Figure 9 shows the variations of the coercive force (j*H*c) with changing heat treatment (HT) temperature. As can be seen from the data, after low-temperature HT in a range of 475–500 ◦C, j*H*c demonstrated an abrupt increase.

**Figure 9.** The dependence of j*H*c on the heat treatment temperature of Nd–Fe–B-based magnet prepared from the powder mixture with 2 wt.% of Dy3Co0.6Cu0.4Hx.

#### **4. Conclusions**

The phase composition of the Dy3Co0.6Cu0.4 alloy in the initial homogenized and hydrogenated states was studied. The alloy in the homogenized state was multiphase and contained the Dy3(Co,Cu) and Dy(Cu,Co) phases. During the hydrogenation of the alloy, the disproportionation or hydrogenolysis process took place, which, regardless of the multiphase composition of the initial alloy, resulted in the formation of DyH2-3 hydride and a fine (Co + Cu) mixture with small trace quantities of Dy3(Co,Cu) and Dy(Cu,Co).

The study of the sintered Nd(RE)–Fe–B magnet prepared from the strip-cast alloy showed that Dy introduction via the two-component blending method (the hydogenated Dy3Co0.6Cu0.4 compound was added to the powder mixture) resulted in the formation of the core–shell structure of 2–14–1 phase grains. The efficient enhancement of the coercivity of Nd(RE)–Fe–B magnets, with a slight sacrifice of remanence, was demonstrated.

The positive effect of REM-alloy hydrogenated additions to the Nd–Fe–B powder mixture allows the possibility of introducing various components to the permanent magnets (heavy REMs, elements structuring grain boundaries, and restricting the magnet grain growth) at the preparation stage, rather than at the alloy-melting stage. This gives the possibility of using a unified initial alloy for the manufacture of magnets with improved (high-coercive or high-performance) magnetic characteristics.

**Author Contributions:** Writing—original draft, K.S., P.A.P., N.B.K.; writing—review and editing, K.S., P.A.P., N.B.K.; investigation (structure, phase composition, thermal dehydrogenation, magnetic properties), K.S., P.A.P., N.B.K., T.C., T.R.; conceptualization and methodology, A.A.L.; software, Y.S.K.; validation of results of hydrogenation, H.D.; resources, N.A.D.; supervision and project administration, G.S.B.

**Funding:** This study was carried out within the project "Development of physico-chemical and engineering foundations for the initiation of innovative resources-economy technology of high-power and high-coercivity (Nd,r)–Fe–B (R = Pr, Tb, Dy, Ho) low-REM permanent magnets", projects No. LTARF18031 funded by the Ministry of Education, Youth and Sports of the Czech Republic and No. 14.616.21.0093 (unique identification number: RFMEFI61618x0093) funded by the Ministry of Science and Higher Education of the Russian Federation. The SEM/EDS investigation and particle size analysis were performed using the research infrastructure of the Regional Materials Science and Technology Centre, VSB–Technical university of Ostrava (Czech Republic), and

XRPD analysis, DTA/TG analysis, and the study of magnetic characteristics were carried out at the Center of Collaborative Access for Functional Nanomaterials and High-Purity Substances, Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences. TheMasterSizer 3000 particle size analyzer (Malvern) was acquired within the "Development of research and development basis of RMSTC" project (No. CZ.1.05/2.1.00/19.0387) within the frame of the operational program "Research and Development for Innovations" financed by structural funds and the state budget of the Czech Republic.

**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* **Densification of Magnesium Aluminate Spinel Using Manganese and Cobalt Fluoride as Sintering Aids**

#### **Ali Talimian 1, Vaclav Pouchly 2,3, Karel Maca 2,3 and Dusan Galusek 1,4,\***

Received: 3 December 2019; Accepted: 20 December 2019; Published: 24 December 2019

**Abstract:** Highly dense magnesium aluminate spinel bodies are usually fabricated using pressure-assisted methods, such as spark plasma sintering (SPS), in the presence of lithium fluoride as a sintering aid. The present work investigates whether the addition of transition metal fluorides promotes the sintering of MgAl2O4 bodies during SPS. At the same time, such fluorides can act as a source of optically active dopants. A commercial MgAl2O4 was mixed with 0.5 wt% of LiF, MnF2, and CoF2 and, afterwards, consolidated using SPS at 1400 ◦C. Although MnF2 and CoF2 promote the densification as effectively as LiF, they cause significant grain growth.

**Keywords:** MgAl2O4; lithium fluoride; cobalt fluoride; manganese fluoride; spark plasma sintering; grain growth

#### **1. Introduction**

Magnesium aluminate spinel is a material of interest for optical applications due to its excellent mechanical and optical properties. MgAl2O4 has low density (3.58 g cm−3), typical fracture toughness of 1.9 MPa·m0.5, and high optical transmissivity in the visible to mid-infrared ranges [1–7]. Moreover, the spinel structure can host optically active elements, e.g., transition metal ions [8–10]. Having a symmetrical-cubic structure, transparent MgAl2O4 ceramics with high optical homogeneity can be fabricated by removing the scattering centers, such as pores and impurities [7,11–13]. Fabricating highly dense MgAl2O4 is, however, difficult because of the slow diffusion of oxygen. Therefore, spinel is usually densified by two-stage sintering, i.e., pressure-less sintering followed by hot isostatic pressing (post HIPing). Alternatively, spinel can be produced via single-stage pressure-assisted sintering, such as hot pressing (HP) or spark plasma sintering (SPS) [2,14–17]. Using the SPS method makes it possible to fabricate highly dense spinel bodies at a significantly lower temperature and a shorter time as compared with the other methods; this enables suppressing grain growth and producing high-quality samples.

Lithium fluoride (LiF) is a conventional sintering aid in processing MgAl2O4; it promotes the densification by producing transient liquid at low temperatures and introducing cation defects into the spinel structure. Moreover, LiF removes carbon contamination by forming volatile CFx species. [18–22]. However, lithium incorporation into the MgAl2O4 structure can have a detrimental effect on optical properties, especially when spinel is doped with optically active elements [22,23].

Transition metal fluorides that melt at low temperatures are other suitable candidates to be used as additives for sintering of magnesium aluminate spinel. Such dopant provides double benefits. They assist densification through the formation of a transient liquid and, at the same time, introduce

an optically active element into the MgAl2O4 structure. In this study, magnesium aluminate spinel bodies were fabricated by spark plasma sintering of a commercial aluminate spinel powder using LiF, MnF2, and CoF2 as sintering additives. The effect of sintering aids on densification behavior and final microstructure was investigated.

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

A commercial magnesium aluminate spinel powder, Baikalox S30CR (Baikowski, Paris, France) was used as the starting material in this study. The powder is characterized by a BET specific surface area of 26 m2g−1 and a median particle size (d50) of 0.2 μm according to the data provided by the supplier. The spinel powder contains minute amounts of impurities, mainly, S(600), Na(41), and Ca(15) in wt. ppm. Lithium fluoride (LiF), manganese fluoride (MnF2), and cobalt fluoride (CoF2), ACS grade > 99.0, purchased from Sigma-Aldrich (St. Louis, MO, USA) were used as sintering aids.

MgAl2O4 ceramics doped with the sintering aids (0.5 wt%) were prepared by dispersing and mixing of powders in isopropanol, ACS grade >99.0%, using an ultrasonic homogenizer (UW2200, BANDELIN, Berlin, Germany). Then, the mixtures were transferred to a rotary evaporator and dried. Ready-to-press (RTP) powder was prepared by passing the dried mixture through a sieve with a screen mesh of 500 μm.

Samples were consolidated using a spark plasma sintering machine (Dr. SINTER SPS-625, FUJI, Tokyo, Japan). The RTP powder was filled in a graphite die with an inner diameter of ca. 12 mm.  The die was then wrapped in a carbon felt and placed between the moving rams of the SPS. The sintering schedules consisted of fast increases of the temperature to 600 ◦C in 3 min followed by heating of the sample at a constant heating rate of 100 ◦C min−1 to 1400 ◦C at which the shrinkage stops; therefore, sintering processes were carried out with no dwelling time to avoid unnecessary grain growth. The sintering was carried out under vacuum (5 to 9 Pa). A constant uniaxial pressure of 75 MPa was applied above 800 ◦C.

The displacement of punches and the temperature were recorded during the whole heating/cooling step. The pellets' temperature was measured constantly by using an optical pyrometer focused on the hole drilled into the die wall. The coefficient of thermal expansion (CTE) of the system (e.g., graphite die, paper, and punches) was determined separately (i.e., in a run without the specimen) throughout the temperature range of this study (600 to 1400 ◦C) in order to account for the instrumental error. The sintered pellets were subsequently subjected to a heat treatment at 800 ◦C (heating rate 2.5 ◦C min−1) for 60 min in air in a muffle furnace to remove the residual carbon from the surfaces.

The bulk density and apparent porosity of sintered bodies were measured using Archimedes' method in deionized water according to the ASTM standard (C329-88(2016)) [24]. All provided values are the means of at least 10 independent measurements.

The melting temperature of sintering aids and their reactions with the spinel powder were studied using thermal analysis. The measurements were performed by a simultaneous thermal analyzer (STA 449 F1 Jupiter, Netzsch, Selb, Germany) in Differential Thermal Analysis, DTA, configuration, using alumina crucibles in flowing N2 (20 mL min−1). Thermal Gravimetric analysis, TG, was performed simultaneously. Data were collected on ca. 100 mg of mixtures containing 10 wt% of sintering aids upon heating at a constant rate of 20 ◦C min−1 to 1350 ◦C.

The samples' microstructure was examined using a scanning electron microscope, SEM, (JEOL 7600F, JEOL, Tokyo, Japan) equipped with an energy dispersive X-ray spectrometer (EDXS, Oxford Instruments, Abingdon, UK). Small fragments were collected from the fractured surface of samples and fixed on aluminium sample holders using conductive adhesive tape and coated with carbon to prevent charging.

#### **3. Results**

Figure 1 shows the SEM micrograph of the magnesium aluminate spinel powder; the powder consists of submicron agglomerates comprised of smaller nanoparticles with a median diameter of 90 ± 15 nm. However, the specific surface area indicates a somewhat smaller primary particle size of approximately 64 nm. Similarly, Maca et al. examined the primary particle size of the same commercial MgAl2O4 powder, and reported an average particle size of 58 nm, by assuming that the primary particles have a spherical shape. The median particle size provided by the producer is, therefore, related to the size of agglomerates [25].

**Figure 1.** Scanning electron microscope (SEM) image of magnesium aluminate spinel powder.

Table 1 summarizes the measured density and porosity of samples produced from the powder mixture containing 0.5 wt% of additives and the additive-free sample; the theoretical density of samples was calculated using the density of magnesium aluminate spinel (3.58 g cm−1) and the density of a respective additive following the rule of mixtures. The density of LiF, MnF2, and CoF2 are 2.64, 3.98, and 2.70 g cm−3, respectively. The measured density of an additive containing samples is within the range of experimental error comparable to the density of additive-free samples. While the residual porosity of additive-free samples is almost zero, the doped samples are characterized by limited amounts of closed porosity. Such behavior can be related to the evaporation of additives at high temperatures.

**Table 1.** Relative density and apparent porosity of additive-free and transition metal fluoride-doped samples produced by spark plasma sintering (SPS) at 1400 ◦C (no isothermal dwell). The numbers in parenthesis represent standard errors.

Figure 2 shows the temperature dependence of shrinkage and shrinkage rates of the powder mixtures containing 0.5 wt% of the additives and of the additive-free spinel powder during SPS; the shrinkage was determined by measuring the punch displacement upon heating at the constant heating rate of 100 ◦C min−1, between 600 ◦C and 1400 ◦C. The shrinkage rate was calculated point-by-point, using Equation (1):

$$\frac{d(\frac{\Delta l}{l\_0})}{dt} = \frac{1}{l\_0} \frac{\Delta l\_{T+\delta T} - \Delta l\_{T-\delta T}}{2\delta T} \dot{T} \tag{1}$$

where Δ*l* represents linear shrinkage measured at the temperature *T*, *l*o is the original length of the sample, *t* represents time, and the variable . *T* stands for heating rate. The shrinkage curves are characterized by two main regions, a rapid shrinkage of ~5%, occurring around 800 ◦C, followed by continuous shrinkage up to ~23%, after which the curve reaches a plateau. While the latter is related to the densification by sintering, the former is attributed to the powder particles' rearrangement when pressure was applied [26]. The densification of all samples starts at around 850 ◦C. The sintering aids clearly decrease the temperature at which the densification is completed. The shrinkage curve of additive-free spinel reaches a plateau indicating the end of shrinkage, at 1350 ◦C, whereas the shrinkage of samples doped with LiF, CoF, and MnF2 stops at 1170, 1195, and 1250 ◦C, respectively. Moreover, the shrinkage rate of doped samples is significantly higher than that of the pure spinel, particularly at temperatures higher than 1000 ◦C (Figure 2b).

**Figure 2.** (**a**) Relative shrinkage of additive-free spinel and spinels doped with LiF, MnF2, and CoF2 (0.5 wt%) with temperature and (**b**) first derivative of the shrinkage calculated using Equation (1).

Figure 3 summarizes the results of DTA and TG analyses of powder mixtures containing 10 wt% of fluorides in the temperature interval between 600 to 1350 ◦C. These were carried out as reference measurements elucidating thermal processes on fluoride doped powders. The DTA curve of LiF-doped samples is characterized by a sharp endothermic peak at ~830◦ attributed to chemical reactions and melting of lithium fluoride, as discussed below (Equations (6) and (7)). In contrast, the sample

containing MnF2 exhibits no clear endothermic effect at the melting temperature of MnF2 (856 ◦C). The behavior of the CoF2 containing sample is similar, showing no thermal effect, which could be attributed to melting of CoF2 (i.e., at ca 930 ◦C). However, all samples exhibit an endothermic peak at 1240 ◦C, attributed to the eutectic melting of magnesium fluoride, indicating chemical reactions between sintering aids and MgAl2O4, yielding MgF2, as pointed out in the following text. The TG curves show that the weight of LiF samples decreases rapidly above 1050 ◦C, while samples containing MnF2 and CoF2 exhibit slower weight loss in the following two steps: a slow decline above 850 ◦C followed by a rapid decrease over 1050 ◦C. The onset of weight loss can be correlated with the small endothermic effect on DTS curves associated with melting of MgF2 (1263 ◦C). The observed weight loss was then associated with evaporation of MgF2 from the melt. On the basis of the literature data, vapor pressure of molten MgF2 reaches ~13 Pa at 1270 ◦C and ~130 Pa at 1434 ◦C, so its loss is expected to be significant.

**Figure 3.** DTA (**a**) and TG (**b**) records of samples comprising 10 wt% of LiF, MnF2, and CoF2, heating rate = 20 ◦C min−1. Vertical line at 1230 ◦C represents onset of melting for TM fluorides doped spinel ceramics.

Figure 4 shows the X-ray diffraction pattern of pellets produced by SPS at 1400 ◦C; the XRD pattern of as-received spinel powder is also shown for comparison. According to the XRD experiments, magnesium aluminate spinel is the only crystalline phase present in the samples; the sintered samples are characterized by sharp and narrow diffraction maxima that imply the sintering procedure (heating up to 1400 ◦C with the heating rate of 100 ◦C min−1, with no dwell time) increases the size of coherently diffracting domains (crystallites). The XRD patterns were analyzed further by using Rietvel refinement [27,28]. The lattice parameter of additive-free spinel is estimated to be 8.0798 ± 0.0002 Å while the lattice parameter for the samples doped with LiF, MnF2, and CoF2 are 8.0814 ± 0.0001 Å, 8.0833 ± 0.0003 Å, and 8.0833 ± 0.0001 Å, respectively. The incorporation of dopants into the spinel structure results in a slight increase of the lattice parameter due to size mismatch of doping cations and Mg2+ and Al3+ in the spinel crystal lattice.

**Figure 4.** XRD patterns of additive-free and doped samples with 0.5 wt% (LiF, MnF2, and CoF2), spark plasma sintered at 1400 ◦C. The diffraction pattern of additive-free powder is also shown for comparison. The experimental data are fitted by the model patterns obtained by Rietveld refinement of experimental data.

Figure 5 shows the fracture surface of additive-free and doped samples with the 0.5 wt% of LiF, MnF2, and CoF2. All doped samples exhibit significant grain coarsening. The LiF-doped spinel is, interestingly, characterized by a smaller grain size as compared with the spinel doped with MnF2 and CoF2.

Bright spots observed on fracture surfaces were studied by EDX. Figure 6 shows a typical EDX spectrum collected from a bright spot in a CoF2 doped sample. The spots also contain, along with doping ions, a significant concentration of Sulphur, implying that the sulphate impurities in the spinel powder reacted with the dopant yielding sulfate phases during sintering. However, the content of sulphates was below the detection limit of X-ray diffraction, and the size of sulphate inclusions too small to be identified by EBSD.

**Figure 5.** Scanning electron micrographs of the fracture surfaces after spark plasma sintering at 1400 ◦C. The images were recorded in a BSE mode from (**a**) additive-free sample, and the samples doped with 0.5 wt% of (**b**) LiF, (**c**) MnF, and (**d**) CoF2.

**Figure 6.** (**a**) Scanning electron micrograph showing in detail bright spots in CoF2 doped spinel after spark plasma sintering and (**b**) the EDX spectrum of the selected areas/spots: (I), spinel grain and (II), bright spot.

#### **4. Discussion**

Figure 2 shows that the onset of densification of all studied samples occurs at a lower temperature than in conventional sintering. The application of pressure during spark plasma influences densification in two ways. First, powder particles rearrange under pressure. Secondly, the densification mechanism is also affected, due to grain boundary sliding [24,27]. Consequently, the maximum in densification rate of SPS samples is achieved at a lower temperature as compared with that reported for conventional sintering of the same powder (1350 ◦C, S30CR, Baikowski) [29,30].

The LiF-doping results in a higher densification rate, and larger grains than in the additive-free spinel are formed. The interaction between LiF and MgAl2O4 above the melting LiF point, 840 ◦C, and the formation of liquid are described as follows [31,32]:

$$\text{LiF}(l) + \text{MgAl}\_2\text{O}\_4(\text{s}) \rightarrow \text{LiF} : \text{MgF}\_2(l) + \text{LiAlO}\_2(\text{s}) \tag{2}$$

The transient liquid enhances the densification through two main mechanisms, liquid redistribution facilitates particle rearrangement and the fluorine rich liquid enhances the mass transport. Moreover, lithium aluminate spinel can produce solid solution with magnesium aluminate spinel and introduces structural defects according to:

$$\text{LiAlO}\_2 \xrightarrow{\text{MgAl}\_2\text{O}\_4} \text{Li}\_{\text{Mg}}^{\prime} + V\_{\text{Al}}^{\prime // \prime} + \text{Al}\_{\text{Al}}^{\times} + 2\text{O}\_{\text{O}}^{\times} + 2V\_{\text{O}}^{\times} \tag{3}$$

Introduction of structural defects, such as oxygen vacancies, facilitates the movement of ions, particularly oxygen ions and, in turn, promotes the densification, as well as the grain growth. With the further increase of temperature, the evaporation rate of the transient liquid accelerates (above ca. 1100 ◦C, Figure 3b) and the liquid phase is effectively removed from the system. Consequently, the densification rate decreases.

$$\text{LiF}: \text{MgF}\_2(l) \rightarrow \text{LiF}(\text{g}) + \text{MgF}\_2(\text{g})\tag{4}$$

Finally, gaseous MgF2 reacts with LiAlO2 forming spinel again:

$$2LiAlO\_2(s) + MgF\_2(g) \to 2LiF(g) + MgAl\_2O\_4(s) \tag{5}$$

Considering densification behavior of MnF2 and CoF2 similar to that of the LiF doped samples, a similar sintering mechanism can also be expected for the transition metal fluorides. Surprisingly, although CoF2 and MnF2 have higher melting temperatures than LiF, the CoF2 and MnF2 doped spinel exhibits more extensive grain coarsening than the LiF doped ones.

Interestingly, the DTA records show no endothermic effect, which would indicate melting of pure transition metal fluoride additives around the expected temperatures of their melting. Only a small endothermic effect corresponding to melting of MgF2 implies chemical reactions between MnF2 or CoF2 and MgAl2O4 yielding transition phases. Due to the similar ionic radius of magnesium, manganese and cobalt (*rMg*2+ = 72 pm, *rMn*2+ = 70 pm, and *rCo*2+ = 75 pm), it can be assumed that the transition metal ions replace magnesium ions in the spinel crystal lattice, producing MgF2:

$$\rm MnF\_2(s) + MgAl\_2O\_4(s) \to MnAl\_2O\_4(s) + MgF\_2(s) \tag{6}$$

$$\text{CaF}\_2(\text{s}) + \text{MgAl}\_2\text{O}\_4(\text{s}) \rightarrow \text{CaAl}\_2\text{O}\_4(\text{s}) + \text{MgF}\_2(\text{s})\tag{7}$$

The formation of a solid-solution with different divalent cations within the spinel structure results in spinel structure strain, as well as introduction of point defects, such as oxygen vacancies as a result of hosting divalent ions in octahedral sites. [7,33]. The TG results, Figure 3b, confirm that the weight loss of MnF2 and CoF2 doped samples begins at higher temperatures than in the LiF doped material (1250 ◦C vs. 1100 ◦C). The transient liquid is, therefore, present at grain boundaries for a longer time, providing a faster diffusion path for the elements, and resulting not only in efficient densification but also in more pronounced grain growth. This resulted in the increase of the median grain size from 0.8 μm in undoped ceramic to 10.3 μm, 14.0 μm, and 11.6 μm in LiF, MnF2, and CoF2 doped spinel, respectively, as shown in Figure 5. Apart from the presence of transient liquid in spinels doped by transition metal fluorides, the finer grain size of LiF doped samples can be attributed also to Zener pinning effect of LiAlO2 (Equation (1)) precipitated at grain boundaries [34].

Transition metal fluorides act as sintering aid during the densification of magnesium aluminate spinel and produce spinel structures containing optically active ions, e.g., Mn2+ and Co2+, that can be used in applications such as white LEDs or Q-switches. [35,36] Further studies are required to evaluate whether and how the addition of transition metal fluorides affects the densification and final properties of magnesium aluminate spinel ceramics.

#### **5. Conclusions**

Highly dense magnesium aluminate spinel bodies doped with LiF, MnF2, and CoF2 were produced using spark plasma sintering. Although the contribution of CoF2 and MnF2 to the densification of MgAl2O4 is more complicated as compared with LiF, they promote the densification almost as efficiently as LiF, despite the higher melting points of transition metal fluorides. The MnF2 and CoF2 containing samples exhibit larger grains as compared with LiF-doped spinel spark plasma sintered under the same conditions.

**Author Contributions:** Conceptualization, D.G. and A.T.; methodology, A.T. and V.P.; formal analysis, A.T., V.P., and D.G.; investigation, A.T. and V.P.; resources, D.G. and K.M.; data curation, A.T. and V.P.; writing—original draft preparation, A.T. and V.P.; writing—review and editing, D.G. and K.M.; visualization, A.T. and V.P.; supervision, D.G. and K.M.; project administration, D.G.; funding acquisition, K.M and D.G. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the European Union's Horizon 2020 research and innovation programme, grant agreement no. 739566. Financial support of this work by the grants SAS-MOST JRP 2015/6 and VEGA 2/0026/17 is gratefully acknowledged. We appreciate the support provided by the Czech Ministry of Education under grants LTT18013 (Inter-Transfer) and CEITEC 2020 (LQ1601).

**Acknowledgments:** This paper is a part of the dissemination activities of the project FunGlass.

**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*

## **Rapidly Solidified Aluminium Alloy Composite with Nickel Prepared by Powder Metallurgy: Microstructure and Self-Healing Behaviour**

#### **Alena Michalcová \*, Anna Knaislová, Jiˇrí Kubásek, Zden ˇek Kaˇcenka and Pavel Novák**

Department of Metals an Corrosion Engineering, Faculty of Chemical Technology, University of Chemistry and Technology, Prague 16628, Czech Republic<EMAIL_ADDRESS>(A.K<EMAIL_ADDRESS>(J.K<EMAIL_ADDRESS>(Z.K<EMAIL_ADDRESS>(P.N.)

**\*** Correspondence<EMAIL_ADDRESS>
Received: 13 November 2019; Accepted: 11 December 2019; Published: 13 December 2019

**Abstract:** Composite material prepared by spark plasma sintering (SPS) from a powder mixture of AlCrFeSi rapidly solidified alloy and 5 wt. % of Ni particles was studied in this work. It was proven that during SPS compaction at 500 ◦C, no intermetallic phases formed on the surface of Ni particles. The material exhibited sufficient mechanical properties obtained by tensile testing (ultimate tensile stress of 203 ± 4 MPa, ductility of 0.8% and 0.2% offset yield strength of 156 ± 2 MPa). Tensile samples were pre-stressed to 180 MPa and annealed at 450 and 550 ◦C for 1 h. Annealing at 450 ◦C did not lead to any recovery of the material. Annealing at 550 ◦C caused the full recovery of 0.2% offset yield strength, while the ductility was decreased. The self-healing behaviour originates from the growth of intermetallic phases between the Ni particle and the Al matrix. The sequence of NiAl, Ni2Al3 and NiAl3 intermetallic phases formation was observed. In particular, the morphology of the NiAl3 phase, growing in thin dendrites into the Al matrix, is suitable for the closing of cracks, which pass through the material.

**Keywords:** self-healing; aluminium alloy; microstructure

#### **1. Introduction**

Self-healing materials are designed in order to enrich structural materials by the ability of closing crack that was formed during material utilization. Closing of the crack will lead to the restoration of the mechanical cohesion of the material. This may decrease the impacts of an accident caused by material failure.

Several mechanisms of the self-healing behaviours have been proposed. Precipitation mechanism at low temperature was observed in an Fe-Au alloy [1]. In several other alloys, precipitation self-healing of the crack occurred after an exposure to elevated temperature [2–5]. Phase transformation may also be a driving force for the self-healing behaviour. Composite materials with the shape memory alloy (SMA) reinforcement and the matrix with off-eutectic composition were described in [6,7]. In this case, the phase transformation of the SMA (smart material, in this case shape memory alloy) reinforcement leads to the crack closing and a partial matrix melting heals the already closed crack. Another type of the self-healing behaviour leading to crack closing was observed in Al-Ag alloy, where the phase transformation was between the Ag2Al phase and the fcc-Al [8]. A different strategy for the crack healing is a usage of an encapsulated healing agent. A low melting metal or alloy is filled into a hollow diffusion barrier, typically composed of oxides [5]. The idea is that the spreading crack will break the diffusion barrier and enter the area with the low melting healing agent. After annealing, the low melting healing agent will fill the crack and the hollow diffusion barrier remains as a spherical pore in the microstructure.

The idea presented in this paper is to modify the encapsulated-healing-agent strategy by adding Ni powder into rapidly solidified Al alloy. The diffusion barrier will be formed by nickel aluminides and may provide the material with a better cohesion, compared to the oxides. During annealing, intermetallic phases will be formed by a reaction of Ni and Al. This process is accompanied by volume changes [9,10], so a decreased level of porosity may be expected compared to the self-healing by the encapsulated low-melting metal/alloy mechanism. By this innovative process, self-healing materials will be prepared by the two-step powder metallurgy processing (1) preparation of rapidly solidified alloy, (2) compaction. No further steps are necessary to supply self-healing agent to the material. Simple powder mixing before compaction is required.

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

Powder alloy with intended composition AlCr6Fe2Si1 (given in wt. %) was prepared by a gas atomisation and mixed with 5 wt. % of a gas atomized Ni. In both cases, the finest granulometric fraction of the size less than 63 μm was used. The powder mixture was consolidated using the spark plasma sintering (SPS) device (FCT Systeme HP D 10, Rauenstein, Germany) forming a bulk sample of 40 mm in diameter and approximately 13 mm in height. Chemical composition measured by X-Ray fluorescence spectroscopy (XRF, ARL 9400 XP, Thermo ARL, Switzerland) of the sintered sample denoted as "AlCrFeSi + Ni" is given in Table 1.

**Table 1.** Chemical composition of AlCrFeSi + Ni.

Bulk material was observed by scanning electron microscopy (SEM, TESCAN VEGA 3 equipped by energy dispersive spectroscopy (EDS) detector by Oxford Instruments, Abingdon, UK). SEM samples were prepared by mechanical grinding (P320-P4000), by polishing using diamond paste (D3, D0.7) and by etching in 0.5 % HF solution. TEM samples were prepared by ion polishing using a Gatan PIPs device (Gatan, Pleasanton, CA, USA). Observation by the transmission electron microscopy (TEM) was done by Jeol 2200 FS equipped by EDS detector by Oxford Instruments. Tensile testing was performed by the Instron 5882 universal loading machine with the strain rate of 1 mm / min. Dog-bone shaped samples with the length of 16 mm, width of 4 mm, and height of 2 mm with 20 mm radius were used. Phase composition was determined by XRD using diffractometer PANanalytical X´Pert PRO, Co lamp (Almelo, The Netherlands).

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

#### *3.1. Microstructure of As-Sintered AlCrFeSi* + *Ni Material*

Bulk sintered material was composed of Al matrix, Al13Cr2 and Al13Fe4 intermetallic phases and Ni as proven by XRD pattern, shown in Figure 1.

Figure 2 shows the SEM micrograph of the as-sintered AlCrFeSi + Ni composite material. In the Al matrix, fine particles of intermetallic phases Al13Cr2 (round) and Al14Fe4 (needles) are visible. The initial particles of Al alloy were deformed during the SPS process, while Ni particles remained spherical.

**Figure 1.** X-Ray diffraction pattern of as-sintered AlCrFeSi + Ni.

**Figure 2.** Micrograph of as-sintered AlCrFeSi + Ni (SEM/BSE).

To describe the boundary between the Al matrix and the Ni particle, it is necessary to use the TEM. The TEM micrograph is given in Figure 3. The light grey Al matrix contains fine intermetallic particles. The internal structure of the Ni particle is dendritic.

**Figure 3.** Detailed micrograph of as-sintered AlCrFeSi + Ni (TEM).

More detailed view was obtained in STEM mode (Figure 4). Sufficient contact of the Al matrix with the Ni particle was observed, but the connection was not uniform. EDS line profiles were measured in a location with the good connection (Line 1) and in a location where the gap is evident (Line 2).

**Figure 4.** Detailed micrograph of boundary between Al matrix and Ni particle in the as-sintered AlCrFeSi + Ni (STEM) with indication of line profiles.

The EDS profile of the Line 1 is plotted in Figure 5. Although it is the location with a good connection, no plateau in any profile was observed. This means that no intermetallic phase with a stoichiometric composition was formed. Enrichment of the Ni particle surface by Cr was observed. This impurity might originate from the initial Ni powder.

**Figure 5.** Line profile 1, location shown in Figure 4.

The EDS profile of the Line 2 is plotted in Figure 6. At the position of 0.3 μm from the beginning of the line, the minimum of all the elements' profile lines can be seen. This is the position of the gap. No plateau was observed again in the vicinity of the gap. This proves that the sintering process caused only partial connection of Al matrix and Ni particles. The other possible explanation of the presence of a gap would be a formation of the gap caused by different thermal expansion of intermetallic interlayer between Al matrix and Ni particle during cooling of sintered samples. As the composition changes gradually (Figure 5), there is no evidence of phase transformation during the SPS process.

**Figure 6.** Line profile 2, location shown in Figure 3.

#### *3.2. Microstructure of AlCrFeSi* + *Ni Material Annealed at 550* ◦*C for 1 h*

Figure 7 illustrates microstructure of AlCrFeSi + Ni composite after annealing at 550 ◦C for 1 h. The composite behaved like diffusion couple that was already well described in several works [10–16]. The Ni residual core is covered by a thin layer of NiAl phase. A large layer of Ni2Al3 is on the top of the NiAl layer. Dendrites of NiAl3 grow into Al matrix. Porosity caused by diffusion (Kirkendall porosity) can be observed between NiAl3 dendrites. The intermetallic particle has an irregular shape compared to the initial Ni sphere. This observation verifies the volume change during formation of intermetallic phases, which is a sign of the self-healing behaviour potential of the AlCrFeSi + Ni material.

**Figure 7.** Micrograph of AlCrFeSi + Ni annealed at 550 ◦C for 1 h (SEM/BSE).

The phase composition was determined by point EDS analysis. Actual positions of each measured spectrum are shown in Figure 8. Theoretical content of Al in NiAl is 50 at. %, in Ni2Al3 60 at. % and in NiAl3 75 at. %. The values of Al content presented in Table 2 are slightly higher that the teoretical ones. It is due to interaction volume of the EDS point spectra measurement (several cubic micrometers). It is highly probable that the EDS spectra contain a partial signal from the Al matrix, which increases the measured values.

**Figure 8.** Micrograph of AlCrFeSi + Ni annealed at 550 ◦C for 1 h with labelled EDS point spectra.

**Table 2.** Chemical composition of points labelled in Figure 8 given in at. % (EDS).

#### *3.3. Tensile Testing*

Figure 9 plots the results of tensile testing of AlCrFeSi + Ni composite material. The black curve corresponds to a testing of the material in the as-sintered state. The results give the essential description of behaviour of the material. The ultimate tensile stress was 203 ± 4 MPa, the ductility was 0.8 %. The yield stress cannot be estimated clearly, so the 0.2% offset yield strength was measured and it reaches the value of 156 ± 2 MPa. Subsequently, the samples were pre-stressed to 180 MPa, where the tensile test was stopped. The pre-stressed samples were annealed for 1 h at temperatures of 450 ◦C and 550 ◦C. The pre-stressed and annealed samples underwent the tensile testing under the same conditions as the as-sintered AlCrFeSi + Ni material. Annealing at 450 ◦C caused an increase in ductility, but the 0.2% offset yield strength was lower. This means that the most important property form the structural point of view was not preserved. The material annealed at 550 ◦C had a lower ductility, but the 0.2% offset yield strength was exactly the same as the one of the as-sintered material.

For the self-healing materials with SMA reinforcement, is was described that the material after healing reached 95% of UTS value and the ductility decrease from 6.4% to 2.2% (corresponding to 34% of the original ductility value) [6]. In our case, 63% of UTS value was achieved but what is more important by preserving the value of 0.2% offset yield strength. The value of ductility decreased in our case from 0.8% to 0.27%. This means that the self-healed material had ductility reaching 33% of the value of sintered material, which is in good agreement with other materials with self-healing properties [6].

**Figure 9.** Tensile testing of AlCrFeSi + Ni, black—tensile tested in as-sintered state, green—prestressed up to 180 MPa, annealed at 450 ◦C / 1 h and tensile tested, red—prestressed up to 180 MPa, annealed at 550 ◦C/1 h and tensile tested.

Figure 10 shows the SEM micrographs of fracture surfaces of the tensile tested materials. The Ni particle visible in the as-sintered material (a) has a dendritic surface morphology. The SPS process was successful, as the Ni particle is an integral part of the material and does not serve as stress concentrator (the Al matrix is still surrounding the Ni particle). The morphology of the Ni particle after annealing at 450 ◦C was not changed (b) and the whole fracture surface looks similar to the as-sintered material. The fracture surface of the material annealed at 550 ◦C is significantly different. The fracture surface of Ni2Al3 particle is shown in the middle (c) with small round island of NiAl. Dendrites of NiAl3 are interconnecting Ni2Al3 particle with Al matrix in the vicinity. In the areas where initial Ni particles were close (top left corner), the NiAl3 are not observed. In these places, polyhedrons of Ni2Al3 are formed. This may be the reason for the decrease of the ductility of the material annealed at 550 ◦C.

**Figure 10.** Fractography of AlCrFeSi + Ni, (**a**) tensile tested in as-sintered state, (**b**) prestressed up to 180 MPa, annealed at 450 ◦C / 1 h and tensile tested, (**c**) prestressed up to 180 MPa, annealed at 550 ◦C / 1 h and tensile tested (SEM).

#### **4. Conclusions**

AlCrFeSi + Ni composite material was successfully prepared by the SPS process. During sintering, no intermetallic phase was formed on the surface of Ni particles, as proved by TEM observation. The results of tensile testing showed sufficient mechanical properties of the AlCrFeSi + Ni material (ultimate tensile stress: 203 ± 4 MPa, ductility: 0.8%, 0.2% offset yield strength: 156 ± 2 MPa).

The tensile testing samples were pre-stressed to 180 MPa (the value above 0.2% offset yield strength) and annealed. On the annealed samples, the tensile testing was performed again. The sample annealed at 450 ◦C, which is under the temperature of intermetallic formation, exhibits a slightly higher ductility, but the 0.2% offset yield strength did not reach the value of the as-sintered material. The value of 0.2% offset yield strength for the sample annealed at 550 ◦C was similar to the value of the as-sintered material. The self-healing behaviour was enabled by the formation of a sequence of nickel aluminides from initial Ni particle into Al matrix. The formation of Ni2Al3 polyhedrons caused the decrease of materials ductility. This problem might be solved by an optimization of the Ni content in the material.

**Author Contributions:** A.M. wrote the paper and performed TEM measurements. A.K. performed SEM measurements. J.K. performed tensile testing. Z.K. prepared samples for TEM. P.N. analysed the data connected with intermetallic growth.

**Funding:** This research was financially supported by the Czech Science Foundation, project No. GJ17-25618Y. The authors acknowledge the assistance in TEM sample preparation provided by the Research Infrastructure NanoEnviCz, supported by the Ministry of Education, Youth, and Sports of the Czech Republic under Project No. LM2015073.

**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* **Zn-Mg Biodegradable Composite: Novel Material with Tailored Mechanical and Corrosion Properties**

#### **Jiˇrí Kubásek 1,\*, Drahomír Dvorský 1, Jaroslav Capek ˇ 2, Jan Pinc 2 and Dalibor Vojt ˇech 1**

Received: 7 November 2019; Accepted: 24 November 2019; Published: 27 November 2019

**Abstract:** Zinc-based alloys represent one of the most highly developed areas regarding biodegradable materials. Despite this, some general deficiencies such as cytotoxicity and poor mechanical properties (especially elongation), are not properly solved. In this work, a Zn-5Mg (5 wt.% Mg) composite material with tailored mechanical and superior corrosion properties is prepared by powder metallurgy techniques. Pure Zn and Mg are mixed and subsequently compacted by extrusion at 200 ◦C and an extrusion ratio of 10. The final product possesses appropriate mechanical properties (tensile yield strength = 148 MPa, ultimate tensile strength = 183 MPa, and elongation = 16%) and decreased by four times the release of Zn in the initial stage of degradation compared to pure Zn, which can highly decrease cytotoxicity effects and therefore positively affect the initial stage of the healing process.

**Keywords:** biomaterials; metallic composites; powder technology; zinc

#### **1. Introduction**

Binary zinc–magnesium alloys have been studied with regard to the improvement of corrosion protection of steel [1–3], and the development of arterial stents [4–10] and various biodegradable fixation devices [11,12]. As a relatively new biodegradable material, they can compete with magnesium products [7–10]. Although zinc (ρ = 7.14 g·cm−3) is much heavier than Mg (ρ = 1.74 g·cm−3), it is less susceptible to corrosion in the human environment. Various Zn-based materials with superior mechanical and corrosion properties have been developed and studied [9,10,13,14] but pure Zn (even in low doses) is relatively toxic to cells [15] compared to Mg. The daily zinc requirement for an adult is estimated to be 15 mg/day [7,16]. On the contrary, even up to 700 mg/day of Mg can be tolerated [7,16]. The degradation of Zn-based materials is connected to the production of Zn2+ ions, which are released to the surrounding extracellular space and spread into different parts of the organism. Zn2+ ions play an important role in various cellular processes (cell proliferation, differentiation, and signaling). Among other things, a high concentration of Zn2+ can induce apoptosis or necrosis or destroy ion-dependent intracellular signaling pathways [14]. On the contrary, zinc ions in adequate doses can enhance regulation of genes, cell survival/growth and differentiation, extracellular matrix (ECM) mineralization, and osteogenesis [17]. Magnesium is well tolerated by the organism in larger doses than zinc and affects activation of many enzymes, co-regulation of protein synthesis and muscle contraction, and stabilization of DNA and RNA [7]. Although the results of in vitro cytotoxicity tests are generally inconsistent [10,15,18] and long term in vivo tests do possess some good results [10], the lower initial release of Zn from the degradable implant is desirable to improve the osseointegration process during the initial stage of healing and suppress the possible formation of necrotic tissue. Unfortunately, alloying generally leads to a minor decrease in the corrosion rate of Zn-based alloys [10,14]. In this

work, we selected a non-traditional way of producing a Zn-5Mg composite material by powder metallurgical methods. Although powder metallurgy is a powerful method for the production of materials with superior mechanical and corrosion properties, it has been very rarely used to prepare biodegradable Zn-based materials intended for medical devices [19,20]. However, our presented results demonstrate that this technique enables the preparation of materials with tailored mechanical and corrosion properties and improved biocompatibility.

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

#### *2.1. Sample Preparation*

Powders of pure Mg in the form of spherical particles 50–200 μm in diameter and pure Zn in the form of elongated irregular particles about 30–50 μm in thickness and 60–300 μm in length were mixed in Turbula® T2F (WAB-GROUP, Muttenz, Switzerland) for 20 min in a ratio equal to the composition of Zn-5Mg (5 wt.% Mg). The mixture was subsequently compressed in a mold 20 mm in diameter by loading 80 kN for 2 min. Green compacts were extruded at 200 ◦C and an extrusion velocity of 5 mm/min. The final diameter of the extruded rod was 6.5 mm, which corresponds to an extrusion ratio equal to 10. All processing steps were performed under an air atmosphere. Pure Zn was prepared in the same way as the reference material.

#### *2.2. Microstructure Characterization*

Mechanical grinding and polishing were used for surface pre-treatment. Final polishing was performed using an Eposil F suspension (Metalco Testing s.r.o, Roztoky u Prahy, Czech Republic). The microstructure was characterized by optical microscopy and scanning electron microscopy (TescanVEGA3, TESCAN Brno, s.r.o, Brno, Czech Republic) equipped with Energy-Dispersive X-Ray Spectroscopy-EDS (AZtec, Oxford Instruments, Abingdon, United Kingdom).

#### *2.3. Mechanical Properties*

Vickers Hardness HV1 was measured on sample planes perpendicular to the extrusion direction. Compression tests were measured on cylindrical samples 5 mm in diameter and 7 mm in height. Classical dog bone specimens with gauge length equal to 10 mm and diameter equal to 4.5 mm were used for tensile tests. Shoulders on the side of specimens were 15 mm in length and 5.5 mm in diameter. The radius between the shoulders and loaded part was 3.5 mm. Both compressive and tensile tests were performed using a LabTest 5.250SP1-VM universal loading machine (LABORTECH s.r.o., Opava, Czech Republic) at a strain rate of 0.001 s−1. Loading was performed parallel to the extrusion direction. Tensile yield strength (TYS), ultimate tensile strength (UTS), and elongation to fracture were evaluated from engineering stress–strain curves.

#### *2.4. Corrosion Properties*

Samples 5 mm in diameter and 7 mm in height were immersed in simulated body fluid (SBF27) prepared according to Müller at al. [11] for 336 h. The ratio between sample surface and volume of solution was 100 mL·cm−2. After immersion testing corrosion products were removed from the samples by repeated immersion of the samples in a solution containing 200 g·L−1 CrO3. This process was performed until the change of the sample weight between steps was less than 0.001 g. Corrosion rates were calculated from the weight changes between the initial sample and dry sample with removed corrosion products and the concentration of released Zn ions into the corrosion media was determined by ICP-MS analyses (Elan DRC-e, Perkin-Elmer, Waltham, MA, USA). Cathodic curves with a general three electrode setup (sample = working electrode, glassy carbon = counter electrode, and Ag/AgCl containing saturated KCl = reference electrode) were measured using Parstat 3000-MC (AMETEK—Measurement, Communications & Testing, Berwyn, Pennsylvania, USA) in SBF. A sample with the same dimensions as for immersion tests mounted through the screw-thread to the Teflon holder

was left in the corrosion media for 60 min to stabilize the open circuit potential (OCP); subsequently, polarization started from +0.02 versus OCP and continued to −0.6 versus OCP with a rate of 1 mV·s−1. Before all corrosion tests, samples were ground on SiC papers of up to P2000 and degreased in ethanol.

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

#### *3.1. Microstructure Characterization*

The microstructure of the extruded materials is depicted in Figure 1. Both Zn (Figure 1a) and Zn-5Mg (Figure 1b,c) can be seen to be characterized by a partially recrystallized bimodal microstructure containing equiaxed grains with grain sizes in the range 10–30 μm and elongated grains 10 μm in thickness and 50–300 μm in length. Oxide shells, which come from the original powder surface, can be observed to surround the deformed particles. No systematic difference between the microstructure of the Zn matrix was observed for pure Zn and Zn-5Mg extruded materials. Magnesium particles (black areas in Figure 1) can be seen to be homogeneously distributed in the Zn matrix. The interface between Zn and magnesium is occupied by an Mg2Zn11 intermetallic phase [21] (Figure 2) which belongs to the gamma brasses type of phase and contains 15.5 wt. % of Mg. This thermodynamically stable intermetallic phase [22] is generally formed during solidification of Zn–Mg alloys as part of the eutectic reaction. The average thickness of the presented phase was 1.2 ± 0.2 μm. Yang et. al. [23] have also studied Zn–Mg composite materials prepared from pure powders by spark plasma sintering (SPS). Final composites were composed of a mixture of stable Mg2Zn11 and metastable Mg2Zn. In addition, no pure magnesium remained in the microstructure [23].

**Figure 1.** The microstructure of (**a**) Zn; (**b**) and (**c**) general view and detail of Zn-5Mg material prepared by powder metallurgy, respectively.

**Figure 2.** Line-scan of chemical composition at the interface of Zn and Mg.

#### *3.2. Mechanical Properties*

Tensile and compressive curves of Zn and Zn-5Mg and evaluated values of mechanical properties are shown in Figure 3 and Table 1, respectively. Due to similarities in the microstructure of the Zn matrix, the difference in mechanical behavior can be attributed to the addition of Mg particles. This improves the yield stress (TYS—tensile yield strength, CYS—compressive yield strength) and ultimate stress (UTS—ultimate tensile strength, UCS—ultimate compressive strength) values under both tension and compression loading. The general increase in presented values is about 40 MPa. Such an improvement is related to the presence of an Mg2Zn11 intermetallic phase formed on the interface between the Mg and Zn particles (Figure 2). On the contrary, the specified phase is brittle and causes a partial decrease in the ductility. Nevertheless, both Zn and Zn-5Mg are still characterized by excellent elongation (35 ± 4% and 16 ± 2%, respectively) compared to various Zn-based alloys [10]. Yang et al. [23] measured CYS for Zn-5Mg (186.63 ± 26.54 MPa) prepared by SPS and indicate the possible use of these materials for some applications without load-bearing requirements. However, the values of 148 ± 6 MPa for TYS and 183 ± 4 MPa for UTS measured for extruded Zn-5Mg are not far from generally required values for fixation medical devices (TYS > 200 MPa, UTS > 270 MPa), and suitable processing of these materials could lead to the desired values. Additionally, the value of modulus of elasticity (E) is slightly lower for Zn-5Mg (81 GPa) as a consequence of the lower value of E for pure Mg (45 GPa) compared to pure Zn (108 GPa). Such behavior is also desirable to prevent the stress shielding effect.

**Figure 3.** Compressive and tensile stress–strain curves for extruded Zn and Zn-5Mg.

**Table 1.** Tensile and compressive mechanical properties of Zn and Zn-5Mg materials after compaction by extrusion at 200 ◦C. Compressive tests were finished at 35% of relative deformation due to the absence of fracture.

#### *3.3. Corrosion Properties*

Corrosion rates evaluated from weight changes were 0.373 ± 0.004 and 0.043 ± 0.020 mg· cm−2·day−1 for Zn and Zn-5Mg respectively. Evaluation from Zn2+ concentration led to lower values, namely to 0.082 ± 0.016 and 0.022 ± 0.010 mg·cm−2·day−1, due to the presence of Zn in the solid state of corrosion products. This result reveals a surprising four-time decrease in released Zn2+ ions for Zn-5Mg. In fact, during in vitro cytotoxicity testing, even 50% dilution generally highly improves cell viability [15]. Regarding the alloying of Zn-based biodegradable alloys, such a huge change in corrosion rate has never been observed, to the best of our knowledge. To explain this behavior, potentiodynamic curves were measured after 1 and 12 h of immersion in SBF. The curves representing the polarization of pure Zn (blue lines in Figure 4) can be seen to be similar after 1 and 12 h in SBF. Both curves contain a specific linear region which indicates a cathodic reaction controlled by the diffusion of oxygen. The Ecor value of Zn-5Mg (−1.42 V) is due to the dissolution of less noble Mg, producing a more negative value compared to the Ecor of pure Zn (−0.98 V). At this potential, the reduction of water (Equation (1)) accompanied by hydrogen release is a dominant cathodic reaction.

$$2\text{H}\_2\text{O} + 2\text{e}^- \rightarrow \text{H}\_2 + 2\text{OH}^- \tag{1}$$

**Figure 4.** Polarization curves for extruded Zn and Zn-5Mg.

Magnesium works as a sacrificial anode and helps to protect the zinc matrix. Consequently, the decrease in current densities on the measured cathodic curve of sample Zn-5Mg immersed for 12 h was observed. The observed behavior caused the decrease in corrosion rate of the alloy during longer immersion tests and also lowered the amount of released Zn2+ ions during the initial stage of the corrosion process, which can particularly suppress the cytotoxic effect and improve the material biocompatibility. In contrary to the present study, Yang et al. [23] observed an increased corrosion rate for Zn-5Mg compared to pure Zn. They claimed that the corrosion rate of Zn-5Mg was increased as a consequence of micro-galvanic cells among MgZn2, Mg2Zn11, and pure Zn. In this study, the intermetallic phase occupied the only low area at the interface of pure Zn and Mg. Pure magnesium is even less noble than the presented phases, works as strong sacrificial anode, and enables the passivation of Zn. This leads in particular to the lower corrosion rates of the Zn-5Mg composite in this specific case and supports the idea that variation in phase composition of these types of materials can highly affect corrosion behavior.

#### **4. Conclusions**

This paper introduces a novel approach for the preparation of Zn-based biodegradable materials using the powder metallurgy technique. In this work, a Zn-5Mg alloy was successfully prepared by extrusion of a green compact containing a mixture of pure Zn and Mg powders. Compared to the product made using Zn powder, the addition of Mg was observed to positively affect TYS and UTS

values while maintaining excellent elongation. Additionally, Zn-5Mg was characterized by a decreased value of modulus of elasticity, which is highly desirable for fixation devices intended for traumatology applications, and a lower corrosion rate compared to pure Zn and Zn-based alloys. A lower amount of Zn ions released by the corrosion process can suppress cytotoxicity effects. Generally, the observed behaviour will lead to a significant improvement in the biocompatibility of the material.

**Author Contributions:** Conceptualization, J.K. and D.V.; methodology, D.D. and J.C.; investigation, J.K., D.D., ˇ and J.P.; writing—original draft, J.K.; writing—review and editing, D.V. and J.C.; supervision, D.V. ˇ

**Funding:** This research was funded by the Czech Science Foundation (project no. 18-06110S) and specific university research (MSMT No 21-SVV/2019) for financial support of this work.

**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*

## **Impact of the Morphology of Micro- and Nanosized Powder Mixtures on the Microstructure of Mg-Mg2Si-CNT Composite Sinters**

#### **Anita Olszówka-Myalska \*, Patryk Wrze´sniowski, Hanna Myalska, Marcin Godzierz and Dariusz Kuc**

Institute of Materials Science, Department of Material Science and Metallurgy, Silesian University of Technology, Krasi ´nskiego 8, 40-019 Katowice, Poland<EMAIL_ADDRESS>(P.W<EMAIL_ADDRESS>(H.M<EMAIL_ADDRESS>(M.G<EMAIL_ADDRESS>(D.K.)

**\*** Correspondence<EMAIL_ADDRESS>Tel.: +48-32-603-4422

Received: 26 August 2019; Accepted: 2 October 2019; Published: 4 October 2019

**Abstract:** The problem of preparing a ternary powder mixture, which was meant to fabricate sintered heterophase composite, and consisted of micro- and two nanosized powders, was analyzed. The microsized powder was a pure magnesium, and as nanocomponents, a silicon powder (nSi) and carbon nanotubes (CNTs) with 2% and 1% volume fractions, respectively, were applied. The powder mixtures were prepared using ultrasonic and mechanical mixing in technological fluid, and four mixing variants were applied. The morphology of the powder mixtures was characterized with scanning electron microscopy (SEM), and then, composite sinters were fabricated in a vacuum with hot temperature pressing at 580 ◦C under 15 MPa pressure, using a Degussa press. The reaction between the nSi and the Mg matrix, which caused the creation of the Mg2Si phase in the fabricated Mg-Mg2Si-CNT composite, was confirmed with X-ray diffraction (XRD). The porosity and hardness of the composite sinters were examined, and optical microscopy (OM) and quantitative image analyses were carried out to characterize the microstructure of the composites. In the manufacturing process of the Mg-nSi-CNT mixtures, the best results were the following: first separate de-agglomeration of nanocomponents, then their common mixing, and finally, the deposition of nanocomponents at the surface of the microsized magnesium powder. The applied procedure ensured the uniform layer formation of de-agglomerated nanocomponents on the Mg powder, without re-agglomerated nSi and CNTs. Moreover, this type of powder mixture morphology allows to obtain sinters with lower porosity and higher hardness, which is accompanied by precipitation of a finer Mg2Si phase. In the Mg-Mg2Si-CNT composite, the carbon phase was present, and it was located in the magnesium matrix and in silicide.

**Keywords:** heterophase magnesium matrix composite; Mg2Si; carbon nanotubes; nanopowders de-agglomeration; sintering

#### **1. Introduction**

Development in the fabrication of nanopowder materials, and their increased availability, is gaining momentum around the world. Nanopowder materials, such as carbon nanotubes (CNTs) [1–5], graphene [5–7], and nanosilver [8,9], due to their extraordinary properties, are applied as reinforcing components in composites with polymer [10–12], metal [13–17], and ceramic matrices [18,19]. Research focused on the use of nanocomponents concerns composites with the magnesium matrix, due to the low density of this metal [20–22]. However, proper use of the unique properties of nanocomponents is limited, due to the components' high tendency to agglomerate, which means that the consolidation of agglomerated nanofibers or nanoparticles with the matrix component can cause

a decrease in, for example, mechanical properties, as well as thermal or electric conductivity. In the case of metal matrix composites manufactured with powder metallurgy methods, powder mixtures with 0.5–4.0 vol.% of nanocomponents are usually applied, and much attention has focused on the formation of homogenous mixtures. Mixing in the ball mills is a method well-known in the literature for preparing powder mixtures, but this method can cause a loss of nanocomponents, the generation of internal defects and their re-agglomeration. Another solution, which has been applied to Mg-CNTs [23–25] and Mg-nSiO2 [26,27] mixtures, is ultrasonic mixing in technological fluids.

In the present work, ultrasonic mixing was applied to develop a beneficial procedure for preparing ternary mixtures, for a system that included a microsized powder and two nanopowders. That issue is crucial in the design of a new multiphase material technology. In the experiments, carbon nanotubes and silicon nanopowder (nSi) were used. During sintering, the silicon reacts with magnesium, and an intermetallic Mg2Si phase is formed [28–32]. Thus, from the Mg-nSi-CNT powder mixture, a composite with Mg2Si particles and nanotubes will be created. Due to the mechanical properties, magnesium silicide is often used as a reinforcement phase in ex situ and in situ composites with aluminum [33–36] and the magnesium [26–30] matrix. The intermetallic phase also exhibits good thermoelectric properties, and many studies have focused on that aspect of application [29,30]. However, the structural effect of the formation of Mg2Si in an environment enriched with CNTs is not known in the literature. Some movement of CNTs induced by synthesis reaction can be expected in the composite. The thermodynamic data indicate the reaction 2Mg + Si→ Mg2Si, and Gibbs free energy ΔG at 580 ◦C (the calculations were carried out using HSC Chemistry 4) reaches the value of −71.76 kJ. However, transformation of the CNTs into SiC is also possible, because ΔG for the reaction Si + C→SiC is also negative, and equals −64.96 kJ.

In the here presented studies, during the preparation of the ternary powder mixtures, it was assumed that de-agglomeration of the main nanocomponents would occur due to the ultrasound action in technological liquid (alcohol) [37] supported by mechanical mixing. The issue that requires detailed experiments and analyses is the correct order of the mixing steps which may limit uncontrolled re-agglomeration of the nanocomponents. Four variants of the mixing procedure were applied, and the effectiveness of the de-agglomeration was evaluated based on the powder mixture morphology. The porosity, hardness and microstructure of the composite sinters were also examined. Special attention was paid to Mg2Si phase dispersion, because the presence of nSi clusters in a ternary powder mixture induces relatively massive silicides in the composite, and on this basis, the effectiveness of the de-agglomeration of the nanocomponents can be characterized. The work concerns the general problem as the simultaneous use of fibrous and particulate nanocomponents in the conventional powder metallurgy processes, where a microsized powder is the main ingredient. This issue will occur in materials design regardless of the components chemical and phase composition. The powder ultrasonic mixing used in technological procedure allows to de-agglomerate primary nanofibers and nanoparticles agglomerates, and then their mixture with microsized component formation. The morphology of the mixture is controlled by interaction processes occurring in systems: nanofiber-microparticle, nanoparticle-microparticle, and nanofiber-nanoparticle. This type of powder mixture can be potentially applied to fabricate sintered composites or semi-products, in form of cold-pressed or hot-pressed moulds. In the literature, an application of multiphase moulds obtained by powder metallurgy was proposed for metal matrix composite fabrication with the different methods of plastic working processes [38–44] or casting methods [45–47]. The goal of such technological solutions is an increase of the reinforcing phases homogeneity, and its dispersion, in the case of in situ formed phases.

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

As a raw material in the experiments, magnesium powder (63–250 μm, Sigma Aldrich, 13112, Saint Louis, MO, USA), silicon nanopowder (average particle size 50 nm, Sigma Aldrich, Saint Louis, MO, USA) and carbon nanotubes (diameter 50–85 nm, length 10–15 μm, Multi-Walled Carbon Nanotube Powder, GRAPHENE SUPERMARKET, Ronkonkoma, State NY, US) were applied. Micrographs of

the fabricated Mg, nSi and CNT powders observed with a scanning electron microscope (SEM) are presented in Figures 1–3. The surface of the magnesium granules (Figure 1) is irregular, which might be a beneficial feature during the mixing process and the deposition of the nanocomposites on a single granule. The SEM micrographs of nSi (Figure 2) and the CNTs (Figure 3) (raw materials) revealed a few or dozen micrometer-sized clusters, and that disadvantage indicates the necessity of de-agglomeration in processing ternary powder mixtures. Otherwise, correct exploitation of unique nanoaddition properties will be ineffective, due to the submicro- and nanopores between single particles or fibers consolidated with the metal matrix, independent of the technology.

In the experiments, the same composition of the powder mixture, 97% Mg; 2% nSi; 1% CNT (vol.%), was chosen. Based on previous experiences regarding preparation of mixtures with nanocomponents [26, 27,33], four variants of mixture processing consisting of ultrasonic and mechanical mixing in alcohol were applied (Table 1). The main differences were in the sequence of de-agglomeration of the nanocomponents and then mixing with the microsized powder. When the mixing procedure was completed, the liquid alcohol was removed from the suspension, and then the powder mixture was dried for 18 h at 60 ◦C.

The microstructure of the ternary powder mixtures was examined with a field emission (FE)SEM (FE Hitachi S-4200, Hitachi Group, Tokyo, Japan), and the distribution of the components was analyzed. Then the Mg-nSi-CNT powder mixtures and pure Mg powder as a reference material were sintered in a Degussa press, in a vacuum atmosphere, at 580 ◦C under 15 MPa pressure. Sinters 20 mm in diameter and 10 mm high were obtained.

The porosity measurements of the fabricated composite samples were carried out with the Archimedes method, and Vickers hardness HV0.2 was determined with the Zwick 110 hardness tester (Zwick, Ulm, Germany). For the phase composition examination of the sinters, X-ray diffraction (XRD) was applied (X'Pert 3 Powder X- ray diffractometer, Malvern Panalytical Ltd., Royston, UK). For microstructure characterization, polished samples without additional etching were prepared. The observations were carried out with an optical microscope (OM; GX71 Olympus, Tokyo, Japan) and scanning electron microscope (SEM, Hitachi 3400N, Tokyo, Japan) equipped with wavelength-dispersive X-ray spectroscope (WDS Thermo Scientific Magna Ray, Waltham, Massachusetts, US). By WDS method, the elemental mapping of magnesium (TAP crystal), silicon (TAP crystal), carbon (NiC80 crystal), and oxygen (NiC80 crystal) were obtained. The quantitative metallography examinations were focused on Mg2Si phase characterization, the volume fraction evaluated with the area fraction and dispersion evaluated with the particle cross-section area. The examinations were performed using Met-Ilo software (J. Szala, Silesian University of Technology, Katowice, Poland), to analyze the influence of the morphology of the Mg-nSi-CNT powder mixture on the microstructure of the composite sinter. The stereological parameters were determined using OM images, although the resolution of method excludes identification of a single CNT, and allows observation only of bigger objects, due to the characteristic blue of the Mg2Si phase, identifying this silicide was very simple, in contrast to the SEM observation. The size of the Mg2Si phase particles and the quantity were measured, and then the particles divided into five size classes (left closed intervals: 10–100, 100–1000, 1000–10,000, 10,000–100,000 and > 100,000 μm2). The measurement procedure was conducted in a relatively large area (15 different areas were chosen for investigation, and they were conducted at 50X magnification). The image analyses allowed to indicate the most effective mixing method for preventing the re-agglomeration phenomenon. In the case of this phenomenon, the methodology is more accurate in comparison with thin foils, for example, where only a small area of the material can be analyzed. Additionally, the silicide particles detected with the OM, and characterized with quantitative metallography, were only a portion of the nSi and Mg reaction product. The other particles were smaller, even nanosized, and more difficult to distinguish.

**Figure 1.** SEM micrographs of Mg powder as-fabricated: (**a**) particles, (**b**) surface of single particle.

**Figure 2.** SEM micrographs of silicon nanopowder (nSi) as-fabricated: (**a**) agglomerates of different size, (**b**) morphology of single agglomerate.

**Figure 3.** SEM micrographs of as-fabricated carbon nanotubes (CNT): (**a**) agglomerates of different size, (**b**) morphology of single agglomerate.

**Table 1.** Applied procedure of ternary powder mixture preparation and condition of their sintering process (subscript D means the ultrasonic de-agglomeration in alcohol).

#### **3. Results**

#### *3.1. Characterization of the Powder Mixtures*

Examinations of the Mg-nSi-CNT powder mixtures with SEM revealed that the main effect of the mixing procedure was the deposition of the nanopowders on the Mg powder. This phenomenon is known from the literature for a binary mixture of microsized and nanosized components and occurs due to adhesion forces [23–28]. Moreover, the macroscopic observations of the mixtures did not reveal a significant residue of nanocomponents on the mixer walls after mixing, and generally, the metallic color of the magnesium granules disappeared. Additionally, the technological liquid was transparent after the process.

The microscopic examination of the dried ternary powder mixtures (Figures 4–7) showed differences in the distribution of the CNTs and nSi, which was related to the mixing procedure. Magnesium powder was coated with nanocomponents, but an unfavorable phenomenon was observed in the case of the (Mg + nSiD) + CNTD and (Mg + CNTD) + nSiD samples (Figures 6a and 7a). Agglomerates of CNTs and nSi, a few micrometers and larger, were detected, and they were poorly connected to the Mg grains, and even separate from the metallic granules. This effect was more intense for the (Mg + nSi)D + CNTD mixture, but in general, the tendency to re-agglomeration of two different nanocomponents in the ultrasonic mixed suspension was noticed. It occurred when the microsized component (Mg) was not mixed at the same time with the two previously de-agglomerated nanocomponents. The main reason for that type of re-agglomeration likely is the fibrous morphology of the CNTs. In the procedure for mixing the (Mg + nSiD) + CNTD powder, the nSi particles initially deposited at magnesium were caught by the de-agglomerated fibers, while in the processes for preparing (Mg + CNTD) + nSiD, the nSi accumulated around the CNTs.

**Figure 4.** SEM micrographs of Mg+(nSiD+CNTD) powder mixture: (**a**) morphology of powder mixture, (**b**) surface of microsized Mg coated with CNT and nSi.

**Figure 5.** SEM micrographs of Mg + (nSi + CNT)D powder mixture: (**a**) morphology of powder mixture, (**b**) surface of microsized Mg coated with CNT and nSi.

**Figure 6.** SEM micrographs of (Mg + CNTD) + nSiD powder mixture: (**a**) morphology of powder mixture, (**b**) surface of microsized Mg coated with CNT and nSi.

**Figure 7.** SEM micrographs of (Mg + nSD) + CNTD powder mixture: (**a**) morphology of powder mixture, (**b**) surface of microsized Mg coated with CNT and nSi.

Moreover, the microstructure observations indicated that in the proposed mixing procedures, for one micro- and two nanosized components, selective deposition of previously de-agglomerated nanocomponents could not be expected. This means that the formation of a layered structure, where first, nSi covers the Mg particles, and then the CNTs create another layer, or vice versa, cannot be achieved. That excludes formation of the CNT nanozone at the Mg-nSi interface, and limits further de-agglomeration of the fibrous component, induced in conditions where the silicide forms not by reactive diffusion, but by self-propagating high-temperature synthesis (SHS).

In comparison with the powder mixtures described previously, the results of the FE-SEM examinations for the two other mixtures, Mg + (nSi + CNT)D and Mg + (nSiD + CNTD), revealed some differences in the microstructure. Generally, the Mg surface was coated with distinctly separate CNT and nSi grains, and the layer of nanocomponents was uniform and relatively thin (Figure 4b). However, when the CNT and nSi were de-agglomerated together before being introduced in the Mg-alcohol suspension (the Mg + (nSi + CNT)D mixture), the agglomerates of CNT + nSi at the Mg surface were revealed. That effect may confirm a tendency of CNTs and nSi to mutual agglomeration, induced by the fibrous component.

#### *3.2. Characterization of Composite Sinters*

An example of the XRD pattern obtained for the composites is shown in Figure 8, and it confirms the presence of αMg and a new Mg2Si phase, formed in situ as a result of the nSi reaction with the Mg matrix. A weak signal coming from MgO phase was identified as well. The presence of a carbon component, or SiC, was not detected due to the low carbon component content and insufficient method sensitivity. Therefore, in a future investigation, the identification of nanostructural phases in composite will be performed, and the high-resolution transmission electron microscopy and selective area electron diffraction methods will be conducted.

**Figure 8.** X-Ray Diffraction Pattern of Diffraction Mg + nSiD + CNTD composite.

Results of composite microstructure examinations with the OM without additional etching are shown in Figures 9–12. Two main structural elements were revealed: rounded magnesium areas (light) with characteristic MgO boundaries derived from the initial magnesium powder, and an irregular dark blue phase of similar or greater size (Figures 9a, 10a, 11a and 12a). At higher magnification (Figures 9b, 10b, 11b and 12b), it is visible that the irregular dark blue phase contains very fine black and elongated phases, and this suggests that the particles are a mixture of Mg2Si and CNTs. Similar black elongated phases were detected in the initial magnesium powder grains except the oxides (beige), and inside the Mg grains as well. Within the Mg grains, light blue, irregular and very fine precipitations of Mg2Si can be observed, and that suggests the SHS reaction. Moreover, that explains the movement of CNTs, previously deposited on the microsized Mg powder, in the magnesium matrix. The comparison of the microstructure of all composites indicates that in the material obtained from the Mg + (nSiD + CNTD) mixture, the Mg2Si phase is the finest (Figure 9), and more beneficial conditions for matrix reinforcement formation are created.

**Figure 9.** OM micrographs of Mg-Mg2Si-CNT composite fabricated from Mg+(nSiD+CNTD) powder mixture: (**a**) Mg2Si agglomerates in Mg matrix, (**b**) phases of different size and morphology.

**Figure 10.** OM micrographs of Mg-Mg2Si-CNT composite fabricated from Mg + (nSi + CNT)D powder mixture: (**a**) Mg2Si agglomerates in Mg matrix, (**b**) phases of different size and morphology.

**Figure 11.** OM micrographs of Mg-Mg2Si-CNT composite fabricated from (Mg + nSiD) + CNTD powder mixture: (**a**) Mg2Si agglomerates in Mg matrix, (**b**) phases of different size and morphology.

**Figure 12.** OM micrographs of Mg-Mg2Si-CNT composite fabricated from (Mg + CNTD) + nSiD powder mixture: (**a**) Mg2Si agglomerates in Mg matrix, (**b**) phases of different size and morphology.

Results of composite microstructure examinations with the SEM and WDS are presented in Figure 13. The results revealed distribution of Mg, Si, O and C. The overlapping areas, enriched in Mg and Si, confirmed in situ formation of Mg2Si particles in the composite. Moreover, in the regions containing Mg2Si particles, an increase of carbon concentration was observed on the carbon elemental mapping (Figure 13e), which may suggest the presence of CNT's. Furthermore, the microareas with higher concentration of oxygen were found in the magnesium matrix. The oxygen distribution indicates oxides presence originating from the initial Mg powder surface. Obtained WDS elemental mapping results are consistent with the OM and XRD results.

Results for the porosity and microhardness measurements are presented in Table 2, and the differences depending on the mixing procedure and the effectiveness of the de-agglomeration of the nanocomponents are demonstrated. The lower porosity of the composite obtained from the initial powder mixture with the uniform nanocomponents distribution, that is, Mg + (nSiD + CNTD), indicated that this type of powder mixture morphology is the most effective in material compaction, and it influences the composite microhardness, which is a little bit higher for a sinter obtained from this mixture. The highest porosity was obtained for the reference Mg sinter, in comparison with the composites obtained. This result can be explained by the impact of an exothermal reaction between Mg and nSi, which induced a local increase in the temperature in the composite sinter, and better compaction under the same pressure applied.

**Table 2.** Hardness and open porosity of composite sinters obtained from Mg-nSi-CNT mixtures of the same composition and prepared by different way.

An example of the image transformation procedure applied for the quantitative metallography examination of the Mg2Si phase with Met-Ilo software is presented in Figure 14. For each sample, 15 areas were measured, and the results are presented in Table 3 and in Figure 15.

Differences in Mg2Si synthesis depending on the preparation of the powder mixture were revealed. The main difference was in the value of the silicide area fraction AA, which is two times lower for the composites obtained from Mg + (nSi + CNT)D and Mg + (nSiD + CNTD) mixtures in comparison to (Mg + nSiD) + CNTD and (Mg + CNTD) + nSiD. That effect was obtained for the same initial powder composition and sintering parameters, which suggests that very fine Mg2Si particles smaller than 10 μm2 were formed, and many below the OM resolution. The analysis of the detected number of particles in size classes also exhibited an evident difference for the composites from the (Mg + nSiD) + CNTD and (Mg + CNTD) + nSiD mixtures. For those composites, the number of particles was greater compared to the two other samples. This effect was directly connected to the morphology of the mixture of the initial powders, and it can be explained as a result of the Mg reaction with few micrometer-sized nSi/CNT agglomerates, located either on the Mg microsized powder surface or separately from the metallic particles (Figures 6 and 7).

The preliminary research on the hybrid composite obtained with powder metallurgy showed that the parameters applied for Mg-nSi-CNT sintering and the mixture composition required optimization.

**Figure 13.** SEM micrographs of Mg-Mg2Si-CNT composite fabricated from Mg+(nSiD+CNTD) powder in composition mode (**a**) with WDS elemental mapping of Mg (**b**), Si (**c**), O (**d**) and C (**e**).

**Figure 14.** Procedure applied for Mg2Si particles detection in Mg-Mg2Si-CNT composite sinters: (**a**) initial OM image, (**b**) grey image-histogram equalization, (**c**) numerical inversion, (**d**) automatic binarization (k-means; white phase), (**e**) opening 2 and (**f**) complete image - initial + binary.

**Table 3.** Quantitative characteristics of Mg2Si particles sized of more than 10 μm2 formed in Mg-Mg2Si-CNT composite sinters.

**Figure 15.** Quantity of Mg2Si particles in composite sinters divided into size classes.

From the literature [13,14,21–23] it is known, that due to CNT's tendency for agglomeration, an application of more than 1 vol.% of CNT's may cause a decrease of mechanical and electrical properties in metal matrix composite. Thus, in the present work, the highest effective reinforcing amount of fibrous carbon nanocomponent was used. In our previous experiments, the procedure of Mg-nSi-CNT mixture preparing with proposed method was successfully examined to volume fraction less than 2% of nSi and 1% of CNT. However, the effects related with Mg2Si formation can be useful

in CNT de-agglomeration processes. Therefore, in further experiments two main aspects should be considered. The first aspect is the ratio of nanosilicon to CNT in the powder mixture, and the maximum number of nanocomponents, which can introduced into the Mg-base powder mixture. The second important issue is the adjustment of sintering parameters such as heating speed, sintering temperature and pressure.

Generally, the experiment revealed the possibilities of a design for micro- and nanopowder mixture morphology, and its influence on the final product microstructure and properties, including ex situ and in situ nano-reinforcement in composites.

#### **4. Conclusions**

A novel approach for processing ternary powder mixtures intended for fabricating hybrid composites, which consisted of a microsized component and two nanosized components, was presented in this work. The main conclusions are drawn as follows:

**Author Contributions:** Conceptualization, A.O.-M.; methodology, A.O.-M. and H.M.; validation, H.M., P.W. and M.G.; formal analysis, A.O.-M. and P.W.; investigation, H.M., P.W. and M.G.; resources, D.K.; data curation, H.M., P.W., M.G. and D.K.; writing—original draft preparation, A.O.-M.; writing—review and editing, A.O.-M., H.M. and P.W.; visualization, P.W.; supervision, A.O.-M.

**Funding:** The study was carried out as a part of statutory research of the Institute of Materials Science at the Silesian University of Technology for 2019 (BK-205/ RM0/2019).

**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/).

## *Review* **The Critical Raw Materials in Cutting Tools for Machining Applications: A Review**

#### **Antonella Rizzo 1,\*, Saurav Goel 2,3, Maria Luisa Grilli 4, Roberto Iglesias 5, Lucyna Jaworska 6,7, Vjaceslavs Lapkovskis 8, Pavel Novak 9, Bogdan O. Postolnyi 10,11 and Daniele Valerini 1**

Received: 20 December 2019; Accepted: 2 March 2020; Published: 18 March 2020

**Abstract:** A variety of cutting tool materials are used for the contact mode mechanical machining of components under extreme conditions of stress, temperature and/or corrosion, including operations such as drilling, milling turning and so on. These demanding conditions impose a seriously high strain rate (an order of magnitude higher than forming), and this limits the useful life of cutting tools, especially single-point cutting tools. Tungsten carbide is the most popularly used cutting tool material, and unfortunately its main ingredients of W and Co are at high risk in terms of material supply and are listed among critical raw materials (CRMs) for EU, for which sustainable use should be addressed. This paper highlights the evolution and the trend of use of CRMs) in cutting tools for mechanical machining through a timely review. The focus of this review and its motivation was driven by the four following themes: (i) the discussion of newly emerging hybrid machining processes offering performance enhancements and longevity in terms of tool life (laser and cryogenic incorporation); (ii) the development and synthesis of new CRM substitutes to minimise the use of tungsten; (iii) the improvement of the recycling of worn tools; and (iv) the accelerated use of modelling and simulation to design long-lasting tools in the Industry-4.0 framework, circular economy and cyber secure manufacturing. It may be noted that the scope of this paper is not to represent a completely exhaustive document concerning cutting tools for mechanical processing, but to raise awareness and pave the way for innovative thinking on the use of critical materials in mechanical processing tools with the aim of developing smart, timely control strategies and mitigation measures to suppress the use of CRMs.

**Keywords:** critical raw materials; cutting tools; new materials; new machining methods; modelling and simulation

#### **1. Introduction**

In the era of globalisation and high competitiveness, it is of utmost importance for industries to work on reducing manufacturing costs and simultaneously providing added value in terms of increased life of the product when put into service. It is also imperative for them to have a flawless supply chain which contains a key component of sourcing critical raw materials (CRMs) [1] for them to remain sustainable.

In the EU, the excess of imports of CRMs, especially in the mechanical manufacturing industry—here, we refer to metals that are vital to EU industries such as tungsten, chromium, and niobium to name but a few—has reached an alarming stage, bringing a great degree of dependency on countries that are monopolistic suppliers of these CRMs. To mitigate this foreseeable problem, the EU has already launched several campaigns, including the "Raw Materials Initiative" (RMI) [2].

The issue of CRMs must be tackled with scientific rigour by pursuing different parallel actions; in particular, by (1) improving the production processes of CRMs (increasing sustainable mining, reducing extraction costs, increasing the efficiency of materials, increasing security, etc.); (2) finding suitable candidates to partially or totally substitute the CRMs; and (3) increasing their recycling.

In this context, the development of new materials with superior characteristics or better performance than existing materials is desirable to lead to a longer product life and therefore to reduce the cost of the product. The cost of cutting tools used in the manufacturing industry at present, particularly to machine high-value components such as turbine blades, automotive and aerospace parts, machine parts and biomedical implants, is significant and dictates the total manufacturing cost and thus the final price of the product [3]. The tooling cost is not only associated with the cost incurred on relapping the tools, but it also involves the cost incurred due to the increased cycle time owing to the unloading and reloading of the new tool. The main focus in manufacturing research currently is to improve productivity, and thus, there is a focus on developing materials (especially for cutting tools) that can withstand higher cutting loads than the current attainable machining limits. The cutting tools discussed here are meant for all mechanical-based contact loading processes, including, for example, milling, drilling, turning, honing, chamfering, hobbing, knurling, parting-off, and so on. The current development of new materials specifically for cutting tools has remained a slow process, particularly due to the high initial investment, but interest has been growing recently towards developing more sustainable ways to de-risk the supply chain.

Cemented carbide is one of the most popular cutting tool materials. Typically, a carbide cutting tool is manufactured with a mixture of tungsten and cobalt (the binder that holds tungsten carbide together), with a wealth of variations in carbide grain size and the ratio of carbide to binder. Preferred blends have been developed over time to achieve effectiveness with different cutting depths and widths, as well as workpiece materials. The distribution of cutting tools in the global market in 2018 by the cutting method and workpiece material is shown in Figure 1a,b, respectively. It can be seen that milling, turning and drilling are the processes mostly used in primary machining operations, and the tools for these operations capture almost 87% of the total tooling market. In terms of the type of cutting material, cemented carbides capture one half of the market, followed by high speed steel. Ceramics, cermets and superhard materials such as polycrystalline diamond (PCD) and polycrystalline cubic boron nitride (PCBN) capture the remaining share of the tool material market. The two important ingredients of the carbide tools mentioned above—namely tungsten and cobalt—were earlier identified (in 2011) as being among the list of the 14 critical raw materials (see Figure 2) vital to EU industries [4]. More recently, in 2014 and in 2017, they have continued to remain on this list on account of their economic importance and risk of supply interruption [5,6]. Another driving force for Co substitution lies in its well-known genotoxic and carcinogenic activities. A pressing need has thus emerged for the

EU to develop alternative solutions to the existing cutting tool materials to avoid a complete closure of machining businesses in the absence of the raw material of the cutting tools.

**Figure 1.** Cutting tool global market distribution by cutting technology (**a**) and by workpiece material (**b**) as presented by Dedalus Consulting. Data taken from [7].

**Figure 2.** Critical raw materials list for 2011–2017 overlaid on the periodic table of the elements [5,6].

Perhaps, the most direct impact of the new era of digital manufacturing [8] on future cutting tools is the potential use of additive manufacturing technologies for the generation of quasi-network-shaped tools with geometric structures and challenging sub-structures and the potential to produce under-functionally different structures and optional materials with graded properties (4D printing). One of the four themes defined by the Roadmap of High Performance Cutting (HPC) [9] points at the use of integrated tooling based on cycle time reduction and quality improvement in high-end batch production. This theme was further developed, and the term "smart tooling" has been coined [9], which refers to the insertion of sensor actuators [10] within the cutting tool as a step forward in the process monitoring of the tool, allowing users to collect the necessary data for the creation of more accurate digital twins for the machining processes (Figure 3).

**Figure 3.** (**a**) Composition of the smart cutting tool, (**b**) assembly of the tool, (**c**) cross-section view of the tool, as taken from [10].

An example of the additional functions which are commercially available in this regard is the use of low-frequency vibration-assisted drilling (LFVAD) [11] for improving the quality of the machining of brittle materials such as carbon fibre reinforced composite (CFRP) via an interrupted cutting mechanism.

Besides such alterations in machining mechanisms, there is remaining interest regarding the successful realisation of novel tool coating materials as well as the development of enhanced cutting technologies, including laser-assisted cutting and cutting with cryogenic assistance. The idea behind the use of these methods depends on whether the main wear mechanism of the tool is governed by the high physical hardness of the workpiece or the high chemical affinity. In the former case, laser heating ahead of the cutting tool can reduce the cutting resistance, making the workpiece more compliant to cutting, while in the latter case, the cryogenic environment helps delay the kinetics of the chemical reactions, thus inhibiting the occurrence of diffusion-induced chemical wear.

As an additional strategy, there are many existing coatings that protect cutting tools working under extreme conditions. The most desired attribute of tool coatings is to provide high hardness and toughness for high wear performance and thermal and chemical stability to withstand extreme cutting environments, especially during high-speed machining. Most of the widely used coatings do not contain CRMs, but multi-elemental and high-entropy coatings designed for specific extreme conditions may include them [12]. Transition metals are one of the biggest groups of chemical elements in the periodic table and, as can be seen in Figure 2, half of the refractory metals group are marked as CRMs. However, the amount of material needed for a coating is significantly lower than for analogous bulk CRMs used in cutting tools. Moreover, the nano and microstructural design of protective coatings, such as nanocomposites or multilayers, suppresses the usage of certain CRMs in coatings.

Several recent studies have addressed technical–scientific actions with different methods that can be pursued to address the problems of CRMs in cutting tools; in particular, to achieve the following:

In this paper, we summarise the current understanding regarding the use of CRMs in the tooling industry, considering the four aspects listed above, without any intentional effort to propose an immediate solution to this problem, but with the aim of stimulating the development of effective strategies in this field.

#### **2. Attempts to Expand the Life of Tungsten Carbide (WC)–Co Based Tools**

There are several methods to increase the lifetime and efficiency of tools in machining applications, such as the proper modification of the base tungsten carbide (WC)–Co material of the tool, the use of advanced processing techniques, and the application of protective coatings. These methods are reviewed in the following subparagraphs.

#### *2.1. WC–Co Based Materials*

Cemented carbides belong to the most common and the longest-used tool materials produced by powder metallurgy methods. Sintered carbides are characterized by their high strength and abrasion resistance and include one or more high-melting metal carbides constituting the basic component together with the metallic binding phase [13]. According to this standard, sintered carbides are divided into groups: those used for the production of tools for metal machining, plastic forming and for the use of mining tools. The basic component of cemented carbides is WC, which, depending on the manufacturer and group of material applications, can constitute from 50% to 90% by weight of the sintered content. The other ingredients are carbides of titanium (TiC), tantalum (TaC) and niobium (NbC), the content of which can be from 0% to 35% by weight. These carbides dissolve each other and can also dissolve a large amount of tungsten carbide. The rest of the composition is usually cobalt. Cobalt is characterized by very good wettability with most materials that are components of sintered tool materials and a fairly high melting point, which is 1493 ◦C. Such a microstructure, which is characteristic for sintered carbides, allows for the presence of a ductile binding phase; additionally, the hard and brittle carbide phase allows the bonding of opposing features in one material, such as high abrasion resistance and hardness with high strength and fairly good ductility. Many carbide applications feature very good abrasion resistance, which depends on the chemical composition.

Two-component cemented carbides of the WC–Co type with a low content of cobalt are characterized by the highest abrasion resistance. These grades can be used if there are no impacts during operation, and abrasion is the main wear mechanism.

The universality of WC–Co results from its very good mechanical and tribological properties. The technological process of obtaining cemented carbides is characterized by relatively easy formation, sintering temperatures lower than ceramic materials and electrical conductivity, and these factors have a positive effect on the possibility of shaping products with complicated geometry using erosive treatment as well as on the ease of applying anti-wear and anti-corrosion coatings. The disadvantage of this material is that cutting tools made of cemented carbides work at relatively low cutting speeds, and these materials tend to oxidize already at temperatures above 400 ◦C. The most commonly used carbides are tungsten, titanium, tantalum or other high-melting metal elements in an amount of 75 to 94 wt.% and cobalt, nickel or molybdenum, and sometimes other metals are usually used as the binding matrix. The toughness of WC–Co can be increased by increasing the cobalt content, whilst the wear resistance is increased by decreasing the cobalt content and decreasing the carbide grain size. The substitution of cubic carbides (TiC, NbC and TaC) for WC leads to improvements in wear resistance and resistance to plastic deformation. The microstructure of cemented carbide WC–Co consists of tungsten carbide particles combined with cobalt, which are obtained in the sintering process with the participation of the liquid phase. When machining a metal with high plasticity, such as pure iron, using conventional WC–Co cemented carbide, a chip tends to adhere on the rake face of the cutting tool, resulting in serious adhesive wear due to the existence of cobalt, which has a lower melting point compared to WC. The presence of cobalt in cemented carbides during the machining of steel, for example, causes the chip to stick to the cutting blade. Crater wear is caused by the chemical interaction between the rake face and hot chip. Wear occurs by the diffusion of the tool material into the chip or by the adhesion between the chip and tool followed by a fracture below the adhered interface within the tool material. Crater wear can be reduced by increasing the chemical stability of the tool material, decreasing the solubility of the workpiece or barrier protection by substrate alloying or coating [14]. Thermal shock cracks are caused by large temperature gradients at the cutting edge. There is a large

difference in the coefficients of thermal expansion between cobalt and tungsten carbide, which is why cracks may appear on the blade during its operation. Therefore, cemented carbides are often used as substrates for coatings.

However, studies are being carried out on uncoated WC–Co intended for machining materials. An example consistent with the idea of reducing the use of critical materials in cemented carbides is cobalt-free carbide. The first cobalt-free carbides were obtained by conventional methods from micropowders [15]. These materials were characterized by high hardness and abrasion resistance, but exhibited greater brittleness compared to cemented carbide WC–Co. For binderless tungsten carbide, the sintering took place in a solid phase, which for conventional, pressure-free sintering means higher sintering temperatures and/or longer sintering times, which in turn cause grain growth in the polycrystalline material. The use of nano-powders and pressure sintering methods (spark plasma sintering, hot pressing sintering) improved the properties of the cobalt-free cemented carbides [16–19].

The addition of free carbon to the binderless WC reduces the amount of the brittle phase of W2C, Co3W3C and oxide creation [20–22].

Spark plasma sintering (SPS, also called "Field Assisted Sintering Technology") and hot pressing sintering, thanks to the shorter sintering time and pressure, allow the reduction of grain growth and of the porosity of materials sintered in the solid phase [23].

Composite materials with the addition of ceramic powders are another idea which has been developed in relation to cemented carbides. In the case of ZrO2, additional reinforcement is used, which is the result of stresses created during the ZrO2 phase transformation. They are usually added in an amount of 5% to 15 wt.% in the form of oxides, carbides and nitrides [24,25].

This idea is consistent with the substitution of the critical materials of W and Co. Composite materials with a WC or WC–Co matrix have recently been developed very intensively thanks to the SPS method [26]. There are already individual examples of the industrial application of such solutions.

Research related to the reduction of cemented carbide consumption is also associated with attempts to increase their durability. Superhard materials such as diamond and cubic boron nitride (cBN) are introduced into the WC–Co [27]. Composites containing ultrafine tungsten carbide/cobalt (WC–Co) cemented carbides and 30 vol.% cBN were fabricated mainly by FAST (field-assisted sintering technology) methods. WC–Co/cBN composites have been considered as a next-generation material for use in cutting-tool edges and are characterized by an optimal combination of hardness and toughness. The major challenge in sintering these composites is to produce a well-bonded interface between the WC–Co matrix and cBN particles [28–31].

Whisker toughening is mainly used for binderless tungsten carbide. There are known studies in which SiCw, Si3N4w and Al2O3 whiskers were used. The whisker participation was up to 10 vol.% [32,33].

The literature describes research related to toughening using nanotubes and graphene to improve the thermal conductivity of cemented carbides. The study found that thermal stress is the main reason for the failure of cemented carbide shield tunnelling tools when shield tunnelling is carried out in uneven soft and hard soil [34].

#### *2.2. Advanced Machining Techniques*

The machining of "difficult to cut materials" such as the nickel-based superalloys used in turbine blades requires an extended tool life for unhindered machining [35]. Traditional carbide tools are limited to working in the range of 30 m/min to 70 m/min because of their poor thermochemical stability; however, they can be used at high feeds due to their high toughness. Improvements in the poor thermochemical stability of cemented carbide tools could be aided by advanced modelling. Additive manufacturing technology [36] and new machining techniques have now opened newer possibilities of making complex tool shapes. This section presents a brief review of recent attempts made to expand tool life either by reducing the load on the cutting tool or by deferring the known pathway of the accelerated wear mechanism. In either case, the end result is the improved quality of the machined part.

#### 2.2.1. Machining by Laser Assistance (Thermal-Assisted Machining)

In pursuit of overcoming the challenges of difficult-to-machine materials, laser assistance has been adopted worldwide for incorporation with mechanical micromachining. To date, seven major patents (in the years 1982 [37,38], 2006 [39], 2013 [40], 2011 [41], 2014 [42], 2016 [43], and 2017 [44]) have been granted in the US concerning the use of laser assistance during mechanical micromachining. All these patents can be connected by a single chain of a process which is now well known as thermal-assisted machining (TAM) [45]. The concept of TAM is schematically shown in Figure 4a and relies on pre-heating the cutting zone of the material being cut ahead of the cutting tool. This methodology reduces the physical hardness of the workpiece, making it more compliant to cutting by reducing the specific cutting energy (i.e., the work done by the tool in removing a unit volume of the material).

**Figure 4.** Important advancements made in machining technology. (**a**) Thermal-assisted machining [46]; (**b**) cryogenic machining [47]; (**c**) schematic diagram to illustrate the mechanism of vibration-assisted machining [48].

#### 2.2.2. Cryogenic Machining

As shown in Figure 4b, cryogenic machining relies on freezing the cutting tool to extreme temperatures of about −196 ◦C by means of either liquid nitrogen or liquid carbon dioxide [49]. Cryogenic machining helps to serve two purposes: (i) it elevates the relative hardness gradient between the tool and the workpiece; and (ii) it delays the kinetics of any chemical diffusion that may likely trigger the tribo-chemical wear of the tool. S. Rakesh highlighted how cryogenic technology is ecological, non-toxic and non-explosive [50].

Cryogenic cooling has been executed in cutting operations in different ways by using liquid nitrogen for precooling the workpiece, cooling the chip, and cooling the cutting tool and cutting zone [51]. Numerous studies have compared conventional cutting strategies and cryogenic cooling methods. However, studies conducted in an attempt to determine the best technique found many contradictions, as the conclusions described above could change in relation to tool–work pairs, to the cutting conditions and to the general evaluation parameters.

The literature dealing with tungsten carbide cutting tools in cryogenic processing is significantly less numerous compared to the literature on steel tools, and only a few papers consider an evaluation parameter of the tool's lifetime [52]. In this context, all these studies show that the cryogenic treatment has positive effects both on the lifetime of the tools and on the surface finishing of the product. The cryogenic treatment of the materials showed significant positive effects such as increased wear resistance, reduction of residual stresses, increased hardness and fatigue resistance.

Seah et al. [53] carried out a series of experiments with the aim to study the aspects of cold and cryogenic treatments on uncoated WC inserts for carbon steel ASSAB 760. They showed that at different cutting speeds the cryo-treated inserts exhibited greater resistance to wear than the untreated and recovered counterpart. In addition, they found that the cold and cryogenic treatment significantly increased the resistance of the cutting insert to chip removal, which became increasingly important as the cutting speed increased.

Yong et al. [54] subjected uncoated WC inserts to a cooling treatment down to −184.5 ◦C for 24 h and then heated the insert to room temperature, keeping the rate of 0.28 ◦C/min unchanged both when increasing and decreasing the temperature. They developed a series of face milling operations using different cutting speeds but kept all the other processing parameters constant using untreated and cryo-processed inserts. Two pieces of information were highlighted: the first concerned the cryo-processed inserts, which generally performed better than their untreated counterparts; and the second focused on the increase in tool life of 28%–38% during the cryogenic treatment in wet machining compared to dry machining.

Sreeramareddy et al. [55] studied the tool wear, cutting forces and the surface finish of parts worked using a WC insert coated with a multilayer and subjected to cryogenic treatment. They showed that cryogenic processing reduced the wear on the side of the inserts as well as the cutting forces and the surface roughness of the workpiece machined with untreated inserts.

These studies suggest that the cryogenic treatment of carbide tools is capable of improving the productivity and quality of the final product because it guarantees greater resistance to wear and surface finish. Improvements have been reported in the red-hardness of cryogenically treated inserts, which resulted in low flank wear [56].

Bryson [57] claimed that the increase in wear resistance, with a consequent increase in the lifetime of the carbide tool, was due to the greater strength of the binder after cryogenic treatment.

Thakur et al. [58] highlighted that WC tools undergo a less-strong microstructural modification under cryogenic treatment compared to that detected with conventional heat treatments; some physical transformations actually occur concerning the densification of cobalt, which induces an increase in the gripping of carbide particles and an improvement in the tool life amounting to a wear resistance increase of 27%. The cryogenic treatment of the tool is one approach to enhance its properties by introducing microstructural changes. The formation of complex compounds such as Co6W6C or Co3W3C might have increased the hardness in the samples due to forced air cooling and oil quenching.

In another work, the effect of cryogenic treatment on the tool was shown to include an increase in tool life, lower cutting force and better surface finish compared to the untreated condition [59].

Reddy et al. [60] examined the workability of C45 steel with untreated and treated (−110 ◦C for 24 h) ISO P-30 tungsten carbide inserts by measuring the flank wear, main cutting force and surface finish. They concluded that the best machinability found was caused by the increase of the thermal conductivity of tungsten carbide induced by cryogenic treatment. The surface roughness of the workpiece was lower by approximately 20% when the workpiece was machined with deep cryogenically treated tungsten carbide tool inserts in comparison with untreated inserts for cutting speeds in the range between 200 and 350 m/min. Vadivel et al. [61] studied the microstructure of cryogenically treated (TiCN + Al2O3) coated and untreated inserts in turning nodular cast iron. Their results highlighted that coated and treated tools have better properties that help the cutting tool to operate under hostile conditions for a longer time.

A. Swamini et al. [62] drew up an overview of the metallurgy behind the cryogenic treatment of cutting tools; their results regarding WC–Co tools can be summarized as follows:

M.I. Ahmed et al. [63] used a modified tool holder for the efficient use of cryogenic cooling for machine cutting. The modified tool holder uses the direct continuous contact of liquid nitrogen with the cutting insert for the perfect cooling of the cutting tool. The results showed an average 30-fold increase in the lifetime of the carbide tools.

M. Dhananchezian [64] described the machinability characteristics of turning Hastelloy C-276 with a Physical Vapour Deposition (PVD) coated nano-multilayer TiAlN carbide cutting tool using dry turning and liquid nitrogen cooling methods. The improvement of cutting tool performance was achieved under liquid nitrogen cooling by the control of the wear mechanisms, which in turn reduced the wear rate. Yildirim [65] examined the effect of some cooling conditions on machinability which may be an alternative to conventional cooling. The results showed that the 0.5 vol.% hBN cooling method used in conjunction with liquid nitrogen gave the best results in terms of the machining performance and lifetime of carbide tools.

Particularly noteworthy is the recent work of Biswal et al. [66] which, based on experimental results, highlighted how uncoated cryogenically treated tempered cermet inserts perform better than other cermet inserts thanks to their better wear resistance, micro-hardness and toughness. Despite all this, the cryogenic processing technique has not yet displaced conventional industrial processing. Further studies on cryogenic processing are necessary in order to highlight the increase in the lifetime of WC tools with the consequent saving of critical materials such as W and Co.

2.2.3. Vibration-Assisted Machining

The literature refers to vibration-assisted machining (VAM) as the process of intermittent cutting. The reported benefits include the low wear rate of the tools, reduced burr formation on the workpiece and higher cutting depths being achieved.

The attainable cutting speeds during VAM are limited by the hardware and system, and hence ultrasonic-assisted machining methods are classed as low-speed machining techniques [67]. Also, the process of vibration assistance can be implemented in two ways:

The resonant system operates at discrete frequencies, normally higher than 20 kHz and at amplitudes of less than 6 μm, whereas the non-resonant system operates at frequencies between 1 to 40 kHz and with amplitudes 10 times higher than the resonant system. As shown in Figure 4c, the tool is prescribed an oscillatory motion by a wiggle function or the tool is vibrated at a high frequency. Some of the reported benefits are highlighted in Table 1.

**Table 1.** Tabulated summary of recent efforts made in the vibration-assisted machining (VAM) of titanium alloys and steel.

#### 2.2.4. Surface Defect Machining

It has been demonstrated that the surface defect machining (SDM) method harnesses the combined advantages of both porosity machining [74] and pulse laser pre-treatment machining [75], as shown in Figure 5a, by machining a workpiece initially by generating surface defects at depths less than the uncut chip thickness through either mechanical means and/or thermal means followed by a routine conventional machining operation. SDM enables ease of deformation by shearing the material at a reduced input energy [76,77]. Also, due to the large proportion of stress concentration in the cutting zone—rather than the sub-surface—a reduction in the associated residual stresses on the machined surface is enabled.

(**b**)

**Figure 5.** (**a**): Schematic diagram indicating the differences between the mode of deformation during conventional machining and surface defect machining (SDM) observed through an FEA (Finite Element Analysis) simulation of hard steel and MD (Molecular Dinamics) simulation of silicon carbide, respectively; (**b**) effect of providing nanogrooves on the tool [78]. (**a**) Surface defect machining. (**b**) Providing nanogrooves on the tool.

Figure 5b shows the provision of the structured surfaces on the rake face of the cutting tool. It has been implied particularly that when these structured surfaces are fabricated at a direction of 90◦ to the cut surface, they are helpful in reducing the extent of friction at the tool–chip interface. In contrast, probing these structures along the direction of cutting is rather disadvantageous as it destroys the integrity of the edge of the tool. Overall, a summary of these recent developments in this area of machining, as an add-on to the existing machines to improve the workpiece's machinability, is summarised in Table 2.

#### *2.3. Protective Coatings*

Several types of failure mechanisms, such as delamination, abrasion, oxidation, diffusion, etc., result from the surface of the cutting tool. Failure occurs due to the interactions at the interface between tool and workpiece or tool and ambient medium, respectively. This leads to the conclusion that it is possible to protect a tool by surface treatment, creating some additional interface or coating with individually designed features. Figure 6 shows the distribution of uncoated and coated cemented carbide cutting tools, where the growth of the latter is clearly visible in a timeline perspective.

#### **Figure 6.** World market of cemented carbide cutting tools. Data from Dedalus Consulting, taken from [7].

Protective coatings may significantly increase the lifetime of tools and thus reduce the consumption of CRM content in bulk materials. Furthermore, most modern coatings are CRM-free. The evolution of the protective effect of coatings by the extension of tools' lifetimes and a comparison of coatings with available data from selected published works are shown in Figure 7. It is seen that coated cutting tools may have an extra lifetime of 200%–500% or more at the same cutting velocities. This can also lead to an increase in operational velocities (by 50% to 150%) with the same lifetime of cutting tools [93]. Coatings improve the endurance-related properties, and eventual wear failures may be dealt with through the wise application of the techniques outlined in Section 4.

**Figure 7.** Comparison of enhancement factors of coated tools over uncoated tools (based on their lifetime). Analysed data are taken from selected published works [94–98]. The results strongly depend on operational conditions, such as the cutting method and speed, the workpiece material, and the thickness of the protective coating.

Each application area requires specific properties of protective coatings to reach maximum performance. The extreme conditions of service of cutting tools mean that working parts face high wear, high pressure, elevated temperatures caused by high machining speeds, oxidation, and corrosion from lubricants or cooling agents. The main criteria of the evaluation of protective coatings are summarised in Figure 8, even though a compromise among the different prerequisites must be achieved. Usually, the most successful in one of the criteria may perform poorly in many of the others. The enhancement of all necessary properties is still a challenge for modern materials science and a strong driving force for computational screening modelling efforts worldwide (see Section 4).

**Figure 8.** Criteria required for the success of the protective coatings of cutting tools.

Based on principal properties and functionality, protective coatings may be grouped mainly into hard coatings, coatings with enhanced thermal stability, coatings with high oxidation and corrosion resistance and thermal barrier coatings. High hardness is essential for cutting tool coatings. In addition to the usually extended tool lifetime, it helps to achieve a smooth high-quality surface and shape of machined parts. However, hard materials are often brittle and prone to cracking, which is why it is crucial for protective coatings to have both high hardness and toughness. Only a mixture of several parameters leads to high wear performance and the long lifetime of cutting tools under extreme conditions. In the next section, a brief description of the main properties of the hard coatings mainly used to coat WC–Co tools and other tool materials is reported.

#### 2.3.1. Diamond and Diamond-Like Carbon (DLC) Coatings

Diamond is the intrinsically hardest material because of its strong nonpolar covalent C-C bonds (sp3 bonding) and short bond length. It belongs to the class of ultrahard materials, with H = 70–100 GPa [99]. Diamond exhibits the highest room temperature thermal conductivity of ∼20 W/cm·K and an extremely low coefficient of thermal expansion of ∼0.8 × 10−6 K−1 at 300 K. At room temperature, it is inert to attack from acids and alkalis, and it is resistant to thermal shock. The topography of the diamond surface can be very different depending on the extent of polishing and orientation of the crystallographic planes, varying from an extremely smooth surface with a friction coefficient as low as 0.1 in air to a very rough surface with protruding edges [100]. Because of these properties, diamond has tremendous application in the field of tribology, especially as a protective coating for cutting tools. Polycrystalline diamond films are best synthesized by chemical vapour deposition (CVD) techniques (but not limited to this) [101,102]; the most common are hot filament-assisted CVD (HFCVD), direct current plasma-assisted CVD (DC PACVD), microwave plasma CVD (MPCVD) and combustion flame-assisted CVD (CFACVD) [100]. However, most of the metals and ceramics have a much higher thermal expansion coefficient than the very low one for diamond coatings. This often may cause residual stress and the further spallation of the coatings, thus limiting the number of workpiece materials for diamond-coated cutting tools. Tungsten carbide (WC) has the closest coefficient of thermal expansion to that of diamond, which allows diamond film deposition on WC substrates in almost stress-free conditions. Nevertheless, the diffusion of cobalt contained in WC tools promotes higher graphitisation at the diamond–carbide interface, which induces coating delamination during machining. To reduce such effects and improve adhesion, the deposition of an interlayer with a specific composition, dopants, the multilayer architecture of coatings, surface etching and other surface treatment techniques may be applied [103–108]. This should also help in the case of the deposition of a diamond coating on surfaces different from WC, with a higher coefficient of thermal expansion.

Diamond-like carbon (DLC), as is clear from the name, exhibits some of the properties intrinsic to diamond. Among them, the most important factors for application in cutting tools are its hardness at a super- and ultrahard level, high wear resistance and low coefficient of friction (∼0.1). Most DLC films are structurally amorphous and can be synthesised by plasma-based PVD and CVD methods. Deposition method and the type of carbon source substantially influence the structural chemistry of the resulting films, which leads to large variations in their properties. Since the 1970s, DLC films have been well developed and discussed [109–111]. Structurally, they are made of sp2- and sp3-bonded carbon atoms and may be classified depending on their structure and hydrogen content as well as the presence of other dopants. DLC coatings with a lower hydrogen percentage (≤40%) and higher proportion of sp3 bonding in the structure have a higher hardness. Thereby, hydrogen-free tetrahedral amorphous carbon (ta-C) with the highest sp3 content (80%–88%) has an ultrahardness of 80 GPa. DLC films are not only hard but also smooth; furthermore, their tribological properties may be manipulated easily by introducing dopants such as nitrogen, silicon, etc. [112]. They also help to achieve high-quality shapes of the machined surface or cutting edge which are clean, smooth, and without chipping or rounding (see Figure 9). However, it is quite difficult to deposit thick coatings (of 2 μm and more) because of delamination due to the internal stress; in addition, there are limitations on the substrate choice with adherent film. Moreover, such coatings are prone to graphitise at temperatures above 300 ◦C with a

further decrease of their hardness [100]. On the other hand, when the outer layer of DLC is turned to graphite when the tool is in service, both the friction and wear rates should decrease [113].

**Figure 9.** Scanning electron microscope (SEM) micrographs of the rake and flank face after the dry milling test for an AlCu2.5Si18 alloy [109].

#### 2.3.2. Transition Metal Compounds

Transition metal nitrides (TMN), carbides (TMC) and borides (TMB) are largely employed as hard protective coatings in the cutting and forming tool industry. They attract interest due to their exceptional properties, such as their high hardness, chemical inertness, electronic properties, high melting point and thermal stability under harsh environments (oxidation, radiation, etc.). Such properties are mainly due to the variety of chemical bonding [114–116]. Based on the prevailing bonding type in TM nitrides, carbides and borides, Holleck [117] arranged their properties from low to high levels; some of these are shown in Figure 10. Depending on the priority properties and the suitable deposition technique, the coating which fits an application best can thus be chosen.

**Figure 10.** Properties of transition metal nitrides, carbides and borides [117].

The first industrial CVD-coated tool coating on cemented carbide was TiC in 1969, and in 1980 TiN became the first PVD coating. Lower temperatures in the PVD process made the deposition of coatings on steel tools possible [7]. Since the 1970s–80s, and up to the present, TiN, TiC, TiB2, CrN and ZrN have been the most frequently used binary coatings.

TiN has been the most widely studied TMN protective coating and has been in wide use since the late 1960s. However, it has some limitations and barely overcomes the modern challenges of thermal stability and oxidation resistance. At temperature above 500 ◦C, in fact, an oxide layer may be formed on the surface, which develops stress in the coating, which is high enough to damage or destroy the protective layer.

#### 2.3.3. Multi-Elemental Compounds and High-Entropy Alloy Protective Coatings

To overcome the modern challenges related to performance in extreme conditions, new, more complex compositions of coatings, such as ternary and quaternary compounds, have been proposed, which exhibit some superior and specific properties (e.g., thermal stability and oxidation resistance). An important example of this are Al-containing TMN ternary compounds of the TMxAl1−xN type, where TM is a transition metal, such as Ti, Cr, Zr, Nb, Hf, Ta, V, etc. The most popular are Ti-Al-N and Cr-Al-N systems, which have remained the "state-of-the art" coatings for a long time and are well-discussed in research and review papers [118–120]. Their exceptional hardness and oxidation resistance are mainly caused by the supersaturated solid-solution of hexagonal B4-structured AlN in the cubic B1 structure of TiN, with the resulting large volume mismatch, elastic strain energy and solid solution strengthening [119]. The mechanical and chemical properties of TMxAl1−xN coatings usually improve with the increase of the Al fraction, but only up to a certain critical Al content (usually around 40%–50%). The formation of Al, Ti or Cr oxides increases the oxidation resistance at elevated temperatures. Moreover, metastable supersaturated films such as TMxAl1−xN systems exhibit the phenomenon of age-hardening caused by decomposition with annealing temperature or time, leading to an increase in hardness (which is more significant in the Ti-Al-N system than in Cr-Al-N) [119]. This particular feature may have a high potential for the efficient enhancement of tool lifetimes.

Recently, the groups of T. Polcar and A. Cavaleiro performed studies on the addition of Cr to the Ti-Al-N system and its influence on thermal resistance, oxidation stability, tribological and cutting performance by the deposition of quaternary Ti-Al-Cr-N coatings [121,122]. A coating which demonstrates relatively low wear resistance at room temperature (Ti-Al-Cr-N system) in comparison to the other (in this case, the Ti-Al-N system) can exhibit much higher wear performance at elevated temperatures (650 ◦C) corresponding to realistic conditions for cutting tools' working temperatures.