**Mechanical Alloying: Processing and Materials**

Editor **Joan-Josep Su ˜nol**

MDPI • Basel • Beijing • Wuhan • Barcelona • Belgrade • Manchester • Tokyo • Cluj • Tianjin

*Editor* Joan-Josep Sunol ˜ Universitat de Girona Spain

*Editorial Office* MDPI St. Alban-Anlage 66 4052 Basel, Switzerland

This is a reprint of articles from the Special Issue published online in the open access journal *Metals* (ISSN 2075-4701) (available at: https://www.mdpi.com/journal/sustainability/special issues/Bioenergy Biofuels).

For citation purposes, cite each article independently as indicated on the article page online and as indicated below:

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## **Contents**


## **About the Editor**

**Joan-Josep Su ˜nol** is full professor in Applied Physics at the University of Girona, Spain (since 2020). He works in the Department of Physics of the Polytechnic Engineering School. He obtained the Ph.D. degree at the Autonomous University of Barcelona in 1996. Regarding publications and congresses, more than 200 articles are indexed (SCI) journals and more than 300 communications has been published in scientific congress. Dr. Sunol has been the president of the Spanish GECAT ˜ (Thermal Analysis and Calorimetry Group) since 2015 and the coordinator of the Materials and Thermodynamics research group of the University of Girona. He was head of the Physics Department of the University of Girona from 2007 to 2010.

## **Preface to "Mechanical Alloying: Processing and Materials"**

This book is a compilation of recent articles linked to the production and the structural and functional characterization of alloys and compounds produced by mechanical alloying. In one of the works in this Special Issue, a non-complex model of mechanical alloying was applied to compare the final microstructure of two nanocrystalline alloys as a function of the energy transfer in two milling devices: planetary and shaker. Regarding the production of materials, some examples are: (a) milled Mn–Al-based alloys were introduced in dissolutions with azoic dye (wastewater treatment redox processes); (b) controlled annealing provoked the relaxation of mechanically induced strain or the recrystallization from an amorphous phase.

Mechanical alloying can be considered as a step-in powder metallurgy process (high-temperature press, spark plasma, or microwave sintering). Usually, the main objective is to obtain the desired microstructure to optimize the mechanical and functional properties of the material.

As Guest Editor of this Special Issue, I am very happy with the final result, and hope that the selected papers will be useful to researchers working on mechanical alloying as a processing technique of materials with improved functional properties. I would like to warmly thank the authors of the eight articles in this Special Issue for their contributions, and all of the reviewers for their efforts in ensuring high-quality publications. Finally, thanks to the editors of Metals for their continuous help, and to the Metals editorial assistants for their valuable and inexhaustible engagement and support during the preparation of this volume.

> **Joan-Josep Su ˜nol** *Editor*

## *Editorial* **Mechanical Alloying: Processing and Materials**

**Joan-Josep Suñol**

Department of Physics, C/Universitat de Girona 3, Universitat de Girona, 17003 Girona, Spain; joanjosep.sunyol@udg.edu; Tel.: +34-972-419-757

## **1. Introduction and Scope**

Mechanical alloying is a technique involving the production of alloys and compounds, which permits the development of metastable materials (with amorphous or nanocrystalline microstructure) or the obtention of solid solutions with extended solubility. The elements or compounds to be mix (usually as powders) were introduced in jars, together with a few numbers of balls.

Regarding the scope of this Special Issue, so many options were given to the potential authors:


Finally, only height articles have been published. Nevertheless, the set of materials, characterization and applications described in the manuscripts provides a wide spectrum of the potential of this processing technique.

## **2. Contributions**

Regarding the modelling of the milling process, the main problem is due to the high quantity of processing parameters to be controlled, which include the filling factor of the jars, the material of the jars and balls, the milling atmosphere, the milling time, the milling intensity, the ball to powder weight ratio (BPR), the number and diameter of balls, the temperature inside the jars, the local temperature on interactions between powder and balls, the optional change in the sense of the rotation of the jars, the on–off switch periods, the controlled addition of a process control agent (PCA) that can help in grain refinement and act as a surfactant, the frequency of collisions between balls, in which powdered particles are involved, and so on. Thus, it is quite difficult to model the energy or powder transfer during the milling process. Furthermore, there are ball milling devices with different geometries: shaker mills, planetary mills. Likewise, the interaction between of the powders with balls (and/or jar internal wall) can be facilitated by abrasion or percussion. For kinetic energy, the velocity of the balls has a broad distribution. For this, all models are usually based upon estimation. One of the works in this Special Issue applies a non-complex model to compare the final microstructure of two Fe-X-Nb-Cu (X = Nb, Ni-Zr) alloys as a function of the energy transfer in two milling devices: planetary and shaker. In this work, the shaker mill is more energetic [1].

Regarding the production of materials, the alloys and compounds that are produced are obtained in a powder shape. Milling usually favors a reduction in the grain size (except for very ductile materials) and the formation of smooth surfaces with high specific

**Citation:** Suñol, J.-J. Mechanical Alloying: Processing and Materials. *Metals* **2021**, *11*, 798. https:// doi.org/10.3390/met11050798

Received: 6 May 2021 Accepted: 13 May 2021 Published: 14 May 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the author. 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 (https:// creativecommons.org/licenses/by/ 4.0/).

surface/volume ratio. The size distribution of the powders can be checked by scanning electron microscopy. One of the problems associated with the milling process is contamination from the milling tools and atmosphere. Additional oxygen contamination can be induced after the extraction of the powders from the jars. Thus, a shift in the composition can be produced. This effect is checked with microanalysis techniques.

Sometimes the powdered compounds can be directly used in specific applications without additional treatments. As an example, Mn–Al-based alloys were introduced in dissolutions with azoic dye. The interaction with the metallic particles favors the decolorization process of the dyes by breaking the azo bond of the macromolecule [2].

In order to obtain an improvement in the functional properties of the alloys and compounds, sometimes controlled annealing is needed. Furthermore, annealing provokes the relaxation of the mechanical induced strain. In one of the articles in this Special Issue, annealing was performed at 700–800 ◦C in high-nitrogen chromium-manganese steels [3]. The austenite phase of the steel was stabilized. Likewise, the annealing objective is the recrystallization of an amorphous phase [4]. The development of the desired crystallographic phase is associated with the influence of the microstructure in the functional response of the alloy. Some Mn-Co(Fe)-Ge(Si) alloys have a martensitic transformation coupled with a magnetic transition favoring an improved magnetocaloric effect.

Mechanical alloying can be a step-in powder metallurgy process. The powders (as obtained after milling) can be compacted at high pressure. An alternative is the spark plasma sintering process (SPS). Refractory high-entropy alloys are produced to maximize the strength, yield strength and fracture strain [5]. An innovative technique is the microwave sintering of previously compacted powders [6]. Al-Y2O3 nanocomposites produced by mechanical alloying and pressing were sintered in a microwave sintering oven. The processing conditions were heating rate of 10 K/min until 550 ◦C and a dwell time of 30 min. The main objective is to optimize the mechanical properties: hardness, yield strength, ultimate compression strength and compressive strain.

Two of the selected articles are reviews. One is devoted to Fe-Cr based alloys and their consolidation at high temperature [7]. In these materials, the technological objective is to improve the resistance to corrosion. Nanocrystalline alloys have higher resistance than microcrystalline alloys. A system with improved resistance is Fe-Cr-Ni-Zr.

A second article revises the hydrogen absorption behavior and the absorption/desorption kinetics of metal hydrides produced by mechanical alloying [8]. It is a critical overview on the effect of mechanical alloying in binary (CaH2, MgH2, etc.) and ternary (Ti-Mn-N and Ca-La-Mg-based systems) hydrides. Sometimes the technological process has multiple steps, involving: mechanical alloying, heat treatment, a second mechanical alloying process, degassing and, finally, extrusion.

## **3. Conclusions and Outlook**

As a main conclusion, it is necessary to acknowledge the variety of alloys and compounds produced by mechanical alloying: Fe-X-B-Cu (X = Nb, NiZr) nanocrystalline alloys, mixtures of the binary Fe-Mn and Fe-Cr alloys with the nitrides CrN (Cr2N) and Mn2N, Mn-Al-Co and Mn-Al-Fe alloys, non-equiatomic refractory high entropy alloy (W35Ta35Mo15Nb15)95Ni5, nanocrystalline MnCo0.8Fe0.2Ge1−*x*Si*x*, nanocrystalline Fe-Cr alloys, Al–Y2O3 nanocomposites and hydride-forming alloys. Regarding the study of their properties, it is important to improve mechanical properties, hydrogen absorption, magnetocaloric effect and resistance to corrosion. The processing parameters affect the final microstructure of the material, and the microstructure affects the functional response. Likewise, the powders can be consolidated (press, spark plasma sintering, microwave sintering) to obtain bulk materials. Further investigations should be performed to gain a deeper knowledge of the influence of the milling parameters and to analyze the option to develop new advanced materials for specific applications.

As Guest Editor of this Special Issue, I am very happy with the final result, and hope that the present selected papers will be useful to researchers working on mechanical

alloying as processing technique of materials with improved functional properties. I would like to warmly thank the authors of the eight articles in this Special Issue for their contributions, and all of the reviewers for their efforts in ensuring high-quality publications. Finally, thanks also to the editors of *Metals* for their continuous help, and to the *Metals* editorial assistants for the valuable and inexhaustible engagement and support during the preparation of this volume. In particular, my sincere thanks go to Toliver Guo for his help and support.

**Conflicts of Interest:** The author declares no conflict of interest.

## **References**


## **Microstructure and Compressive Behavior of Al–Y2O3 Nanocomposites Prepared by Microwave-Assisted Mechanical Alloying**

## **Manohar Reddy Mattli 1, R. A. Shakoor 1,\*, Penchal Reddy Matli <sup>2</sup> and Adel Mohamed Amer Mohamed <sup>3</sup>**


Received: 4 March 2019; Accepted: 3 April 2019; Published: 5 April 2019

**Abstract:** In this study, Al–Y2O3 nanocomposites were synthesized via mechanical alloying and microwave-assisted sintering. The effect of different levels of yttrium oxide on the microstructural and mechanical properties of the Al–Y2O3 nanocomposites were investigated. The density of the Al–Y2O3 nanocomposites increased with increasing Y2O3 volume fraction in the aluminum matrix, while the porosity decreased. Scanning electron microscopy analysis of the nanocomposites showed the homogeneous distribution of the Y2O3 nanoparticles in the aluminum matrix. X-ray diffraction analysis revealed the presence of yttria particles in the Al matrix. The mechanical properties of the Al–Y2O3 nanocomposites increased as the addition of yttria reached to 1.5 vol. % and thereafter decreased. The microhardness first increased from 38 Hv to 81 Hv, and then decreased to 74 ± 4 Hv for 1.5 vol. % yttria. The Al–1.5 vol. % Y2O3 nanocomposite exhibited the best ultimate compressive strength and yielded a strength of 359 ± 7 and 111 ± 5 MPa, respectively. The Al–Y2O3 nanocomposites showed higher hardness, yield strength, and compressive strength than the microwave-assisted mechanically alloyed pure Al.

**Keywords:** aluminum; yttrium oxide (yttria); mechanical alloying; microwave sintering; microstructure and mechanical properties

## **1. Introduction**

Metal matrix composites (MMCs) find noteworthy applications in many engineering sectors due to their superior properties such as high strength, high-temperature capability, specific modulus, and good wear resistance compared to monolithic base materials. The mechanical performances of MMCs often show greater improvement than can be achieved by conventional strengthening methods in monolithic alloys [1–4].

Aluminum (Al)-based metal matrix composites (AMMCs) are an excellent choice for automotive, aerospace, defense, and nuclear power sectors because of their lightweight and favorable mechanical, thermal, and physical properties. Aluminum (Al)-based metal matrix composites are capable of achieving high strength, high-fatigue resistance, high-wear and corrosion resistance, and good compatibility with various manufacturing processes [5–8].

At present, ceramic particle-reinforced Al-matrix nanocomposites have been prepared primarily by mechanical alloying, forging, and casting routes [9–11]. Among these methods, mechanical alloying (MA) has been widely used to fabricate Al-matrix nanocomposites due its cost-effectiveness, simplicity, and its ability to improve the properties vis-a-vis those of the unreinforced matrix [12,13]. There are many sintering techniques such as conventional, spark plasma, vacuum, and microwave sintering processes [14–17]. Among these techniques, the microwave sintering process is a heating method that offers the ability to balance the radiant and microwave heating effects. In this process, heat is generated within the sample by rapid oscillation of dipoles at microwave frequencies. Microwave sintering provides efficient internal heating, and energy is supplied directly to the material. Therefore, this process avoids the significant temperature gradient between the surface and interior. Microwave sintering is a high-technology heating process that can save both energy and time [18].

In AMMCs, the most common types of reinforcement that can be used are SiC, Si3N4, Y2O3, TiC, and Al2O3 [19–23]. Among these ceramics, Y2O3 was selected as the reinforcement to be used in this study due to its high strength, hardness, melting point, and thermal conductivity [24–26]. Yttria is an air-stable particle, white in color and solid in substance. By adding the yttria to the aluminum, the strength, corrosion resistance, and wear properties are improved [27]. Yttria is well sintered to a high density and low coefficient of thermal expansion, and has excellent strength properties [28,29]. According to the authors' knowledge, there are no reports in the literature on Al–Y2O3 nanocomposites processed by mechanical alloying and microwave sintering.

Therefore, in this current research, Al–Y2O3 nanocomposites were prepared by mechanical alloying and microwave heating, and the effect of Y2O3 addition on the microstructure and mechanical performance of Al–Y2O3 nanocomposites were investigated.

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

Pure Al (99.5% purity, with an average particle size of 10 μm) and Y2O3 nanoparticles (99.99% purity, with an average particle size of 50–70 nm) were purchased from Alfa Aesar (Tewksbury, MA, USA) and selected as raw materials for the synthesis of Al–Y2O3 nanocomposites.

Aluminum–yttria composites were prepared with 0, 0.5, 1.0, 1.5, and 2.0 vol. % yttria nanoparticle contents. The mixture of powders was blended at room temperature using a Planetary Ball Mill (PM 200) for 2 h, with a rotation speed of 200 rpm. No balls were used during the blending of powders. The mixed powder (~1.0 gm) was compacted into cylindrical pellets by applying a pressure of 50 MPa with a holding time of 1 min. The compacted cylindrical pellets were sintered in a microwave sintering furnace at a temperature of 550 ◦C with a heating rate of 10 ◦C/min and providing a dwell time of 30 min. The microwave furnace had an alumina insulation and silicon carbide susceptor. The silicon carbide susceptor was used to increase the heating rate and hybrid heating. Alumina insulation prevents heat loss and is used as well to protect the interior walls of the microwave oven. The compacted pellets were placed at the center of the cavity and sintering was conducted at the multimode cavity [30]. Figure 1 shows the schematic representation of the microwave sintering furnace.

The density of the sintered samples was calculated using Archimedes' principle. The porosity of the samples was calculated by the theoretical and experimental density of the composite samples. The X-ray diffraction (XRD, PANalytical X'pert Pro, PANalytical B.V., Almelo, The Netherlands) analysis was performed to identify the phases present in Al–Y2O3 nanocomposites. The XRD patterns were recorded in the 2θ range of 20–90◦ with a step size of 0.02◦ and a scanning rate of 1.5◦/min. The microstructural characterization and determination of the distribution of the yttria nanoparticles in the aluminum matrix were carried out using scanning electron microscopy (SEM, JeolNeoscope JSM6000, Tokyo, Japan) and energy dispersive X-ray spectroscopy (EDS, Tokyo, Japan).

The microhardness of the Al–Y2O3 nanocomposites was determined using Vickers microhardness tester (MKV-h21, USA). Microhardness analysis was carried out to investigate the effect of yttria on the hardness of the Al–Y2O3 nanocomposite, carrying the load of 25 gf and a dwell time of 10 s, for each sample with an average of five successive indentations. Compressive strength analysis was performed at room temperature using a universal testing machine (Lloyd), under an engineering strain rate of 10−4/s.

The respective data of each sample were obtained by an average of three successive values of test results. From the load–displacement curves, 0.2% offset compressive yield strength (CYS), ultimate compressive strength (UCS), and compressive strain were determined.

**Figure 1.** Schematic diagram of a microwave sintering furnace.

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

## *3.1. Density and Porosity of Al–Y2O3 Nanocomposites*

Density and porosity values of the microwave sintered Al–Y2O3 nanocomposites with different contents of yttria in the Al matrix are shown in Table 1.


**Table 1.** Density and porosity of Al–Y2O3 nanocomposites.

It can be observed that the density of the composite gradually increased with the increase of the yttria content since the density of yttria (5.01 g\cc) is higher than that of Al (2.70 g\cc). Generally, the higher relative density of sintered samples influences the mechanical properties of the composites. The porosity of the composites decreased by increasing the amount of yttria content. The decrease in porosity with increasing yttria content shows that the presence of the hard yttria particles did not impair the densification of the Al powder [31]. Microwave heating was one of the main reasons for the low porosity of the synthesized composites.

## *3.2. XRD Analysis of Al–Y2O3 Nanocomposites*

The X-ray diffraction (XRD) patterns of the microwave sintered pure Al and Al–Y2O3 nanocomposites with different amounts of Y2O3 are shown in Figure 2a. Figure 2b shows the enlarged patterns of the Al–1.5 vol. % Y2O3 nanocomposite. The XRD patterns clearly indicate the presence of Y2O3 nanoparticles in the Al composite matrix. Due to the small volume of yttria reinforcement present in these composites, the yttria peaks were very small compared to the aluminum matrix peaks. Also, it can be seen that the intensity of the yttria diffraction peaks increased with the increasing of yttria percentage. The XRD results show that the main elements of Al (higher peak) and Y2O3 (lower peak) are present in Al–Y2O3 nanocomposites.

**Figure 2.** (**a**) X-ray diffraction (XRD) pattern of Al–Y2O3 nanocomposites, (**b**) enlarged pattern of Al–1.5vol. %Y2O3 nanocomposites [32,33].

## *3.3. SEM Analysis of Al–Y2O3 Nanocomposites*

The SEM and EDS images of the microwave sintered Al–Y2O3 nanocomposites with different contents of yttria are shown in Figure 3. The results of microstructural characterization revealed that yttria particulates were present individually and in relatively smaller clusters indicating an improvement in their distribution. The EDS analysis confirms the aluminum and yttria particles present in the Al matrix. The EDS mapping spectrum of all nanocomposites were mainly composed of Al, Y, and O elements, as shown in Figure 3b,d,f. The microcracks were restricted by the presence of hard and homogeneous yttria particles in the Al-matrix and influenced the microstructure and mechanical properties of Al–Y2O3 nanocomposites. The specimen with 2 vol. % of yttria particles shows the decreasing of the interparticle distances as the concentration of the nanoparticles increased.

## *3.4. Microhardness of Al–Y2O3 Nanocomposites*

Vickers microhardness was measured on all specimens to study the effect of Y2O3 content on the microhardness. Figure 4 shows the results of the microhardness of the Al–Y2O3 nanocomposites with different content of yttria. From the Table 2, the microhardness of the composite increased as the yttria increased of up to 1.5 vol. % and then decreased at 2.0 vol. % Y2O3. The considerable increase in hardness could be attributed to the presence of homogeneously distributed hard ceramic nanoparticles and dispersion hardening effect [34]. Al–2.0 vol. % Y2O3 nanocomposites show a decreased microhardness value, which was mainly due to the agglomeration of the yttria and increasing presence of clustering of yttria in the case of the Al matrix [35]. The microhardness of the microwave sintered samples in this study was found to be higher than the vacuum sintering and arc-melting samples [36].

The increment of microhardness in the composite materials was due to the presence of hard ceramic particles.

**Table 2.** Microhardness, yield strength, and ultimate compressive strength of Al–Y2O3 nanocomposites.


**Figure 3.** Typical micrographs and corresponding energy dispersion elemental mapping analysis of (**a**–**f**) Al–Y2O3 (1, 1.5, and 2 vol. %) nanocomposites.

**Figure 4.** Microhardness of Al–Y2O3 nanocomposites.

## *3.5. Compressive Analysis of Al–Y2O3 Nanocomposites*

The compressive test was conducted on the microwave sintered pure Al and Al–Y2O3 nanocomposites and strengths were compared. Figure 5a shows the engineering stress–strain curves of the Al–Y2O3 nanocomposites with different content of yttria. Figure 5b shows the corresponding mechanical data of Al–Y2O3 nanocomposites.

**Figure 5.** (**a**) The compressive stress–strain curves and (**b**) strength (yield and ultimate) of the Al–Y2O3 nanocomposites.

The yield strength and ultimate compressive strength of Al–Y2O3 nanocomposites show increased values up to 1.5 vol. % of yttria then decreased as shown in Table 2. Al–1.5 vol. %Y2O3 nanocomposites show the maximum yield strength (YS) of 126 ± 5 MPa and ultimate compressive strength (UCS) of 374 ± 6 MPa at a uniform strain of ~60%. These results show the improvement of mechanical properties of Al–Y2O3 nanocomposites compared to the pure Al. The increased mechanical properties of the Al–Y2O3 nanocomposites are attributed to the dispersion hardening effect and homogeneous distribution of hard reinforcements in the Al-matrix [37]. Al–2.0 vol. % Y2O3 nanocomposites show a decreased microhardness value, mainly due to the agglomeration of nanoparticles and grain growth [38]. Reinforcement amounts, density, heating mechanisms factors also govern the variation of the mechanical properties. However, compression properties of the microwave sintered Al–1.5 vol. % Y2O3 nanocomposites are interestingly superior to those of other reinforced AMMCs [39–43].

There are several strengthening mechanisms to enhance materials' mechanical properties like hardness and compressive strength of the composite materials. The strengthening of the composites is not only dependent on unique strengthening mechanisms, but it also depends on several strengthening mechanisms.

In the present study, the strengthening mechanism of the Al–Y2O3 nanocomposites mainly depended on dispersion hardening due to the hard yttria particles present in the aluminum matrix. The increase in strength and hardness may be attributable to Orowan strengthening [44,45].

## *3.6. Fractography of Al–Y2O3 Nanocomposites*

Figure 6 shows the fracture surface images of microwave sintered pure Al and Al–Y2O3 nanocomposites under compressive loading. The SEM observations in nanocomposites show typical shear mode fractures and cracks obtained at a 45◦ to the fracture surfaces with respect to the compressive loading axis. It can be observed that the compressive deformations obtained in pure aluminum and aluminum composites with yttria are different, due to the work hardening behavior. The plastic deformations are restricted by the presence of the second phase in Al–Y2O3 nanocomposites [46].

**Figure 6.** Compression fracture surfaces of (**a**) pure Al and (**b**) Al–1.5 vol. % Y2O3 nanocomposites.

### **4. Conclusions**

The Al–Y2O3 nanocomposites were successfully synthesized by mechanical alloying and microwave sintering method. The influence of yttria nanoparticles on the microstructure and mechanical properties of the Al–Y2O3 nanocomposites were investigated in detail. The density of the composites increased with the increasing of yttria content while porosity decreased. The SEM analysis showed the homogeneous distribution of yttria particles in aluminum composites. The Al–Y2O3 nanocomposites exhibited better mechanical properties compared to pure Al. The optimum hardness (81 ± 3 Hv), yield strength (126 ± 5 MPa), and ultimate compression strength (374 ± 6 MPa) and compressive strain (~60%) values were obtained for the Al–1.5 vol. % Y2O3 nanocomposite. This significant enhancement in mechanical properties in Al–1.5 vol. % Y2O3 nanocomposites make them potential candidates for automotive applications.

**Author Contributions:** A.S. and A.M.A.M. proposed the original project and supervised the investigation. M.R.M. and P.R.M. performed the experiments, analyzed the data, and wrote the paper with assistance from all authors. All authors contributed to the discussions in the manuscript.

**Funding:** This publication was made possible by NPRP Grant 7-159-2-076 from the Qatar National Research Fund (a member of the Qatar Foundation). The Qatar National Library funded the publication cost of this article.

**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/).

## **Mechanical Amorphization and Recrystallization of Mn-Co(Fe)-Ge(Si) Compositions**

## **Antonio Vidal-Crespo, Jhon J. Ipus, Javier S. Blázquez \* and Alejandro Conde**

Departamento Física de la Materia Condensada, ICMSE-CSIC, Universidad de Sevilla, 41080 Sevilla, Spain; antvidcre@alum.us.es (A.V.-C.); jhonipus@us.es (J.J.I.); conde@us.es (A.C.)

**\*** Correspondence: jsebas@us.es; Tel.: +34-954-556-029

Received: 18 April 2019; Accepted: 6 May 2019; Published: 8 May 2019

**Abstract:** Mechanical alloying using a planetary ball mill allowed us to obtain two homogeneous systems formed by units with nanometer size and MnCo0.8Fe0.2Ge1−*x*Si*<sup>x</sup>* stoichiometry (*x* = 0 and 0.5). The phase evolution of the systems with the milling time was analyzed using X-ray diffraction. Thermal stability of the final products was studied using differential scanning calorimetry. Room temperature 57Fe Mössbauer spectroscopy was used to follow the changes in the Fe environments. A paramagnetic Co-based amorphous phase developed in both alloys as milling progressed. However, while the presence of Si stabilized the Mn-type phase, mechanical recrystallization was observed in a Si-free composition leading to the formation of a MnCo(Fe)Ge intermetallic (*Pnma* space group) with a crystal size of 7 ± 1 nm. Mössbauer results indicate that Fe atoms migrate from the initial bcc phase to the amorphous and intermetallic phases.

**Keywords:** half-Heusler alloys; mechanical alloying; Mössbauer spectroscopy

## **1. Introduction**

Half-Heusler MnCoGe alloys can show a martensitic transformation from an orthorhombic TiNiSi-type structure (*Pnma* space group) to a hexagonal Ni2In-type structure (*P63*/*mmc*, although it can be also interpreted as a different orthorrombic *Pnma* structure with different lattice parameters to those of the TiNiSi-type structure [1]). Coincidence of such a structural transformation with a magnetic one has been proposed to enhance the magnetocaloric effect exhibited by these systems [2], which can be achieved by compositional tailoring with partial substitution of Fe for Co [3]. However, the formation of the intermetallic phase of interest is not straightforward and long-duration annealing at high temperatures is needed (typically several days at ~1125 K [3–5]).

On the other hand, high entropy alloys (HEAs) are homogeneous solid solutions formed by at least five different elements with atomic fractions between 5 and 35 at. % [6]. In such HEAs, bcc and fcc solid solutions as well as amorphous phases can be observed as product phases when produced by rapid quenching [7] or mechanical alloying [8]. Both amorphous and supersaturated solid solutions are very attractive homogeneous precursor systems to develop stoichiometric intermetallic phases, strongly reducing the annealing time required with respect to the samples obtained by conventional methods [9]. The development of HEAs from half-Heusler compositions has been recently reported for Ti(NiCo)(SnSb) [10], CoMnSn(Cu) [11] and (TiZr)Ni(SnSb) [12] systems.

The aim of this study is to produce homogeneous systems starting from pure powders with MnCo0.8Fe0.2Ge1−*x*Si*<sup>x</sup>* stoichiometry (with *x* = 0 and *x* = 0.5) using mechanical alloying.

## **2. Materials and Methods**

Pure Mn (99.6%, Alfa Aesar, Karlsruhe, Germany), Co (99.99%, Chempur, Karlsruhe, Germany), Ge (99.99%, Chempur, Karlsruhe, Germany), Si (99.9%, Alfa Aesar, Karlsruhe, Germany) and Fe (>99%, Alfa Aesar, Karlsruhe, Germany) powders (5 g) were mixed in hardened steel vials with 10 mm steel balls in an argon atmosphere and ball milled up to 100 h at 250 rpm in a Pulverisette Vario 4 mill (Fritsch, Idar-Oberstein, Germany) with a frequency ratio of −2 and a ball mass to powder ratio, BPR = 10. Compositions were checked using EAGLE III (EDAX, Mahwah, NJ, USA) X-ray microfluorescence equipment. X-ray diffraction (XRD) experiments were performed using a powder diffractometer D8 Advance A25 (Bruker, Karlsruhe, Germany) at room temperature and the radiation employed was Cu Kα. Experimental patterns were fitted using TOPAS software (Version 6, Bruker, Karlsruhe, Germany). No preferential orientation was allowed to preserve the intensity ratio in our powder samples. Transmission 57Fe Mössbauer (MS) spectra at room temperature were obtained using a Wissel spectrometer (Wissel, Starnberg, Germany). Isomer shifts were measured relative to that of a standard foil of pure Fe. Differential scanning calorimetry (DSC) experiments were performed using a DSC7 (Perkin-Elmer, Norwalk, CT, USA) calorimeter at a heating rate of 20 K/min.

## **3. Results**

#### *3.1. X-ray Di*ff*raction*

Figure 1 shows the evolution of the XRD patterns as a function of the milling time. These patterns can be fitted using the Rietveld method assuming the different starting pure phases, except for Si (the lightest among the studied elements), which is not detected even after 1 h milling, indicating its integration to the other phases. Table 1 shows the R-factors of the different fittings and Table 2 shows the lattice parameter of the different phases detected. The diffraction maxima of the starting hcp Co phase rapidly broadens beyond any realistic values of crystal size or microstrains, which is due to the formation of an amorphous phase in both compositions. In order to account for this amorphous phase, we allowed the amorphous halo associated to this phase to evolve directly from the diffraction maxima of the hcp Co phase. Although Rietveld fitting of an amorphous phase could lead to unphysical results of the parameters (e.g., extremely low crystal size or extremely high microstrains), our aim was just to estimate the phase fraction evolution along the milling.

**Figure 1.** X-ray diffraction (XRD) patterns of samples after different times of milling: (**a**) Si-free alloy (**b**) Si-containing alloy. The corresponding differences between the experimental data and the Rietveld fittings are shown below each experimental pattern. The experimental data in black and the fitting in red.

Rietveld fitting (see R-values in Table 1 for each pattern) supplies valuable information concerning phase fraction, lattice parameter, crystal size and microstrains. In the following we will account only for the crystal size as the main factor for peak broadening (i.e., a minimum crystal size is reported). Figure 2 shows the phase fraction evolution with the milling time of the two studied compositions and Figure 3 shows the corresponding crystal size.


**Table 1.** Parameters from Rietveld fittings.

**Table 2.** Average lattice parameters of the crystalline phases detected by XRD. Changes in this parameter with the milling time is of the order of the error bar.


\* Samples heated up to 973 K at 20 K/min in argon flow.

**Figure 2.** Phase fraction from XRD Rietveld analysis as a function of the milling time: (**a**) Si-free alloy (**b**) Si-containing alloy. Lines are a guide to the eye.

**Figure 3.** Crystal size from XRD Rietveld analysis as a function of the milling time: (**a**) Si-free alloy (**b**) Si-containing alloy.

The fraction of the diamond-like Ge phase (*Fd-3m* space group) exponentially decreased with milling time for both studied compositions. The content of bcc-Fe-type phase (*Im-3m* space group) initially reached values above the starting weight fraction of Fe, indicating the migration of other atoms (mainly Co and Si) to this phase. After 10 h of milling, the decrease of the weight fraction for this phase was clear and was no longer detected by XRD after 20 h milling. The evolution of the Mn phase fraction depends on the Si content of the sample. Whereas for Si-free alloy, the Mn-type phase was no longer detected by XRD after 50 h milling; for Si-containing alloy, the Mn-type phase remained almost constant (or even increased) from 10 h up to the maximum time explored in this study (100 h).

As written above, Rietveld fitting showed that the crystal size of the Co-type phase (*P63*/*mmc* space group) rapidly decreased with milling time (below 2 nm after 5 h milling for Si-free alloy and after 10 h milling for Si-containing alloy). Therefore, the diffraction maxima ascribed to this phase no longer describe a crystalline phase but rather amorphous halos ascribed to an amorphous phase. The distinction between crystalline Co-type and Co-base amorphous phase is not clear. Therefore, the two phases are represented together in Figure 2.

Lattice parameters did not change significantly with the milling time but the average values could differ with respect to the values of the pure phases. This should indicate that e.g., Co migration to bcc Fe occurs at the early stages of milling, in agreement with the higher fraction measured for this phase (8–10 wt. %) with respect to the nominal Fe fraction (6.0 and 6.8 wt. % for the alloys without and with Si). In the case of Ge, the measured lattice parameter agreed with that of the pure phase, indicating that this element preserves its purity during its comminution. In the case of the Mn phase, the presence of Si stabilizes it and reduces the lattice parameter with respect to that of the Si-free composition. Average values of lattice parameters of the different phases are shown in Table 2.

### *3.2. Mössbauer Spectroscopy*

Figure 4 shows the evolution of the Mössbauer spectra with the milling time. Two main contributions can be clearly distinguished: a ferromagnetic contribution (FM) and a paramagnetic (PM) one. The FM contribution corresponds to Fe atoms in the bcc Fe(Co) phase as it is confirmed by the hyperfine field HF~33 T. This may indicate that Fe content in crystalline hcp Co at the early stages is negligible (no site is detected). As milling time increased, FM contribution reduced to zero at around 20 h milling, whereas the PM contribution was present since the earlier studied times and progressively increased with milling. Therefore, as the FM contribution is clearly assigned to bcc-Fe type sites, the rest of the phases, which contain Fe, detected by XRD, must be paramagnetic, including the amorphous phase derived from broadening of the hcp-Co diffraction maxima.

**Figure 4.** Room temperature Mössbauer spectra as a function of the milling time of: (**a**) Si-free alloy and (**b**) Si-containing alloy samples.

#### **4. Discussion**

The recrystallization process was detected only in Si-free alloys, leading to the formation of an intermetallic phase: MnCo(Fe)Ge (*Pnma* space group), with a crystal size <10 nm for as-milled samples after 50 and 100 h (see Figure 3a). The presence of the recrystallization phenomenon was confirmed by DSC. Figure 5a shows the DSC scan of Si-free samples milled for 50 h and 100 h, respectively. The transformation heat, |Δ*H*|, of the exothermic peak at ~550 K strongly decreased from the sample milled for 50 h to the sample milled for 100 h (from Δ*H* = −38 to −16 ± 1 J/g, while the amorphous fraction from XRD decreased from 73 to 44%) due to the recrystallization phenomenon. In fact, XRD patterns of samples heated above the exothermic peak showed the intermetallic MnCoGe-type phase as the single phase present except for some traces of MnO (as shown in Figure 6).

In the case of the Si-containing sample after 100 h milling, DSC of Figure 5b shows a minor exothermic peak at ~550 K but the main transformation peak is found at ~620 K (Δ*H* = −114 ± 1 J/g). Samples heated above this temperature transformed to a single bcc solid solution. At higher temperatures, ~850 K, an endothermic and reversible peak was found.

**Figure 5.** DSC scans at 20 K/min for: (**a**) Si-free alloy after 50 and 100 h milling (**b**) Si-containing alloy after 100 h milling.

**Figure 6.** XRD patterns of samples milled for 100 h and heated up to 973 K at 20 K/min: (**a**) Si-free alloy (**b**) Si-containing alloy. Circles identify intermetallic phase, diamonds identify bcc solid solution and asterisks identify MnO phase. The experimental data are in black and the fitting in red. Corresponding differences between the experimental and fitting curves are shown below each experimental pattern.

The formation of such simple structures in Si-containing alloy (amorphous and bcc solid solution) is typically found in HEA. These systems can be characterized by several parameters:

Mixing enthalpy, Δ*Hmix*;

$$
\Delta H\_{\text{mix}} = 4 \sum\_{i,j$$

with *ci* the molar fraction of the *<sup>i</sup>* element in the composition and <sup>Δ</sup>*Hij mix*, the mixing enthalpy between the elements *i* and *j*. For non-metals such as Si and Ge, it is necessary to calibrate Δ*Hij mix* to subtract the extra energy, Δ*Htrans*, required to transform the element from non-metallic to metallic and: Δ*Hij*<sup>∗</sup> *mix* <sup>=</sup> <sup>Δ</sup>*Hij mix* <sup>−</sup> <sup>Δ</sup>*Htrans*/2. For Si and Ge, <sup>Δ</sup>*Htrans* of 34 or 25 kJ/mol [13], respectively.

*Metals* **2019**, *9*, 534

Atomic-size difference, δ:

$$\delta = \sqrt{\sum\_{i=1}^{N} c\_i \left(1 - \frac{r\_i}{\sum\_{i=1}^{N} c\_i r\_i}\right)^2} \,\tag{2}$$

where *ri* is the atomic radius of the *i* element. And Ω parameter:

$$
\Omega = \frac{T\_{\text{m}} \Delta S\_{\text{mix}}}{\Delta H\_{\text{mix}}},
\tag{3}
$$

where *Tm* is the weighted average of the melting temperature of the composition and

$$
\Delta S\_{\rm mix} = -\mathcal{R} \sum\_{i} c\_{i} \ln(c\_{i}),
\tag{4}
$$

Taking the values of *ri* and <sup>Δ</sup>*Hij mix* from [6,13], the results for the quinary composition are Δ*Hmix* = −33 kJ/mol, δ = 5.4% and Ω = 0.60. While the quaternary composition has Δ*Hmix* = −25 kJ/mol, δ = 3.9% and Ω = 0.64. These parameters place our studied alloys close to the bulk amorphous region of HEA depicted in [6]. This could agree with the easy formation of the amorphous phase in our studied samples. However, although the studied quinary composition develops a single bcc phase solid solution after thermal treatment, HEA with such solid solution generally shows larger Ω and less negative Δ*Hmix* [6].

Despite the low content of Fe in the studied compositions (~6.7 at. %), MS can supply some information to confirm the evolution of the phases during milling. As already described, the only FM site detected corresponds to the bcc-Fe(Co) phase and was fitted using a broad sextet centered at HF~33 T, which confirms the rapid transformation of the FM hcp-Co phase to a PM Co-based amorphous phase enriched in Ge and Mn (or at least the negligible Fe content in the residual hcp-Co phase).

The PM contribution was fitted using a doublet with quadrupolar splitting, <*Q*> = 0.48 ± 0.07 mm/s (Si-free alloy) and 0.42 ± 0.15 mm/s (Si-containing alloy). In the case of Si-free alloy, both *Q* and *IS* remain almost constant along the milling process indicating that the Fe sites out of the bcc phase might be similar (i.e., Fe atoms are not expected to migrate to many different phases). In the case of Si-containing alloy, whereas *Q* is almost constant, *IS* becomes more positive as milling time increases. A clear correlation can be observed between the area fraction of the PM contribution and the amorphous fraction measured from XRD for both compositions up to 30 and 20 h of milling for Si-free and Si-containing alloy, respectively. Figure 7 shows this correlation. In the case of Si-free alloy, correlation is preserved for all times when amorphous and intermetallic phase fractions are considered. This indicates that Fe atoms migrate to these two phases in Si-free alloy.

**Figure 7.** Amorphous fraction from XRD Rietveld analysis as a function of PM contribution from MS for Si-free alloy (in red) and Si-containing alloy (in blue). Numbers indicate the milling time in hours for each composition.

#### **5. Conclusions**

Two different compositions MnCo0.8Fe0.2(Ge1−*x*Si*x*) were partially amorphized by mechanical alloying. X-ray diffraction and Mössbauer spectrometry were used to characterize the evolution of the different phases with milling time.

In the case of the Si-free alloy, almost fully amorphization was achieved after 50 h milling, and further milling led to the development of MnCoGe-type intermetallic. Thermal treatment beyond ~650 K led to the growth of this intermetallic, and the alloy became single phase.

In the case of the Si-containing alloy, the Mn-phase fraction remained almost constant from 30 to 100 h milling and the alloy became only partially amorphous during milling. Thermal treatment beyond ~650 K led to the formation of a bcc solid solution, which is characteristic for high entropy alloys.

**Author Contributions:** Conceptualization of the project: A.V.-C., J.J.I., J.S.B. and A.C. Discussion of the results and revision of the paper: A.V.-C., J.J.I., J.S.B. and A.C. Experiments were developed by A.V.-C., J.J.I., and J.S.B.

**Funding:** This research was funded by AEI/FEDER-UE (Project MAT 2016-77265-R) and the PAI of the Regional Government of Andalucía.

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

#### **References**


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