formation, stability, and transformation Ni4Ti3 → Ni3Ti2 → Ni3Ti during aging are presented by a time-temperature-transformation (T-T-T) diagram [17,25,27].

In this work, the NiTi46 (wt.%) alloy was processed by a combination of SHS, milling in a vibratory mill, and SPS consolidation at three temperatures to get fully dense materials. The prepared samples underwent the heat treatment in temperatures ranging from 400 to 900 ◦C. The characterization of samples was focused on phase composition, observing the transformation temperatures, and changing the mechanical properties depending on the heat treatment regime.

#### **2. Materials and Methods**

The metallic powders with the following particle sizes and purities were used as starting material for the NiTi46 (wt.%) alloy: nickel (particle size < 150 μm, 99.9 wt.% purity, Sigma-Aldrich, St. Louis, MO, USA) and titanium (particle size < 44 μm, 99.5 wt.% purity, STREM CHEMICALS, Newburyport, MA, USA). The powders were mixed manually corresponding to the chemical composition of the NiTi46 powder mixture, which was uniaxially compressed at room temperature to cylindrical green bodies of 12 mm in diameter at a pressure of 450 MPa for 5 min using LabTest 5.250SP1-VM universal loading machine (Labortech, Opava, Czech Republic). The SHS reaction of the pressed powder mixture was carried out in the fused silica ampoules evacuated to 10−2 Pa and sealed, which were placed in the preheated region to 1100 ◦C electric resistance furnace. The duration of the reaction was 20 min with the

following cooling in air. The properties of the samples prepared this way were described in a previous paper [28]. The microstructure is composed of the two phases (NiTi austenite – cubic, Ti2Ni – cubic), hardness 276 HV10, area fraction of the Ti2Ni phase 11.7%, and transformation temperatures: AS = 56 ◦C, AF = 86 ◦C, and MS = 21 ◦C. The SHS product was milled in a vibratory cylinder mill VM4 (OPS Pˇrerov, Pˇrerov, Czech Republic) in atmosphere with a duration of 7 min and the powder fraction with a particle size <355 μm was selected by sieving using Fritsch Analysette 3 device (FRITSCH GmbH, Germany). This pre-alloyed NiTi46 powder was consolidated by using the SPS method (FCT Systeme HP D 10, Frankenblicke, Germany) at three various temperatures (900, 1000, and 1100 ◦C) under the pressure of 50 MPa with a holding time of 10 min. The high heating rate was chosen as 300 ◦C/min at the beginning and the last 100 ◦C with the heating rate of 100 ◦C/min (for example, sintering temperature of 1000 ◦C: applied heating rate of 300 ◦C/min up to 900 ◦C, between the temperatures 900 and 1000 ◦C and the applied heating rate of 100 ◦C/min). The conditions of SPS consolidation as the temperature regime, compaction force, height reduction, and current flow are shown in Figure 1.

**Figure 1.** SPS parameters during consolidation at the temperature of 900 ◦C.

The heat treatment in the temperature range from 400 ◦C to 900 ◦C was applied to the SPS-ed samples. The duration of heat treatment was 60 min. Two variants of cooling were used including a high cooling rate with quenching in water and a slow cooling rate, which was provided by cooling in the closed furnace (average cooling rate approximately 2.5 ◦C/min between 700 ◦C and 300 ◦C).

The metallographic samples were prepared by grinding and polishing and the microstructure was revealed by etching in Kroll's reagent (5 mL HNO3, 10 mL HF, and 85 mL H2O). The microstructure was observed using scanning electron microscopes (SEM) equipped with the EDS (Energy Dispersive Spectroscopy) analyzers for identification of the chemical composition of the individual phases: VEGA 3 LMU (TESCAN, Brno, Czech Republic) equipped with the OXFORD Instruments X-max 20 mm2 SDD EDS analyzer (Oxford Instruments, HighWycombe, UK) and JEOL IT 500 HR 500 (JEOL, Tokyo, Japan). The phase composition was analyzed by the X-ray diffraction analysis (XRD) using a X'Pert Pro (PANalytical, Almelo, The Netherlands) X-ray diffractometer with CuKα radiation and a LynxEye XE detector (PANalytical, Almelo, The Netherlands). The mechanical properties of the samples were evaluated by measuring Vickers hardness with a load of 10 kg and compression tests (LabTest 5.250SP1-VM universal loading machine Labortech, Opava, Czech Republic) with a strain rate of

0.3 s−1 on samples measuring 3.3 mm × 3.3 mm × 5 mm. Compression tests were conducted in both the direction (longitudinal and perpendicular) to the direction of SPS. A longitudinal direction is parallel to compressive force during the SPS process. Differential scanning calorimetry (DSC) analysis of the prepared alloys was performed using Setaram DSC 131 (Setaram, Caluire, France) to determine the transformation temperatures in products. Measurements for determining temperatures austenite start (AS) and austenite finish (AF) were carried out in the temperature range of −20 ◦C to 200 ◦C at a heating rate of 10 ◦C/min and cooling from 200 ◦C to −5 ◦C for detecting the martensite start (MS) and martensite finish (MF) temperatures.

The samples compacted by SPS at 900 ◦C and with following heat treatments at 600 and 700 ◦C for 1 h and slow cooling in the closed furnace were also investigated using transmission electron microscopy (JEOL JEM 2200FS, JEOL, Tokyo, Japan, accelerating voltage of 200 kV). Standard 3 mm samples prepared by a slow-speed diamond blade cutting were mechanically dimpled and ion polished in a Gatan PIPS 691 device (Pleasanton, CA, USA).

#### **3. Results and Discussion**

#### *3.1. Microstructure, Phase Composition, and Phase Transformation*

First, the influence of used sintering temperature on the quality of sintering individual particles was investigated by a porosity measurement. Since it is visible in Figure 2 and acquired by a light microscope, there are differences between samples compacted at various temperatures. The direction of observation plays an important role. The non-deformed grains were observed on the perpendicular cut (perpendicular to the direction of compression) whereas the elongated shape of grains after loading during SPS remained in the microstructure on the longitudinal cut (parallel to SPS compression). The highest values of porosity were determined at the samples sintered at 900 ◦C. The porosity was high at the perpendicular level and also at the longitudinal cut. The SPS temperature of 1000 ◦C was sufficient to the reduction of porosity in comparison to the temperature of 900 ◦C. Mainly in the longitudinal direction, the value of porosity decreased rapidly to a similar value, which was measured after sintering at 1100 ◦C. The values of porosity are compared in Table 1 and, based on the porosity measurement, it is clear that the higher temperature of the SPS process leads to superior sintering of individual particles.

(**a**) (**b**)

**Figure 2.** *Cont.*

**Figure 2.** Microstructure and porosity of the NiTi46 alloy SPS consolidated at various temperatures: (**a**) 900 ◦C-perpendicular, (**b**) 900 ◦C-longitudinal, (**c**) 1000 ◦C-perpendicular, (**d**) 1000 ◦C-longitudinal, (**e**) 1100 ◦C-perpendicular, and (**f**) 1100 ◦C-longitudinal direction.

**Table 1.** Influence of SPS temperature on porosity of the sample and area fraction of the Ti2Ni phase in the microstructure.

The SPS temperature influences the phase composition and also mechanical properties. The effect of sintering temperature was investigated in the previous paper [29]. The area fraction of the undesirable Ti2Ni and Ni3Ti phases increased with an increasing sintering temperature. The high amount of the Ti2Ni phase was formed along the boundaries of the sintered particles. It is necessary to point out that, in the previous paper, pulse current flow through the sample (another SPS device) was used while, in this paper, the regime of direct current flow through the sample was applied and it is the reason for different results. In this case, the lower amounts of the Ti2Ni phase were measured generally and a growing trend with SPS temperature was not observed. The area fraction of the Ti2Ni phase increased slightly by SPS consolidation at 900 ◦C and 1000 ◦C, but, after SPS sintering at 1100 ◦C, the amount of the Ti2Ni phase was reduced to approximately 12%, which means a comparable value to result in samples prepared by the SHS method [29] (see Table 1).

At all SPS temperatures, the SEM observation was performed. Improving the fusion of grain boundaries was confirmed with increasing sintering temperature. The NiTi phase matrix with the fine Ni-rich precipitates in all samples was commonly found within the Ti2Ni and Ni3Ti phases. The microstructures after SPS are shown in Figure 3. In Table 2, there are summarized chemical compositions of individual areas observed by SEM. A good agreement in chemical compositions of the NiTi and Ti2Ni phases to the binary Ni-Ti phase diagram was found out. The chemical composition of the area labelled 3 is close to the Ni3Ti phase. Chemical composition of areas 7 and 10 is approaching the chemical composition of the Ni4Ti3 phase.

A more detailed observation of the microstructure was performed using a transmission electron microscope (TEM). The main goal of this experiment is based on investigating the fine needle-like particles in the NiTi matrix. Figure 4 shows TEM micrographs of sample SPS-ed at the temperature of 900 ◦C. The NiTi and Ti2Ni phase were observed commonly with the Ni4Ti3 phase (determined by electron diffraction) in the NiTi matrix.

**Figure 3.** *Cont.*

**Figure 3.** SEM micrographs of SPS-ed samples at various temperatures: (**a**) 900 ◦C, (**b**) 900 ◦C-detail of Ni-rich precipitates, (**c**) 1000 ◦C, (**d**) 1000 ◦C-detail of Ni-rich precipitates, (**e**) 1100 ◦C, and (**f**) 1100 ◦C-detail of Ni-rich precipitates and formed the Ni3Ti phase.

**Table 2.** Chemical composition of individual areas measured by EDS analysis.

**Figure 4.** *Cont.*

**Figure 4.** Ni4Ti3 phase in the matrix: (**a**) bright-field micrograph, g = 10-1, close to the [4,1,4] zone axis, (**b**) corresponding dark-field micrograph using Ni4Ti3 spot marked by a circle in the diffraction pattern in the inset, and (**c**) bright filed micrograph showing dark Ni4Ti3 particles in the matrix adjacent to a Ti2Ni particle.

The effect of heat treatment on the microstructure and the phase composition was investigated at samples SPS-ed at the temperature of 900 ◦C. The phase compositions of SPS-ed samples were verified by XRD analysis. The diffraction lines were very similar through all SPS consolidation temperatures. The phase composition of SPS-ed samples consists of the NiTi phases (Cubic, Pm-3m), Ti2Ni phase (Cubic, Fd-3 m), Ni3Ti phase (Hexagonal, P63/mmc), and the Ni4Ti3 phase (Rhomboedral, R-3). XRD patterns are displayed. After heat treatment in the temperature range of 600–900 ◦C with a 1-h duration, there were no observed changes in the phase compositions of the samples (see Figure 5). This fact can seem to be strange in comparison with other studies. However, it is necessary to consider the initial state in individual studies (mostly after homogenization annealing, e.g., [17,27]). The initial state of these samples is after SPS consolidation at 900, 1000, and 1100 ◦C. The very high heating rate (approximately 300 ◦C/min) and a very high cooling rate (as shown in Figure 1) was applied during the SPS process, whereas the initial state of the Ni-Ti alloys is after homogenization or solution annealing for tens of minutes or several hours at temperatures at around 1000 ◦C [17,27,30]. Moreover, the transformations in the microstructure and phase composition occur in a bigger extent after heat treatment with longer duration since it is clear in this study [17]. The crystallite sizes of the NiTi phase after spark plasma sintering, which was determined by the means of the Sherrer's formula, range from 20 to 47 nm. The crystallite sizes increased after heat treatment and the highest values of the crystallite sizes were determined at samples heat-treated at a temperature of 700 ◦C. The significant increase of the crystallite sizes of the NiTi phase is related likely to the recrystallization process, grain growth, or the order-disorder transformation around the temperature of 600–700 ◦C reported in References [31,32]. The values of the crystallite sizes of the NiTi phase depending on the regime of the heat treatment are stated in Table 3.

(**a**)

(**b**)

**Figure 5.** XRD patterns of the NiTi46 alloys: (**a**) spark plasma sintered at 900, 1000, and 1100 ◦C, (**b**) spark plasma sintered at 900 ◦C, heat-treated at 600–900 ◦C and slowly cooled, (**c**) spark plasma sintered at 900 ◦C, heat-treated at 600–900 ◦C, and cooled in water.

**Table 3.** Crystallite sizes of the NiTi phase depending on the regime of heat treatment.

From the point of view of the microstructure, the particles of the Ti2Ni, Ni3Ti phases and Ni4Ti3 needles in the NiTi matrix were observed in Figure 6. A higher amount of the Ni3Ti phase was formed in the sample heat-treated at 700 ◦C and cooled in water than in the samples with slow cooling in the closed furnace after heat treatment. The area fraction of the Ti2Ni phase after heat treatment was almost invariable with values between 13% to 16%. Detailed observation of the Ni4Ti3 and other phases after heat treatment at 600 ◦C and 700 ◦C was carried out by TEM again (see Figure 7).

**Figure 6.** SEM microstructures of samples SPS-ed at 900 ◦C and heat-treated: (**a**,**c**) at 700 ◦C followed with slow cooling in the closed furnace, (**b**) at 700 ◦C and cooled in water, and (**d**) at 600 ◦C followed with slow cooling in the closed furnace.

**Figure 7.** TEM observation of samples heat-treated at 600 ◦C (**a**,**b**) and at 700 ◦C (**c**,**d**) and slow cooled: (**a**) Ni4Ti3 precipitates with a typical lenticular cross-section, (**b**) corresponding diffraction pattern in the zone axis [2,-1-1,3] of the precipitate, which is coincident with the zone axis [1,0,2] of the B2 NiTi matrix, (**c**) Ni4Ti3 particles in the NiTi matrix adjacent to very fine grains of the Ti2Ni phase, (**d**) Ni4Ti3 particle adjacent to a coarser grain of the Ti2Ni phase. The inset shows a diffraction pattern of the Ti2Ni phase in the [1,1,2] zone axis.

The phase transformation in the NiTi phase between austenite and the martensite structure was studied using differential scanning calorimetry. The straight lines were obtained for the samples as-SPS sintered at all temperatures and any phase transformation occurs in these samples. The change in the phase transformation behaviour was brought by heat treatment of the samples. When the samples were heat-treated, the temperature of heat treatment and way of cooling are important. Fast cooling in water did not cause the recovery of the phase transformation. Therefore, the microstructure observation and mechanical property investigation are focused on the samples cooled slowly in the closed furnace. The peaks on DSC curves were formed only at samples after heat treatment at temperatures of 600–700 ◦C, which were slow cooled in the closed furnace. The heating and cooling DSC curves of samples SPS-ed at 900 ◦C are shown in Figure 8.

**Figure 8.** DSC heating and cooling curves: (**a**) Heating curves of samples SPS-ed at 900 ◦C and heat-treated between 500 and 700 ◦C. (**b**) Cooling curves of samples SPS-ed at 900 ◦C and heat-treated between 500 and 700 ◦C. (**c**) Heating curves of samples SPS-ed at 1000 and 1100 ◦C and SPS-ed at 1000 and 1100 ◦C with heat treatment at 600 ◦C. (**d**) Cooling curves of samples SPS-ed at 1000 and 1100 ◦C SPS-ed at 1000 and 1100 ◦C with heat treatment at 600 ◦C.

#### *3.2. Mechanical Properties*

The increase of hardness with the increasing temperature of SPS is similar to the previous results [29]. However, an increase of hardness was assigned to the rising value of the Ti2Ni phase. It is contrary to the current study and to the decrease of the Ti2Ni phase amount after SPS at 1100 ◦C. There are two possible reasons. Firstly, the hardness increases with the quality of powder sintering at higher SPS temperature, which is visible in Figure 3. Secondly, the precipitation process of Ni-rich phases occurs during the SPS process while heat treatment at the temperatures of 900–1000 ◦C lead to improved hardness [17]. Values of hardness as SPS-ed samples are stated in Table 4.

**Table 4.** Summary of hardness and compressive stress-strain test of samples after spark plasma sintering at various temperatures. Longitudinal direction is parallel to the direction of compressive force by the SPS process.

The stress-strain behavior was analyzed in tandem with hardness and the same influence (increasing values of UCS) with an increasing temperature of the SPS process was observed. As visible in Figure 9, the samples after SPS reach high values of ultimate compressive strength (UCS) 1900–2300 MPa, but there are not the areas of plastic deformation on stress-strain curves and the samples fail in a brittle manner at a maximum load perpendicularly as well as in the longitudinal direction. The values of elongation at maximum force (Agt) were between 7.4–9.4% for all SPS temperatures and both directions. In the case of the compression test, the lower porosity of samples SPS-ed at higher temperatures is likely connected with increasing values of UCS.

After heat treatment with cooling in the closed furnace, the decrease of hardness was found and the lowest value was measured after heat treatment at the temperature of 700 ◦C. Currently, with the decrease of hardness, increased ductility and plasticity of samples was measured by a compressive test (Table 5). The evolution of hardness depending on the temperature of heat treatment is shown in Figure 9 for samples SPS sintered at 900 ◦C. In case of samples SPS sintered, a 1000 ◦C and 1100 ◦C, hardness after heat treatment at 600 ◦C dropped to values of 504 HV 10 and 509 HV 10. By the compressive test after heat treatment, the difference between annealing temperatures of 600 ◦C and 700 ◦C was observed. In case of annealing temperature of 600 ◦C, the similar values of Agt were measured at longitudinal and perpendicular directions between 9.5% and 10.5%, whereas, at an annealing temperature of 700 ◦C, the increase of Agt was found and the difference between longitudinal (16.3–19.1%) and perpendicular (9.3–13.6%) direction grew up. The UCS of the samples SPS-ed at 1000 ◦C and 1100 ◦C increased after heat treatment about 100 MPa in comparison with the state after SPS. Generally, the values of UCS and Agt increased with an increasing temperature of the SPS process.

**Figure 9.** *Cont.*

**Figure 9.** Compressive stress-strain curves and evolution of hardness: (**a**) Stress-strain curves of samples SPS-ed at 900 ◦C and heat-treated-perpendicular direction. (**b**) Stress-strain curves of samples SPS-ed at 900 ◦C and heat-treated-longitudinal direction. Stress-strain curves of samples SPS-ed at 1000 ◦C, 1100 ◦C, and heat-treated-perpendicular (**c**) and longitudinal (**d**) direction. (**e**) Evolution of hardness during heat treatment of sample SPS-ed at 900 ◦C.

**Table 5.** Summary of mechanical properties of heat-treated samples consolidated by the SPS process.

The explanation of the rapid decrease of hardness could be caused due to removing the deformation strengthening from the milling of an SHS product. This type of decrease has to have the same scale for all samples. The dependence of hardness on the temperature of heat treatment is undeniable. Thus, the different processes and changes in microstructure must occur during heat treatment at various temperatures. However, these changes were not observed in phase composition and microstructure due to the short time of annealing during heat treatment. In this study [17], the longer time heat treatment was applied and the following changes in the microstructure occurred. The high UCS, hardness, and low ductility of samples were attributed to the precipitates of Ni3Ti4 and Ni3Ti2 phases, while, when the microstructure contains the Ni3Ti phase or combination of the Ni3Ti and Ni3Ti2 phases (corresponding to heat treatment between 600 and 800 ◦C while a longer time must be applied at a temperature of 600 ◦C), high mechanical properties and good ductility were obtained. Both studies [17,24] show similar evolution of the hardness values during heat treatment when compared to our results.

#### **4. Conclusions**

The fabrication process composed of Self-propagating High-temperature Synthesis (SHS) and Spark Plasma Sintering (SPS) was chosen to obtain a completely dense material. The highest temperature of SPS (1100 ◦C) led to the highest values of ultimate compressive strength without the formation of an excessive number of undesirable phases. The following heat treatment is necessary to obtain a material with a good combination of strength and ductility. The heat treatment leads to the disappearance of the deformation strengthening coming from milling in the vibration mill and to recover the phase transformation between the austenite and martensite structure of the NiTi phase, which was detected by differential scanning calorimetry (DSC). The heat treatment near the temperatures of 600–700 ◦C with slow cooling is recommended to obtain good ductility, strength, and probable shape memory properties.

**Author Contributions:** P.S. and P.N. designed the experiment and evaluated the phase composition. P.S. and A.Š. provided sample preparation by the SHS reaction, measurement of mechanical properties, and DSC analysis. F.P. consolidated samples by the SPS method. P.S. and J.D. analyzed the microstructure by LM and SEM. M.K. and P.H. observed the microstructure by SEM and TEM. P.S. wrote the paper. P.N. and M.K. reviewed and edited the paper.

**Funding:** European Regional Development Fund (projects Pre-Application Research of Functionally Graduated Materials by Additive Technologies, No. CZ.02.1.01/0.0/0.0/17\_048/0007350 and Nanomaterials Centre for Advanced Applications, No. CZ.02.1.01/0.0/0.0/15\_003/0000485) and specific university research (MSMT No. 21-SVV/2019) supported the research.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Application of a Dy3Co0.6Cu0.4H***x* **Addition for Controlling the Microstructure and Magnetic Properties of Sintered Nd-Fe-B Magnets**

**Katerina Skotnicova 1,\*, Pavel A. Prokofev 2,3, Natalia B. Kolchugina 3, Gennady S. Burkhanov 3, Alexander A. Lukin 2, Yurii S. Koshkid'ko 3,4, Tomas Cegan 1, Henryk Drulis 4, Tatyana Romanova 4 and Nikolay A. Dormidontov 3**

Received: 14 November 2019; Accepted: 15 December 2019; Published: 17 December 2019

**Abstract:** The focus of new technologies on the formation of inhomogeneous distributions of heavy rare-earth metals (REMs) in hard magnetic Nd–Fe–B materials is of scientific importance to increase their functional properties, along with preserving existing sources of heavy REMs. This paper focused on the coercivity enhancement of Nd2Fe14B-based magnets by optimizing the microstructure, which includes the processes of grain boundary structuring via the application of a Dy3Co0.6Cu0.4H*x* alloy added to the initial Nd–Fe–B-based powder mixtures in the course of their mechanical activation. We have studied the role of alloying elements in the formation of phase composition, microstructure, the fine structure of grains, and the hysteretic properties of hard magnetic Nd(R)2Fe14B-based materials. It was shown that the Dy introduction via the two-component blending process (the hydrogenated Dy3Co0.6Cu0.4 compound is added to a powder mixture) resulted in the formation of the core-shell structure of 2–14–1 phase grains. The efficient improvement of the coercivity of Nd(RE)–Fe–B magnets, with a slight sacrifice of remanence, was demonstrated.

**Keywords:** grain boundary diffusion; Nd–Fe–B magnets; hydrogenation; microstructure; magnetic properties

#### **1. Introduction**

Researchers have made many attempts to reduce the heavy rare-earth (RE) consumption Nd–Fe–B sintered magnets with high-coercivity. Some progress has been achieved using Dy and/or Tb in various forms to realize approaches named grain boundary diffusion (GBD) [1–3] and grain boundary structuring (GBS) [4–8]. The application of binary mixtures allows one to improve the structure of the boundary phases and grain boundaries of the main magnetic phase and to realize the diffusion of a required component of the alloy directly through the boundaries. It has been demonstrated that by controlling the process time and temperature of GBD processes, the coercivity of the magnet can be greatly enhanced, without sacrificing the remanence.

It was shown in our previous studies that hydrogenated Tb and Dy additions allowed us to enhance the coercivity with a slight decrease in the remanence [9] and increase the stability of the magnet properties during annealing at the low-temperature [10], respectively.

The grain boundary restructuring, with rare-earth-rich low-melting compounds added to low-alloyed Nd–Fe–B-based compositions in the course of technological processing, was realized when using (Pr,Nd)6Fe13Cu [4], Dy32.5Fe62Cu5.5 [5], Dy69Ni31 [6], Dy88Mn12 (wt.%) [11], Pr34.4Co65.6 (wt.%) [12], and Dy82.3Co17.7 (wt.%) [13], which is a low-melting eutectic composition. It was shown that the intrinsic coercivity evidently increased when using Dy82.3Co17.7 and the maximum intrinsic coercivity was achieved when its content was 2 wt.%. At the same time, the remanence and maximum-energy product decreased slightly as the Dy82.3Co17.7 content increased. By adding a small amount of Dy82.3Co17.7, the coercivity improved greatly, and the irreversible loss decreased sharply. The increase in the Curie temperature of the magnets suggests that Co atoms have been incorporated into the 2:14:1 main phase. A well-developed a core–shell structure is formed in these magnets.

The experiments with REM-M-H compounds (rare earth metal-transition metal(s)-hydrogen), which are added at the stage of mechanical milling and alloying, were performed to realize the optimum microstructure, nano-heterogeneous distribution of heavy REMs (Dy or Tb) within a grain, and economically alloyed composition of magnets, which assumes, in particular, the distribution of heavy REMs within the near-grain boundary areas. Such a heavy-REM distribution allows us to (1) locally increase the coercive force and decrease the probability of the formation of reverse domains at grain boundaries; (2) limit the substitution of heavy REM for neodymium in the matrix phase and, thus, decrease the probability of decreasing magnetization and remanence; and (3) decrease the amount of heavy REMs, which is required to reach the given increase in the coercive force. The latter circumstance determines the possibility of the development of physico-chemical and technological foundations of resource-saving technology, the possibility of decreasing the material costs and prices of products manufactured from the new alloys, and the possibility of substantially widening the functionality of the materials.

Thus, by applying compositions with a heavy rare-earth metal, the outer region of the Nd2Fe14B matrix grains was enriched during the sintering process and substitutes for Nd were used in the matrix grains to form the (Nd,Dy)2Fe14B core–shell phase.

This paper focused on optimizing the microstructure of the near-stoichiometric Nd2Fe14B-based magnet, which included the grain boundary diffusion and grain boundary structuring processes via the application of a hydrogenated Dy3Co0.6Cu0.4H*x* composition added to a powder mixture.

#### **2. Experimental**

The strip casting technique was used for the preparation of the base Nd-24.0, Pr-6.5, Dy-0.5, B-1.0, Al-0.2, Fe-balance alloy (wt.%). The strip-cast alloy was subsequently subjected to hydrogen decrepitation process, which was realized during heating to 270 ◦C in a hydrogen flow at a pressure of 0.1 MPa and holding at this temperature for 1 h.

The Dy3(Co1−xCu*x*) alloy with *x* = 0.4 was produced by the arc melting of the starting components (distilled Dy of 99.9% purity, Co of ≥ 99.25% purity, and oxygen-free Cu of 99.95% purity) in an argon atmosphere using a water-cooled copper bottom and a non-consumable tungsten electrode. The ingot was homogenized at 600 ◦C for 90 h and subjected to hydrogenation under conditions used for the strip-casting alloy, namely, upon heating to 270 ◦C in a hydrogen flow at a pressure of 0.1 MPa and subsequent 1 h heating at this temperature (Regime 1 was used to manufacture the magnet), and upon heating to 700 ◦C in a high-purity hydrogen atmosphere and holding at this temperature for 1 h in a glass Sieverts-type apparatus (Regime 2 was used for investigations). In the case of heating at 700 ◦C, the hydrogenation up to the Dy3Co0.6Cu0.4H*x* composition with *x* = 8.26 was realized. It is expected that such a hydrogen content accords with the complete hydrogenation of dysprosium to a dysprosium hydride.

The mixture of hydrogen-decrepitated strip-cast Nd(RE)–Fe–B alloy and the Dy3Co0.6Cu0.4H*x* alloy (Regime 1) was milled for 40 min to an average particle size of 3 μm using a vibratory mill and isopropyl alcohol medium. After wet pressing of the pulp in a transverse magnetic field of 1500 kA/m, compacts were sintered at 1080 ◦C for 2 h and optimally heat treated (HT) at 500 ◦C for 2 h. Then, samples of the magnet were subjected to low-temperature heat treatment in the temperature range 400–900 ◦C, with subsequent quenching in N2.

The phase composition of the Dy3Co0.6Cu0.4 and Dy3Co0.6Cu0.4H*x* (*x*=8.26) alloys was investigated by X-ray diffraction (XRD) analysis using an Ultima IV (Rigaku») diffractometer (equipped with a "D/teX" detector, CuKα radiation) and a Philips X'Pert 1 diffractometer, respectively; the scanning step was 0.001◦. X-ray diffraction patterns were processed, and the phase composition of the alloy was determined using PowderCell software. Data on the crystal structure type, lattice parameters, and the crystallographic positions of atoms in the Dy–Co, Dy–Cu, and H–Dy systems [14–16] were used to simulate theoretical XRD patterns.

An Quanta 450 FEG high-resolution field emission gun scanning electron microscope (FEI Company, Fremont, USA) equipped with an energy-dispersive spectroscopy (EDS, EDAX Inc., Mahwah, USA) microprobe was used to investigate the structure, chemical composition, and distribution of magnet components (X-ray mapping) of the addition and magnet sample. The mean particle size was evaluated by means of a MasterSizer 3000 laser diffraction particle size analyzer (Malvern Panalytical Ltd, Malvern, United Kingdom). The hysteretic properties of the magnet sample were measured at room temperature (RT) using an automatic hysteresis graph system MH-50 (Walker Scientific Inc., Worcester, USA). The differential thermal analysis (DTA) and thermogravimetric analysis were performed under an argon atmosphere with a heating/cooling rate of 30 ◦C/min using a STA 449 F3 Jupiter installation (Netzsch Holding, Selb, Germany).

#### **3. Results and Discussion**

#### *3.1. X-Ray Di*ff*raction Analysis*

Figure 1 shows the X-ray diffraction pattern of the Dy3Co0.6Cu0.4 alloy subjected to prolonged annealing in an argon atmosphere. The reflections belong to the main Dy3(Co,Cu) phase and the Dy(Cu,Co) phase based on DyCu [14,15]. The analysis of the crystal structures of the found compounds and theoretical XRD patterns constructed for the simulated structures allowed us to determine variations in the lattice parameters of the Dy(Cu1−yCoy) and Dy3(Co1−xCux) phases alloyed with Co and Cu, respectively (see Table 1). As seen, the alloying of the binary compounds with Co and Cu did not change the crystal structure type of the compounds. In accordance with the binary phase diagrams [14,15], the phases present in the alloy are alloyed compositions of the binary compounds.

**Figure 1.** X-ray diffraction pattern of the Dy3Co0.6Cu0.4 alloy.

**Table 1.** The crystal structure type and lattice parameters of the phases in the Dy3(Co0.6Cu0.4) alloy.

The phase composition of the alloy was also confirmed by the EDS microanalysis, see Figure 2 and Table 2. The microstructure consisted of Dy3(Co1−xCux) (*x* ~ 0.4) dendrites (point 1 in Figure 2) and Dy(Cu1−yCoy) + Dy3(Co0.6Cu0.4) mixture (point 2 in Figure 2) found in the interdendritic regions. The composition of the Dy(Cu1−yCoy) phase cannot be accurately determined by the EDS analysis because of its small size, since the surrounding matrix is analyzed along with this very small inclusion. However, the increased content of copper is evident in this mix area.

**Figure 2.** Scanning electron microscopy images of the microstructure of Dy3Co0.6Cu0.4 alloy subjected to prolonged annealing in an argon atmosphere: (**a**) metallographic section, and (**b**) fracture surface.

**Table 2.** The chemical composition (at.%) of phases found in the microstructure of the Dy3Co0.6Cu0.4 alloy (the average value from three analysis).

As is shown in Table 1, the substitution of Cu for Co in Dy3(Co1−xCux) (with regard to the solubility of Cu and Co in Dy3Co and DyCu, respectively) changed the lattice parameters: the lattice parameters *b* and *c* increased as the radius of Cu atoms (0.128 nm) was higher than that of the Co atoms (0.125 nm), whereas the lattice parameter *a* decreased. This is likely to be due to the fact that copper atoms substitute for cobalt atoms only at certain sites.

We assumed that the solidification of the alloy occurs via the primary formation of the Dy3Co-based phase by peritectic reaction; the DyCu-based compound is the secondary phase. According to the Co–Dy phase diagram, the solidification path may include the formation of the Dy12Co7-based phase by peritectic reaction.

#### *3.2. Interaction of Dy3(Co,Cu) Alloy with Hydrogen*

The saturation of the Dy3Co0.6Cu0.4 alloy with hydrogen led to the embrittlement of the alloy (i.e., the powder material suitable for further introduction of the composition into the Nd–Fe–B magnetic alloy powder during cooperative milling was obtained). Figure 3a shows the X-ray diffraction analysis data for the Dy3Co0.6Cu0.4 alloy subjected to hydrogenation (Regime 2). The hydrogenated

composition contained DyH2 [17] and DyH3 [18] hydrides. Other reflections corresponded to the Dy3(Co,Cu) phase; it is likely that small quantities of the Dy3(Co,Cu) and Dy(Cu,Co) phases did not react with hydrogen. After hydrogenation, copper and cobalt may be present in the form of a fine mixture.

**Figure 3.** X-ray diffraction pattern of the Dy3(Co,Cu) alloy after saturation with hydrogen (Regime 2) (**a**) and after the thermal dehydrogenation process (DTA) (**b**).

Figure 3b shows the X-ray diffraction analysis data of the alloy Dy3Co0.6Cu0.4H*x* subjected to thermal dehydrogenation (upon heating during DTA). The sample was heated up to 700 ◦C (Figure 3). After heating, the presence of DyH2 and small quantities of the Dy3(Co,Cu) and Dy(Cu,Co) phases were detected; DyH3 was absent. The presence of a thin mechanical mixture of Cu and Co is also possible.

According to the DTA data (Figure 4), the decomposition of DyH3 started at a temperature of ~314 ◦C, which agreed with the literature data [16]. Between ~314 ◦C and ~690 ◦C, no thermal effects were identified. Above ~690 ◦C, in accordance with the Dy–H [16] diagram, the solid solution of hydrogen in dysprosium decomposed to form dysprosium. However, the thermal effects at temperatures above 600 ◦C can correspond to the melting of one of the metallic phases of the alloy; nevertheless, the thermal effect corresponding to ~690 ◦C is accompanied by a significant weight loss. The observed formation of Dy hydrides indicates the possibility of the hydrogenated Dy3Co0.6Cu0.4 alloy to be used as additions in manufacturing sintered Nd–Fe–B magnets.

**Figure 4.** Thermogravimetric analysis (TG) and differential thermal analysis (DTA) curves of the Dy3(Co,Cu)H8.26 sample.

#### *3.3. Microstructure and Electron Microprobe Analysis of Sintered NdFeB-Based Magnet*

In accordance with the microprobe analysis data shown in Table 3, the microstructure of a magnet prepared from the powder mixture with 2 wt.% Dy3Co0.6Cu0.4H*x* (Regime 1) was characterized by the presence of four structural components differing in the chemical composition, see Figure 5 (the phases are indicated by red numbers).

**Table 3.** Chemical composition of phases observed in the structure of the Nd–Fe–B sintered magnet prepared from a powder mixture with 2 wt.% of Dy3Co0.6Cu0.4Hx (the values averaged for three measurements are presented).

**Figure 5.** The microstructure of the Nd–Fe–B sintered magnet prepared from the powder mixture with 2 wt.% Dy3Co0.6Cu0.4Hx; (**a**) phase 1—2:14:1 phase grains, phase 2—Nd-rich phase, phase 3—oxide phases; (**b**) phase 4—a phase based on Fe–Nb; (scanning electron microscopy; marked phases correspond to those in Table 2).

The chemical composition of matrix grains (Phase 1 in Figure 5a) was close to the stoichiometric (Nd,R)2Fe14B composition. The presence of Dy in the matrix alloy did not allow us to unambiguously conclude the formation of the core–shell structure, but the presence of cobalt in 2:14:1 phase grains demonstrates the possibility of micro-alloying through the use of hydrogenated low-melting Co-containing compounds (the melting temperature was lower than the sintering temperature of Nd–Fe–B magnets). The Nd-rich phase (Phase 2 in Figure 5a) was characterized by a variable composition. Phase 3 (Figure 5a) corresponded to the oxide phases. In accordance with the literature data [19,20], they may be based on NdO, Nd2O3, or NdO2. The presence of a phase based on Fe–Nb in triple junctions (TJ) was observed (Phase 4, Figure 5b). This fact may be related to impurities in the industrially prepared alloy matrix.

The distribution of rare earth elements, Co and Cu in the matrix grains, and in the intergranular Nd-rich phases (phase 2 in Figure 5a) in the sintered magnets prepared from the powder mixture with 2 wt.% of Dy3Co0.6Cu0.4H*x* addition was also investigated by X-ray mapping (see Figure 6). The nonuniform Dy distribution within the 2:14:1 phase grains could be observed. The depletion of triple junctions of Co and their enrichment in Cu should be noted in the case of the addition of Dy3Co0.6Cu0.4H*x*. The presence of reactive Dy powder (originating from DyH2 that was decomposed during sintering) ensures the diffusion of Dy atoms to the 2:14:1 phase lattice, since the atomic radius of Dy atoms is lower than that of Nd atoms. This led to ousting Nd atoms to peripheral areas. The diffusion coefficient of Nd atoms is lower than that of Dy atoms [21]; thus, the diffusion of Dy is more significant. Such an inequality of diffusion flows of atoms caused lattice stresses and resulted in the inhomogeneous Dy and Nd(Pr) distribution over the 2:14:1 phase grains. The core–shell structure (Dy-enriched shell and Dy-depleted core) is evident in Figure 6.

**Figure 6.** Co, Cu, and Dy mapping in 2:14:1 phase grains and triple junction phases of the Nd–Fe–B sintered magnet prepared from the powder mixture with 2 wt.% Dy3(Co,Cu). The red circle indicates the depletion of 2:14:1 phase grain in Dy (i.e. the formation of core-shell structure).

**Figure 7.** (**left**) Line chemical analysis over the triple junction phase and (**right**) corresponding SEM image with the analysis direction marked.

The other components of the Dy3Co0.6Cu0.4H*x* composition (i.e., Cu and Co) are also useful additions for Nd–Fe–B-based magnets. It is evident from Figures 6 and 7 that Co evinced the tendency to incorporate the 2:14:1 phase grains, while the Cu enriched triple junction phases. The role of Cu in the grain-boundary restructuring and positive effects of Co on the coercivities of Nd–Fe–B magnets were reported in our previous work [22] and were also considered in [23–31].

#### *3.4. Dependence of the Coercive Force (jHc) on the Heat Treatment Temperature*

The magnetic properties (j*H*c) of the magnets (see Table 4 and Figure 8) prepared with the hydrogenated Dy3Co0.6Cu0.4 alloy were lower than those in the case of the application of the addition of the DyH2 [31]. One of the causes is the incomplete hydrogenation of the alloy (see Figure 3, XRD data) and, therefore, the incomplete occurrence of the grain boundary diffusion of the available Dy. The small quantity of the Dy3(Co,Cu) phase present in the Dy3Co0.6Cu0.4 alloy was subjected to hydrogenation. However, the value of *B*r in the case of Dy3Co0.6Cu0.4H*x* was higher than that in the case of DyH2, which may be due to a difference in the Dy content in the chemical composition of the 2:14:1 phase. The difference in the rare-earth metal and Cu contents in the Nd-rich phases provided a lower value of *H*k in the case of magnets with 2 wt.% Dy3Co0.6Cu0.4Hx. The hysteretic properties of the Nd–Fe–B magnet, without the addition of hydride after optimal HT, are also shown in Table 4 for comparison.

**Figure 8.** Magnetization reversal portions of hysteresis loop for the Nd–Fe–B sintered magnets prepared from the powder mixture with 2 wt.% Dy3Co0.6Cu0.4Hx.

We assumed that the optimal HT for magnets of this type was in the range of 475 to 500 ◦C, as in the case of the magnets considered in [32–34]. Subsequent HT in this temperature range, which is performed after the optimal heat treatment (500 ◦C), will lead to an increase in the coercive force of magnets with 2 wt.% Dy3Co0.6Cu0.4Hx.

**Table 4.** Hysteretic properties of sintered magnets prepared from the powder mixtures with 2 wt.% Dy3Co0.6Cu0.4H*x* and DyH2 and optimally heat treated at 500 ◦C for 2 h; *B*r = remanence of magnetic flux density; j*H*c = coercivity of magnetic polarization; *H*k = parameter adopted as a criterion of coercivity (i.e., the magnetic field determined at 0.9 × *B*r); (*BH*)max = maximum energy product; HT = heat treatment.

∗ The initial Nd–Fe–B alloy contains 0.5 wt.% Dy.