The commercial ball-milling-crushed Fe78Si9B13 amorphous powders (China Amorphous Technology CO., LTD, Foshan, China) were sieved into three groups with different particle size ranges, and respectively named as P1, P2 and P3 powders for particle size distributions of 140–170 mesh (90–110 μm), 170–270 mesh (53–90 μm) and 270–400 mesh (38–53 μm). The sodium silicate and the talcum powder were employed as the coating and lubricating agents, respectively.

During a typical fabrication procedure of the AMPCs, the Fe78Si9B13 powders, consisting of 20 wt.% P1, 60 wt.% P2, and 20 wt.% P3, were first uniformly mixed with 2 wt.% of sodium silicate and 0.6 wt.% of talcum powder. The sodium silicate was sufficiently dissolved with the help of ethyl alcohol (4 wt.%) as the co-solvents. And then, the mixture was molded into green compacts with the dimensions of 26.9 mm in outer diameter, 14.8 mm in inner diameter and 11.2 mm in thickness (Φ26.9 mm × Φ14.8 mm × 11.2 mm) by cold pressing under a pressure of 2.05 GPa. Finally, the compacted cores were orderly dried at 393 K for 120 min, impregnated in dilute sodium silicate under vacuum condition, and annealed at 730 K for 30 min. In this work, three groups of powder core samples were prepared under various particle size distributions, compaction pressures, and coating agent contents. The particle size distribution of the Fe78Si9B13 powders was tested by a laser particle size analyzer (LPSA, BT-9300ST, Bettersize Instruments Ltd., Dandong, China). The microstructures and phase structures of the MPCs were examined by a scanning electron microscope (SEM, Phenom ProX, Phenom-World BV, Eindhoven, The Netherlands) and X-ray diffraction (XRD, Rigaku MiniFlex600X, Rigaku Corporation, Tokyo, Japan) with Cu Kα radiation, respectively. The MPCs' effective permeability, *μe*, the frequency characteristic (*μe*-*f*) when the frequency was the range of 20–100 kHz, within a constant voltage of 0.3 V, and the DC-bias performance were measured by an LCR meter (Wayne Kerr 3260B, Wayne Kerr Electronics, West Sussex, UK). The core losses were measured by a core loss measuring instrument (MATS-2010M, Linkjoin, Loudi, China) at 50 kHz and 100 kHz with a maximum magnetic flux density of 0.1 T, respectively. Among these, the cores were winded in turns of 30 with copper wires of Φ0.6 mm.

#### **3. Results and Discussions**

#### *3.1. Microstructure of Amorphous Powders*

Figure 1 displays the representative XRD pattern of the ball-milling-crushed Fe78Si9B13 amorphous powders. A broad diffuse peak at around 2θ = 45◦, without any distinct diffraction peaks, could only be seen for the powders, which validated the formation of an amorphous structure. The inset of Figure 1 shows the morphology of the ball-milling-crushed Fe78Si9B13 amorphous powders with the particle size ranging up to 125 μm. The powders were in the shape of a polygon with smooth edges, which was beneficial to the uniform insulation coating, thereby reducing the eddy current loss between the powders.

**Figure 1.** X-ray diffraction (XRD) patterns of the ball-milling-crushed Fe78Si9B13 amorphous powders, the inset shows the scanning electron microscope (SEM) micrograph of the powders

#### *3.2. Particle Size Distribution*

Table 1 lists the characterization of particle size distribution, green compact density, *ρ*, effective permeability, *μe*, and core loss, *Pc*, for the AMPCs with different matching ratios (mass percent) of P1/P2/P3%. In this table, D10, D25, D50, D75, and D90 represent the specific particle size at which the cumulative frequency showed 10%, 25%, 50%, 75%, and 90%, respectively. The green compact density of the AMPCs was analyzed by the principle of Archimedes. Figure 2 displays the morphology of the AMPCs with different matching ratios. Figure 3 displays the matching ratio dependence of the frequency characteristic and DC-bias performance. We observed the following points. (1) The attenuation rate of *μe*, *Rμ*, in the frequency band of 20–100 kHz for the AMPCs was 0.92–0.98%. It was indicated that the AMPCs with different matching ratios, given in Table 1, exhibited good constant magnetic properties with the frequency. (2) The circumferential crack was observed at the outer surface of the AMPCs when the mass ratio of the P1 powders was greater than 30 wt.%. This could be due to the increasing number of large pores induced by the increase of the coarse particles (P1 powders with the particle size distribution of 140–170 mesh), which was detrimental to the formability enhancement of the AMPCs [28]. (3) The AMPCs' green compact density, *ρ*, the effective permeability, *μe*, and the

residual effective permeability at the applied magnetizing field of 7.96 kA/m<EMAIL_ADDRESS>kA/m, increased first and then decreased with the increasing size of D50, and exhibited a maximum *ρ* of 5.66 g/cm3 and a maximum *μe* of 65.54, as the mass ratios of the P1%/P2%/P3% powders were 20:70:10 and 30:60:10, respectively. The AMPCs' *μe* can be calculated using the following equation:

$$
\mu\_{\mathfrak{e}} = \frac{L\_{\mathfrak{s}} l\_{\mathfrak{e}}}{\mu\_0 \mathcal{N}^2 A\_{\mathfrak{e}}},
\tag{1}
$$

where *Ls*, *Ae* and *le* are the inductance, the cross-section area, and the mean flux density path length of the ring sample cores, respectively. *N* is the total number of coil turns and *μ0* is the free space permeability.

It was seen that the magnitude of the *μe* strongly depended on the number and size of pores in the AMPCs according to the *le* in Equation (1). For the AMPCs of numbers 3, 2, and 1, as the mass ratio of fine powders (P3) increased from 10 wt.% to 30 wt.%, the interface between the magnetic particles and the specific surface area of the powders subsequently increased, followed by an increasing number of the pores. Thus, the *μe* and *ρ* gradually decrease from 64.54 to 53.10, and from 5.66 g/cm3 to 5.55 g/cm3. For the AMPCs of numbers 3, 4, and 5, the ratio of the coarse powders (P1) increased from 20 wt.% to 40 wt.%, which would lead to the increasing number of large pores in the AMPCs [28]. Accordingly, the *ρ* gradually decreased from 5.66 g/cm3 to 5.52 g/cm3. The *μe*, however, increased first and then decreased. A small increment of 1.08 for *μe* was observed as the mass ratio of P1%/P2%/P3% changed from 20:70:10 to 30:60:10, which could be attributed to the increase of the average particle size for AMPCs [29].

(4) It is known that the eddy current loss is dominant for the core loss in the high-frequency range. The eddy current loss is inversely proportional to the electrical resistivity but proportional to the square of particle size [30]. It was seen that the *Pc* gradually increased from 315.2 mW/cm3 to 409.7 mW/cm3, at 50 kHz for *Bm* = 0.1 T, as the size of D50 increased from 66.5 μm to 90.2 μm. Moreover, it was also shown that a small increment of 24.31 mW/cm3 for the *Pc* was observed as the ratio of the P3 powders decreased from 30 wt.% to 10 wt.%, with a constant ratio of 20 wt.% for the P1 powder. Based on the discussion above, the matching ratio of 20:70:10 for P1%/P2%/P3% was adopted as a reasonable process parameter for the subsequent fabrication of the AMPCs, because of its higher *μe*, relatively lower *Pc*, and being easier to shape, compared with the other matching ratios.

**Table 1.** Characterization of particle size distribution, green compact density, effective permeability, and core loss for the amorphous magnetic powder cores (AMPCs) with different matching ratios of P1%/P2%/P3%.

**Figure 2.** Morphology of the amorphous magnetic powder cores (AMPCs) with different matching ratios of P1%/P2%/P3%.

**Figure 3.** Dependence of matching ratio on frequency characteristic (**a**) and direct current (DC) bias performance (**b**).

#### *3.3. Molding Pressure*

Figure 4 shows the SEM micrographs of AMPCs' cross-sections under three different molding pressures of 1.95 GPa, 2.15 GPa, and 2.35 GPa, wherein the matching ratio of P1%/P2%/P3% for the AMPCs was 20:70:10. Figure 5 displays the dependence of the molding pressure on the compact density, the effective permeability, the frequency characteristic, the DC-bias performance, and the core loss for the AMPCs. It was observed that the microstructure of the AMPCs' cross-section consisted of coated powders and pores. The pores showed a three-dimensional network. With the increase of the molding pressure, the pore volume and its ratio gradually reduced. Accordingly, the AMPCs' green compact density, *ρ*, the effective permeability, *μe*, and the residual effective permeability at an applied field of 7.96 kA/m<EMAIL_ADDRESS>kA/m, increased from 5.63 g/cm3 to 5.86 g/cm3, from 62.37 to 73.91, and from 42.95 to 46.53, as the molding pressure increased from 1.95 GPa to 2.55 GPa, respectively. The MPCs' *Pc* decreased first and then increased with the increasing molding pressure, and exhibited a minimum value of 301.92 mW/cm3 (at 50 kHz for *Bm* = 0.1 T) at a molding pressure of 2.45 GPa. The attenuation ratio of *μe* in the frequency band of 20–100 kHz, *Rμ*, slowly went down from 1.02% to 0.85% with the increasing molding pressure. This was beneficial to the frequency characteristic improvement of the AMPCs.

Moreover, it was seen from Figure 5a,d that both the *μe* and *Pc* rapidly increased as the molding pressure surpassed 2.45 GPa. This could be contributed to the damage of the coated insulation material and the increasing residual stress caused by greater molding pressure. When the molding pressure was lower than 2.45 GPa, the increasing process of *μe* could be divided into two different stages: The initial fast increasing stage (1.95–2.25 GPa), and the subsequent stably increasing stage (2.25–2.45 GPa). It was also seen from Figure 5d that a small increment of 4.70 mW/cm3 for the core loss was observed as the molding pressure decreased from 2.45 GPa to 2.35 GPa. Thus, it was concluded that the molding pressure range of 2.35–2.45 GPa was reasonable for the AMPCs' fabrication.

**Figure 4.** SEM micrographs of AMPCs' cross-sections under different molding pressures, (**a**) 1.95 GPa, (**b**) 2.15 GPa, and (**c**) 2.35 GPa.

**Figure 5.** Dependence of the molding pressure on the compact density, the effective permeability (**a**), the frequency characteristic (**b**), the DC-bias performance (**c**), and the core loss (**d**) for the AMPCs.

#### *3.4. Insulating Material Content*

Figure 6 shows the SEM micrographs of the coated magnetic powders with three different coating agent (sodium silicate) contents of 0.5 wt.%, 1.0 wt.% and 2.0 wt.%. Figure 7 displays the dependence of the sodium silicate content on the compact density, the effective permeability, the frequency characteristic, the DC-bias performance, and the core loss for the AMPCs with a P1%/P2%/P3% ratio of 20:70:10 and a molding pressure of 2.45 GPa. It was seen that the increase of the coating agent content was beneficial to obtaining an even electro-insulating layer on the surface of the magnetic powders. The attenuation ratio of *μe*, *Rμ*, in the frequency band of 20–100 kHz was in the range of 0.83%–0.91%. That meant that the AMPCs with different sodium silicate contents all exhibited good frequency stabilities of effective permeability. The AMPCs' core loss, *Pc*, as shown in the Figure 7d, gradually decreased from 590.54 mW/cm3 to 248.93 mW/cm3, at 50 kHz for *Bm* = 0.1 T, with an increasing content of sodium silicate from 0.5 wt.% to 3.5 wt.%, due to the improving insulation effect for the AMPCs.

As seen from Figure 7a–c, the AMPCs' *ρ*, *μe*, and<EMAIL_ADDRESS>kA/m initially increased, but was followed by a decline with the increasing content of sodium silicate. Within the content range of 0.5 wt.%–1.5 wt.%, the *ρ* increased with the increase of the sodium silicate content. This was due to the lubrication film on the particle surface, generated by the coating agent, which can enhance the mobility of the particles during the compacting process [27,29]. For the *μe*, too low content of sodium silicate (0.5 wt.%, seen in Figure 6a) would result in the increased number of pores, due to the poor bonding effect for the particles in the AMPCs. Thus, a *μe* of 70.49 was observed. As the sodium silicate content increased to 1.0 wt.%, the uneven coating layer on the powder surface—as shown in the Figure 6b—would cause the poor insulation for particles in the AMPCs, and in turn, the *μe* climbed to 81.87. With the improved insulation effect for particles, the *μe* decreased to 77.01 as the sodium silicate content increased to 1.5 wt.%. Moreover, it was also seen that both the *ρ* and *μe* gradually decreased when the sodium silicate content was greater than 1.5 wt.%, due to the increasing ratio of

non-magnetic material (coating agent) in the AMPCs. Therefore, the sodium silicate additive amount of 1.5 wt.% was proposed for the AMPCs' fabrication with a P1%/P2%/P3% mass ratio of 20:70:10 and a molding pressure of 2.45 GPa. And then, an excellent comprehensive soft magnetic property could be obtained as follows: *μe* = 77.01, 0.85% < *Rμ*, *μe* (at a bias field of 7.96 kA/m) = 49.57, *Pc* (at 50 kHz for *Bm* = 0.1 T) = 305.23 mW/cm3, and *Pc* (at 100 kHz for *Bm* = 0.1 T) = 650.46 mW/cm3.

**Figure 6.** SEM micrographs of the coated magnetic powders with three different sodium silicate contents of 0.5 wt.% (a), 1.0 wt.% (b) and 2.0 wt.% (c).

**Figure 7.** Dependence of the sodium silicate content on the compact density, the effective permeability (**a**), the frequency characteristic (**b**), the DC-bias performance (**c**), and the core loss (**d**) for the AMPCs.

#### **4. Conclusions**

The comprehensive soft magnetic properties of the Fe78Si9B13 AMPCs have been enhanced by the coordinated adjustments of the particle size distribution, the compacting pressure, and the coating agent content in the industrial condition. The optimal magnetic properties of the Fe78Si9B13 AMPCs can be obtained when the technological parameters are as follows: The matching ratio of P1%/P2%/P3% is 20:70:10 in mass percent, the molding pressure is in the range of 2.35–2.45 GPa, and the additive

amount of sodium silicate is 1.5 wt.%. As compared to the original scheme (the P1%/P2%/P3% mass ratio in 20:60:20, the molding pressure in 2.05 GPa, and the sodium silicate content in 2.0 wt.%), the AMPCs' compact density, *ρ*, the effective permeability, *μe*, and the residual effective permeability at 7.96 kA/m<EMAIL_ADDRESS>kA/m, increased from 5.61 g/cm3 to 5.86 g/cm3, from 58.13 to 77.01, and from 40.36 to 49.57, respectively, while the core loss reduced from 380.85 mW/cm3 to 335.23 mW/cm3 (at 50 kHz for *Bm* = 0.1 T).

**Author Contributions:** H.S. and W.C. conceived and designed the experiments; C.W. and J.L. performed the experiments and analyzed the data; H.S. wrote the paper.

**Funding:** This research was funded by the National Natural Science Foundation of China (Grant No. 51704078), the Province Natural Science Fund of Guangdong (Grant No. 2017A030313312), the university key platform funding projects of Guangdong education department (Grant No. gg041002), the engineering technology research center project of Foshan City (Grant No.<PHONE_NUMBER>8) and the talent research start-up program of Foshan University (Grant No. gg040942).

**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*

### **Fabrication of Functionally Graded Materials Using Aluminum Alloys via Hot Extrusion**

**Dasom Kim 1, Kwangjae Park 1, Minwoo Chang 2, Sungwook Joo 3, Sanghwui Hong 3, Seungchan Cho 4 and Hansang Kwon 1,5,\***

Received: 20 December 2018; Accepted: 7 February 2019; Published: 11 February 2019

**Abstract:** In this study, we have attempted to manufacture functionally graded materials (FGMs) using aluminum alloys 3003 and 6063 via a hot extrusion process to realize multifunctionality through achieving high strength and low weight. The FGMs were fabricated using Al3003 powder and Al6063 bulk to improve the interfacial properties. Particle size analysis and X-ray fluorescence of the Al3003 powder were used to analyze the composition of general Al3003; microstructure analysis revealed improved hardness with almost no defects, such as cracks at the interface between the two materials. The experimentally determined tensile strength of the composite was observed to be higher than the theoretical value calculated using the rule of mixtures; the strengthening mechanisms considered for the calculations were grain size reduction and precipitation hardening. In particular, we attempted to predict the strengthening effect resulting from the fine grain size of the powder and grain size reduction due to the extrusion process using the Hall–Petch equation. The Kelly–Tyson equation was also used to calculate the theoretical strength in the presence of the strengthening phases. Based on these results, it was confirmed that FGMs can be successfully produced using the hot extrusion process.

**Keywords:** functionally graded materials (FGMs); aluminum alloy; hot extrusion; Hall–Petch equation; Kelly–Tyson equation

#### **1. Introduction**

Lightweight materials have been extensively developed for use in transport machineries, especially in aircrafts, automobiles, and marine vehicles. In particular, aluminum is generally used to realize lightweight in materials. Therefore, aluminum has been used in the form of an alloy by adding other elements by many researchers, such as Al-Stainless steel (SUS) composite and Al-carbon nanotubes (CNT) composite [1–3].

Recently, multifunctionality, where diverse functions are realized in the same material, has been increasingly desired. In this context, the concept of functionally graded materials (FGMs) has been suggested as a method to realize multifunctionality. FGMs are materials in which a gradient in property is created by changing the composition continuously.

Miyamoto et al. introduced the concept of FGMs [4]. This concept facilitates the improvement of the quality of the interfaces in a heterogeneous material to prevent cracks, which arise due to the residual stress resulting from the difference in the properties of the different components in the heterogeneous material.

The FGMs mitigate the differences in the properties between the two materials by forming interlayers; however, this is difficult to control because the layers formed at the interface must be thin and uniform. Besides, cracks are sometimes also observed in the FGMs.

To address these problems, researchers have attempted new fabrication methods, such as centrifugal casting, semi-solid forming under magnetic field, and spark plasma sintering (SPS) [5–7], while extrusion processes such as co-extrusion, indirection, and constant extrusion have been used as a fabrication method for over 30 years [8,9]. Fan et al. attempted to join Al1060 and Al6063 by welding following die extrusion [10]. However, there were many cracks at the interface due to the difficulty in controlling the stress. Kwon et al. tried a hot extrusion process to fabricate an Al-CNT composite [11]. The microstructure was observed to be elongated in a direction parallel to the direction of the extrusion, with no cracks and pores. Furthermore, although the composite includes only one volume, the mechanical properties were significantly improved.

This study attempted to fabricate FGMs via a hot extrusion process. The Al3003/Al6063 FGM fabricated in this study has the shape of a rod, the center axis of which consists of Al3003 and serves as a skeleton for the material; this center is surrounded by Al6063. While Al6063 is an aluminum-based alloy that has high strength and hardness to protect the material from external shock, the Al3003 alloy has excellent ductility, which gives flexibility to the material itself, thereby enhancing workability. Thus, the FGM is suitable for use as a wire material for an industrial electric pole and can also be used as a skeleton for steel structures or automobile structural materials.

For dissimilar materials to be well bonded, the surface area of the contact between the two materials must be large. For example, if the surface is rough and the bulk is in contact with a rugged bulk surface, the area of the contact will be larger. Uehara and Sakurai confirmed the relation between the bonding strength and surface roughness of a joined part, which was controlled by changing the feed in a common mechanical milling machine [12]. They observed that the tensile strength was higher when the surface roughness was increased to ~3–6 μm than when the surface roughness was <1 μm.

Figure 1a shows a schematic diagram of the contact surface between two materials in the bulk form. Under these conditions, the contact surface comprises the area indicated by the straight line where the two materials come into contact. On the other hand, if one material is bulk and the other is a powder (Figure 1b), the contact surface area significantly increases because the powder has a much larger surface area than the bulk. Furthermore, Figure 1c shows a schematic diagram of the contact surface between two materials in powder form. It can be observed that the surface area of the contact is much greater, although it is difficult to handle such bonding experimentally. Therefore, we fabricated the FGMs, as shown in Figure 1b; and expected better interface properties compared to the case when two bulk materials are used.

**Figure 1.** Illustration of the interface between two dissimilar materials. (**a**) bulk–bulk interface, (**b**) bulk–powder interface, and (**c**) powder–powder interface.

To fabricate a sample, we performed canning followed by extrusion. Subsequently, the microstructure of the sample was observed using scanning electron microscopy (SEM), and the composition of the FGMs was analyzed using energy-dispersive spectrometry (EDS). In addition, we used the value of the crystallite size, which was calculated using the Scherrer equation, in the Hall–Petch equation to elucidate the relation between the grain size and strength.

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

#### *2.1. Sample Preparation*

First, canning was performed as shown in Figure 2. Al3003 powder (965 g; ECKA Granules, purity: 99.5%, particle size: 100 μm) was poured into a bulk Al6063 can (outer diameter: 100 mm and inner diameter: 60 mm) and subjected to a vertical force under a pressure of 5000 kgf. Subsequently, the material was extruded at a temperature of 468 ◦C, which is below the melting point of both Al3003 and Al6063, using an extrusion machine (SKM Co., Ltd., Incheon, Korea). The temperatures of the billet, container, and mold were 468 ◦C, 407 ◦C, and 468 ◦C, respectively. The ram speed was 2.49 mm/s and the extrusion ratio was 100.

**Figure 2.** Schematic view of the Al6063 bulk cover and can.

#### *2.2. Analysis of Raw Materials*

Particle size analysis (PSA) and X-ray fluorescence (XRF) were used to characterize the Al3003 powder. PSA was first conducted (BECKMAN COULTER, LS 13 320, Indianapolis, IN, USA) followed by XRF to analyze the constituent elements (SHIMADZU, XRF-1800, Tokyo, Japan).

#### *2.3. Characterization of the Samples*

First, we used the Archimedes method to measure the density of the FGMs using a densitometer (KERN, ABJ 120-4M, Balingen, Germany). After etching with 5% NaOH solution, SEM and EDS were employed for a detailed analysis of a sectional area of the samples in detail (TESCAN, VEGA II LSU, city, Czech Republic). Elemental mapping was conducted to determine the ingredient elements of the material (TESCAN, VEGA II LSU, Brno, Czech Republic).

Subsequently, mechanical properties such as hardness and tensile strength were measured using the Vickers hardness test (HM-101, Mitutoyo Corporation, Kanagawa, Japan) according to JIS B 7725 and ISO 6507-2 at five points from the core to the clad (three measurements per point) with a load of 0.3 kg. To measure the tensile stress and strain, a tensile test was conducted on two FGM samples: Al3003/Al6063 FGMs without heat treatment (T-0) and T-5 heat-treated Al3003/Al6063 FGMs (Kyungdo precision Co., Ltd., Siheung, Korea). The no. 4 tensile specimen was fabricated according to KS B 0801 T-5 (width: 6 mm). The treatment is the process that artificially elevates the temperature of the specimen to 150–200 ◦C, and then cools it [13]. The tensile test was conducted at 24 ◦C using a 68.65 kN press. Finally, we used the Hall–Petch equation to confirm the relationship between the grain size and tensile strength; the grain size was substituted by the crystallite size, which was calculated using the Scherrer equation and the XRD data (Rigaku, Ultima IV, Tokyo, Japan). Furthermore, the Kelly–Tyson equation was used in this study to calculate the theoretical yield stress of the FGMs.

#### **3. Results**

In general, the dispersion of particle size is an important characteristic of a particulate sample and has significant influence on the packing density and porosity of the products.

Figure 3 shows the PSA graph of Al3003 powder with various particle sizes ranging from 10 μm to 280 μm; the average particle size is 125 μm. The accumulate volume graph also shows that the d10, d50, and d90 data for the particle size distribution is 35 μm, 97 μm, and 156 μm, respectively.

**Figure 3.** (**a**) scanning electron microscope (SEM) image of Al3003 powder and (**b**) Typical Particle size analysis (PSA) graph (black squares) and accumulate PSA graph (outlined triangles) of the Al3003 powder used in this study.

When Al3003 powder is filled into the Al6063 bulk can, a high packing density (about 63%) can be achieved as the Al3003 powder has a significantly high degree of dispersion in particle size, which indicates that the density of the FGMs is high.

Table 1 lists the general chemical composition of Al3003 and Al6063. By comparison, Si (1.10%), Mn (0.96%), Fe (0.87%), Zn (0.23%), Cu (0.17%), and Al (96.66%), which is determined by XRF, are present in the Al3003 powder used in this study.

**Table 1.** General chemical composition of the components Al6063 and Al3003 [14,15].

We observe that the proportions of the components are approximately in the range of the typical chemical composition of the Al3003 powder (presented in Table 1). Consequently, Al3003 powder, with a mean particle size of 125 μm, was filled into a bulk can made of Al6063, and the final FGM was fabricated via hot extrusion.

Figure 4 shows the photograph of the final material, which is in the shape of a cylindrical rod. Figure 4b shows the cross-section of the FGM, where the inner darker area is composed of Al3003 and the outer brighter area is composed of Al6063. The inner diameter is around 1 mm, and the outer diameter is 8 mm. The density of the FGMs was also measured to confirm that there were nearly no flaws in the material manufactured via hot extrusion. The experimental density measured using the Archimedes' method was 2.6538 g/cm3 and the theoretical density calculated using the rule of mixtures was 2.7009 g/cm3. A significantly high relative density of 98.26% was obtained, which confirms the successful fabrication of a dense FGM. Furthermore, it was confirmed that the density of the Al3003 core was nearly 100%. This indicates that hot extrusion is a useful method to form an interface between powders and bulk materials, demonstrating that this process can be used for the fabrication of complex heterogeneous materials.

**Figure 4.** Al3003/Al6063 functionally graded materials (FGMs) fabricated via hot extrusion: (**a**) Photograph of the FGMs and (**b**) high-magnification image of the cross-section of Al3003/Al6063 FGMs.

Figure 5 shows the SEM images of the Al3003/Al6063 FGMs; microstructure analysis was performed to confirm whether the final FGM was fabricated without any cracks and faults. Figures 5d and 5f show the EDS results for Al6063 and Al3003.

By comparing the XRF results, where Mg and Si were determined to be the predominant components of Al6063 and Al3003, respectively, with the EDS result, we infer that Figure 5d corresponds to Al6063 and Figure 5f corresponds to Al3003.

There are no cracks, except for a few pores created via deep etching (Figure 5a) in the region containing Al6063. However, the Al3003 region consists of more pores, which was not observed before conducting the etching process, and moreover, particles were observed, as shown in Figure 5c. This could be because the etching process in Al3003 can be faster and more significant than in Al6063, since the Al3003 is in the form of a powder, which has a higher specific surface area, while Al6063 has higher corrosion resistance compared to Al3003. We also confirmed from the EDS results that the O content is higher in Al3003 than in Al6063. Hence, deep etching was required to observe the microstructure of Al6063, which has high corrosion resistivity. Therefore, these results show that oxidation has occurred more actively in Al3003 than in Al6063. Figure 5b shows the SEM micrograph of the interface between Al6063 and Al3003, demonstrating a distinct and well-formed interface without any flaws, such as cracks. Figure 5e shows a high-magnification micrograph, where no cracks can be observed, except for pores caused by the deep etching similar to Figure 5b.

**Figure 5.** SEM images (**a**–**c**,**e**) and EDS results (**d**,**f**). (**a**) Region constituted by Al6063, (**b**) interface between Al6063 and Al3003, (**c**) region constituted by Al3003, (**d**) EDS of Al6063, (**e**) high-magnification image of the interfacial area between Al6063 and Al3003, and (**f**) EDS of Al3003.

Figure 6 shows the SEM image of a longitudinal section of the Al3003/Al6063 FGM. Figure 6a shows the region containing Al3003, where the pores are larger and longer when the extrusion direction was subjected to pressure and was compressed, whereas the longitudinal section parallel to the direction of extrusion was stretched in the same direction. The interface between Al3003 and Al6063 was analyzed in detail using elemental mapping.

**Figure 6.** SEM images of the longitudinal section of the FGM. (**a**) Region containing Al3003, (**b**) region containing Al6063, (**c**) interface between Al3003 and Al6063, and (**d**) high-magnification image of the interface between Al3003 and Al6063.

Figure 7b shows the distribution of the relative concentrations of the different components at the interface between Al3003 and Al6063. The particles in the longitudinal section are observed to be relatively larger and have an elongated shape when compared with those observed in the cross-section. Therefore, we can consider that the grains of the FGMs are aligned in the direction of extrusion. Near the interface in the Al3003 area, we expect that the degree of oxidation caused by the deep etching is higher than in the other areas because of the higher O content and lower proportion of Al in this area.

**Figure 7.** Elemental mapping of the area near the interface between Al6063 and Al3003. (**a**) Picture of a cross-section of the sample, (**b**) SEM image and mapping of all the elements, and (**c**) mapping of the major elements.

In general, while etching is conducted, the grain boundaries are oxidized more than the grain interior. Similarly, the interface is more corrosive due to the disordered arrangement of the atoms. To analyze the mechanical properties, Vickers hardness test and tensile test were conducted.

Figure 8 shows the Vickers hardness values measured on the cross section and longitudinal section of the FGM. The overall hardness in the transverse cross-section is observed to be in the range of 50–53 HV, which is almost 1.5 times greater than that generally observed for both Al3003-T0 and Al6063-T0, as shown in Table 2, and the hardness values are nearly constant. Al6063 has good corrosion resistance, and furthermore, it is fabricated in the bulk form; therefore, it is only slightly oxidized. The high hardness value of 53.03 HV observed at the interface confirms that the Al3003/Al6063 FGM has superior interface properties. Moreover, in the cross-section, although the hardness of the individual parts is slightly different, this difference is not large; therefore, the hardness is uniform.

However, the measured value of the hardness of the longitudinal section is relatively lower when compared to that of the transverse cross-section. This could be because the particle size is further reduced in the cross-section where the grains are subjected to a compressive pressure from all directions, which is less than that in the longitudinal cross-section. While joining two different materials, the properties of the joint play an important role—if the two materials are well bonded, the interfacial properties and the physical properties of the final product are improved.

In contrast, the longitudinal section is elongated in the direction of extrusion and the particles are compressed in a direction vertical to the direction of extrusion.

**Table 2.** General specifications of mechanical properties of Al6063 and Al3003 [16].

**Figure 8.** Vickers hardness graph measured on the transverse cross-section and longitudinal sections of the FGM at five regions from the center (Al3003) to Al6063, including the interface between Al3003 and Al6063.

Figure 9 shows the tensile stress–strain graph of the Al3003/Al6063 FGMs (T-0 and T-5). Al3003 and Al6063 have an area ratio of 18:82 at the neck of the tensile specimen. In the case of the T-0 Al3003/Al6063 FGM, the experimental tensile stress is improved to 152.19 MPa, as listed in Table 3, and in the case of the T-5 Al3003/Al6063 FGM, the general tensile stress is improved to 237.29 MPa, which is more than twice the generally observed average tensile stress values for Al3003 and Al6063, as shown in Table 2. The tensile stress of the material satisfies the T6 specification reported by Prillhofer et al. [17]. The elongation of the FGMs without heat treatment is more than 25%, which can solve the formability problem of the general aluminum 6000 series studied by Panigrahi et al. [18].

Furthermore, in the case of T-5 Al3003/Al6063 FGMs, the elongation, which is more than 10%, satisfies the abovementioned T6 specification. To analyze the strengthening mechanism, we considered two possible mechanisms: precipitation hardening and reduction of grain size. In the first mechanism, the precipitates present inside the matrix of the metal play a role in capturing dislocations, which increases the strength and hardness of the material.

Generally, Al alloys have precipitates and in the case of heat-treated alloys, the precipitates are formed and strengthened during the heat treatment. Pokov'a et al. showed that the α-Al12-15 (Fe, Mn)3 Si1-2 phase and Al6 (Fe, Mn) phase are generally present in Al3003 [19], and Camero et al. reported that the precipitates in Al6063 such as Mg2Si are generated mainly during heat treatment [20]. These precipitates serve to reinforce the material through a pinning effect. Furthermore, even if the amount and size of the precipitates are the same, the area of the FGMs is reduced by the extrusion process, which indicates that the proportion of the precipitate per unit area increases. Consequently, the movement of the dislocations is significantly inhibited; the strengthening of a material can occur through this mechanism. In the second mechanism, the materials are strengthened due to the reduction in grain size. As the crystal grain size of the material decreases, the grain boundary area increases in general, which disturbs the dislocation migration. Some researchers attempted to calculate the strength of the extruded materials using the Hall–Petch equation. For instance, Uematsu et al. showed that the grain size can be refined by controlling the extrusion conditions (extrusion ratio, working temperature), and showed that the Hall–Petch equation is valid for calculating the strength of the extruded alloys [21].

**Figure 9.** Stress and strain graph of Al3003/Al6063 FGMs (T-0 and T-5).

Furthermore, Kim et al. attempted to control the grain size using equal channel angular pressing, which is an extrusion method, and used the Hall–Petch equation to calculate the yield strength [22]. Thus, it can be concluded that the use of the Hall–Petch equation to calculate the extruded FGMs is valid, and hence, the equation was applied in this study.

To use the Hall–Petch equation, we first calculated the crystallite size by applying the Scherrer equation (Equation (1)). Some researchers applied the crystallite size to the grain size of the Hall-Petch equation, and proved that it is considerably valuable [23]. The Scherrer equation is expressed as follows [24]:

$$
\pi = K\lambda / \beta \cos \theta \tag{1}
$$

where *τ* is the mean size of the ordered (crystalline) domains, which may be smaller or equal to the grain size, and *K* is a dimensionless shape factor with a value close to unity. From the XRD data, the crystallite size of the Al3003 powder is determined to be 48.85 nm, while that of the Al3003/Al6063 FGM is determined to be 44.87 nm. It is thus confirmed that the grain size is decreased owing to the extrusion process and the result shows a possibility of strengthening due to the reduction in the grain size. The values of the crystallite sizes were substituted into the Hall–Petch equation to calculate the yield stress of Al3003 powder and Al3003/Al6063 FGM, as shown in Equation (2). The Hall–Petch equation is expressed as follows [21]:

$$
\sigma = \sigma\_0 + K\_y d^{-1/2} \tag{2}
$$

In this equation, *σ*0 is the materials constant for the stress in the dislocation movement, *Ky* is the Hall–Petch slope, which is a strengthening coefficient, and *d* is the average grain size. It can be observed from this equation that the yield strength in the y-direction increases as the average particle size, *d*, decreases. In the case of the T-0 Al3003 sample, this effect caused by the reduction in particle size is expected to be observed through improving the material properties via solid-solution strengthening, dispersion strengthening, and work hardening. The value of *σ*0 is approximately 58 MPa, while that of

s approximately *Ky* i96 MPa·μm1/2, and the yield strength was calculated by substituting the crystallite size obtained from the Scherrer equation into the Hall–Petch equation.

The calculated value of the yield strength before extrusion is 491.92 MPa and that after extrusion is 511.12 MPa, showing an increase of approximately 20 MPa after the extrusion process. The strength of the final material is increased by ~50 MPa in the sample that was not subjected to the heat treatment. Additionally, precipitation hardening also plays a role in increasing the strength. Therefore, it is considered that a combined effect of these factors increases the final strength by 50 MPa.

Notably, the calculated value is much higher than the experimentally measured value. The considerable difference between the two values is possibly because of using the crystallite size instead of the actual grain size for the calculations. For instance, in the case of a polycrystalline material, the crystallite size is generally smaller than the grain size. If the strength were calculated using the actual grain size, which is expected to be larger than the crystallite size, in the Hall–Petch equation, the result is likely to be similar to the experimental value.

However, notably, the crystallite size was reduced after the extrusion process, and the difference between the strength of the raw materials and extruded materials is similar to the difference between the theoretical values calculated using the Hall–Petch equation. Therefore, the Hall–Petch equation is considered to be effective for explaining the tendency of the reduction in the grain size via the extrusion process and the relatively increased value of the strength. However, a more specific analysis is required to demonstrate the exact strengthening mechanisms, as there are possibly other strengthening factors besides the two factors mentioned in this paper.

To analyze the interactions between the increase in strength owing to the strengthening factors and the decrease in strength owing to the limitations (e.g., difficulty in atomic-level bonding of two other materials), further investigations in addition to quantitative and qualitative analytical data including simulations are required.

$$
\sigma\_y = \sigma\_f V\_f \left( 1 - \frac{l\_c}{2l} \right) + \sigma\_m \left( 1 - V\_f \right) \tag{3}
$$

Moreover, Kelly and Tyson established the equation for the yield strength of FGMs, in which the matrix and fiber are evenly mixed and aligned in a particular direction with perfect interfacial bonding [25].

**Table 3.** Experimental mechanical properties of Al6063/Al3003 FGM (T-0, T-5).

The Kelly–Tyson equation shows that the overall strength is improved by a load transfer to a fiber when a load is applied to a composite material produced by mixing a matrix and fiber, which indicates that the fiber serves as an enhancer. This equation is used for complex composites. When the material is produced through an extrusion process, the grains are elongated in a direction parallel to the extrusion direction, and the material is compressed in the vertical direction.

A previous study confirmed that this phenomenon actually occurs through the micrographs of the longitudinal cross-section of the Al-CNT composite produced via an extrusion process [26].

As shown in Figure 5a, when the pores distributed in Al3003 are examined, the particles are observed to be arranged in the direction of extrusion and the circular shape of the pores is elongated in the direction of the extrusion. Based on this observation, it was concluded that the Kelly–Tyson equation can be applied for the fabricated material used in this study.

Therefore, the Kelly–Tyson equation was employed, assuming that one material plays the role of the fiber and the other one plays the role of the matrix, although the two are not mixed in our case. Assuming that Al3003 is the matrix and Al6063 is the fiber, in the case of the T-0 FGM, the calculated value of the strength is approximately 147 MPa, which is lower than the experimental value (152.19 MPa, as shown in Table 3). In the case of the heat-treated FGM T-5, the calculated value of the strength is expected to increase to approximately 231 MPa. This theoretical value is also lower than the experimental value for T-5 FGM (237.29 MPa, in Table 3). Thus, for both T-0 and T-5 FMGs, the experimental values are observed to be higher than the calculated values, although the difference is small (less than 5%). However, when the theoretical strength was calculated using the rule of mixtures equation, the values were 90.26 MPa (for T-0 FMG) and 183.32 MPa (for T-5 FMG). Thus, although the Al3003/Al6063 FGM fabricated in this study is not a perfectly mixed composite, it can be confirmed that the Kelly–Tyson equation can be applied to this system. Moreover, the higher experimental value indicates that there are additional factors, such as the reduction in particle size and precipitation hardening, which may play a significant role as described earlier.

The Kelly–Tyson equation considers strengthening effects only by fibers; therefore, this equation may not be the most appropriate equation in the present case. However, when considering all the different factors that affect the strengthening of the composites, this equation can be considered as most suitable.

#### **4. Conclusions**

The Al3003/Al6063 FGM was successfully fabricated via hot extrusion to realize multifunctionality including high strength and low weight.

The interface between the two materials is clear, with almost no cracks. It can be considered that the extruded FGM has excellent interface properties based on the high Vickers hardness of the interface. When the powder and bulk are in contact, the surface area of the contact could be larger than that between two bulks, so that the interface was improved.

The tensile strength of the final material subjected to T-5 heat treatment was observed to improve to twice that generally observed for Al3003 and Al6063, and the elongation was greater than 10%. From these results, it can be confirmed that the final product has a high strength and adequate elongation, which are the objectives of this study.

To analyze the strengthening mechanisms, i.e., precipitation hardening and grain refinement, strengthening via grain refinement was emphasized more in this study. Generally, the grain size of the powder is finer than the bulks and the presence of many grain boundaries hinders the grain growth. Therefore, the FGM fabricated using Al3003 powder could be strengthened by the fine grain size. Moreover, the grain size would be reduced by the extrusion process because the grains were subjected to mechanical stress. The Hall–Petch equation was used to explain the tendency for improved strength, which is appropriate for the explanation of the tendency of the increase in strength when the FGM was extruded; this was reported by earlier researchers, as well as observed in our study.

In addition, when the Kelly–Tyson equation was used assuming that Al6063 was a fiber, the theoretical results nearly corresponded to the experimental values. Therefore, the Kelly–Tyson equation could also be used to predict the strength of the extruded FGMs.

In summary, Al-alloy-based FGMs with lightweight and high strength properties were successfully fabricated via hot extrusion. Thus, it could be confirmed that hot extrusion is useful to fabricate FGMs with improved mechanical properties.

**Author Contributions:** D.K. and K.P. conducted the overall experiments; M.C., S.J., and S.H. fabricated materials via a hot extrusion process and conducted the heat treatment; the tensile test was conducted by S.C.; writing—original draft preparation, D.K.; writing—review and editing, H.K.; visualization, D.K.; supervision, H.K.; project administration, H.K.; and funding acquisition, H.K.

**Funding:** This research was funded by Ministry of SMEs and Startups (MSS).

**Acknowledgments:** This research was financially supported by the Ministry of SMEs and Startups (MSS), Korea, under the "Regional Specialized Industry Development Program" supervised by the Korea Institute for Advancement of Technology (KIAT), (2017\_R0006154).

**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* **Balancing Porosity and Mechanical Properties of Titanium Samples to Favor Cellular Growth against Bacteria**

**Ana Civantos 1,\*, Ana M. Beltrán 2,\* , Cristina Domínguez-Trujillo 2, Maria D. Garvi 3, Julián Lebrato 3 , Jose A. Rodríguez-Ortiz 2 , Francisco García-Moreno 4 , Juan V. Cauich-Rodriguez 5, Julio J. Guzman 2 and Yadir Torres 2**

Received: 6 September 2019; Accepted: 22 September 2019; Published: 24 September 2019

**Abstract:** Two main problems limit the success of titanium implants: bacterial infection, which restricts their osseointegration capacity; and the stiffness mismatch between the implant and the host cortical bone, which promotes bone resorption and risk of fracture. Porosity incorporation may reduce this difference in stiffness but compromise biomechanical behavior. In this work, the relationship between the microstructure (content, size, and shape of pores) and the antibacterial and cellular behavior of samples fabricated by the space-holder technique (50 vol % NH4HCO3 and three ranges of particle sizes) is established. Results are discussed in terms of the best biomechanical properties and biofunctional activity balance (cell biocompatibility and antibacterial behavior). All substrates achieved suitable cell biocompatibility of premioblast and osteoblast in adhesion and proliferation processes. It is worth to highlighting that samples fabricated with the 100–200 μm space-holder present better mechanical behavior—in terms of stiffness, microhardness, and yield strength—which make them a very suitable material to replace cortical bone tissues. Those results exposed the relationship between the surface properties and the race of bacteria and mammalian cells for the surface with the aim to promote cellular growth over bacteria.

**Keywords:** bone implant; porous titanium; cellular adhesion; bacteria colonization; osseointegration

#### **1. Introduction**

Longer life expectancy, traumas, and congenital diseases have led to an increase in the use of implants. Bone metal implants are usually made of commercially pure titanium (c.p. Ti) or based on Ti alloys owing to its biocompatibility and mechanical behavior [1]. C.p. Ti has a Young's Modulus of 100–110 GPa, which is higher than cortical bone (20–25 GPa), causing stress shielding phenomena and bone resorption [2]. This problem can be solved using porous material instead of fully-dense one [3–6]. However, in order to achieve a biomechanical balance (stiffness vs. mechanical strength) of the bone tissue to be replaced, an adequate porosity (in terms of size, shape, and distribution of the pores) is required. Up to 34 processing routes to fabricate porous materials have been already reported [7–10]. Among them, the use of space-holders stands out. This technique uses particles—such as salt [11,12], sugar [13], ammonium bicarbonate [14], or magnesium [15]—to reproduce a bone-like porosity structure in which the space-holder particles can be removed by thermal process or dissolution. This step can be performed before or during the sintering process. The particle size and morphology are key factors in the design of the pore structure. One remarkable advantage is that the space-holder is a cost-effective and non-toxic method, which does not release any toxic agent which could affect cellular and bacterial behavior.

On the other hand, several authors have reported that surface properties such as the topography, roughness, chemistry, and free energy of metallic based biomaterials have a strong influence on the initial adhesion and early differentiation of osteoblast cells [16,17]. In 1987, Gristina was the first author to use the term "the race" for the surface, regarding the competition established between bacteria and cells for the biomaterial surface [18]. Also, Gristina remarked that if bacteria attached and colonize faster than osteoblast cells—i.e., bacteria wins the race—an infection will take place, which will be almost impossible to remove. This effect is due to the fast process of bacteria to excrete extracellular polymeric substances (EPS) once they are attached. The size is one the most highlighted difference: osteoblasts are bigger in size, ranging from 20–30 μm, bacteria such as *Escherichia coli* (*E. coli*) or *Staphylococcus aureus* (the most common bacteria presented in titanium dental implants) are around 1–2 μm in length [19]. The stiffness, surface topography (i.e., roughness) and surface chemistry can be tailored in order to promote an advanced cell adhesion process. It has been described the high stiffness behavior of bacteria compared to osteoblast cells, which can be a potential factor in bacteria adhesion and colonization [20]. However, the adhesion mechanisms and the proliferation rate of bacteria are key factors that favor bacteria in the race for the surface. Filia, pili, and the expression of EPS, which constitute the biofilm formation, are some of the mechanisms that participate in the interactions between bacteria and the surface. This EPS covers and protects the bacteria under a biofilm. In this scenario, bacteria becomes more resistant and the antibiotic treatment usually has no effect. On the other hand, osteoblast cells require a more complex mechanism involving integrin receptors and proteins to develop focal adhesion (FA) points. The FA are based of filopodia and lamellipodia cytoplasmatic projections that control and regulate cell to cell interactions and cell surface as well. Therefore, the control of surface properties may reduce or avoid the bacteria attachment resulting in the successful application of the implant on bone regeneration [19].

In a previous work, we reported the manufacture of porous substrates by the space-holder technique which showed a decrease of bacteria attachment [21]. However, the vast majority of studies which focus on Ti implants made use of either bacteria or celld independently and, thus made it difficult to analyze of what type of surface properties may enhance infections and/or osseointegration. Therefore, here we report for the first time the behavior of bacteria and two murine cell lines (myoblast and osteoblast) on c.p. Ti porous substrates with the aim to design a porous implant able to promote bone regeneration and avoid bacterial infections. In this context, this work examines a detailed study of the influence of microstructural characteristics (size, shape, area, and porosity roughness) on the macro and micro-mechanical behavior, as well as the response of the material in the presence of eukaryotic cell cultures (C2C12-GFP and MC3T3E1) and *E. coli* strain bacteria (Gram-negative bacteria).

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

#### *2.1. Fabrication of Ti substrates*

The blends of medical grade 4 c.p. Ti were supplied by SE-JONG (Materials Co. Ltd. Gojan-dong Korea). Fully-dense c.p. Ti samples were obtained by conventional powder technology (PM), while porous substrates were fabricated using space-holder technique (50 vol % NH4HCO3 and different ranges of particle size: 100–200 μm, 250–355 μm, and 355–500 μm). The green samples were produced by pressing at 800 MPa by means of an Instron 5505 universal testing machine (Instron, High Wycombe, UK). Then, the spacer was thermally removed (firstly at 60 ◦C and, then at 110 ◦C under low vacuum

conditions), both stages of the thermal treatment are carried out for 10–12 h and low vacuum conditions of 10−2 mbar). Subsequently, the samples (porous and fully-dense) were sintered in a ceramic tubular furnace during 2 h at 1250 ◦C under high vacuum conditions (~10−5 mbar).

#### *2.2. Microstructural and Mechanical Characterization*

Different techniques were employed to characterize the porosity at different levels (superficial and volumetric), as well as the pore distribution, shape, and roughness previously reported [3–5,14]. Total porosity was determined by both Archimedes' method and image analysis (IA), using a Nikon Epiphot optical microscope (Japan) coupled with a Jenoptik Progres C3 camera and Image-Pro Plus 6.2 analysis software. Confocal laser and scanning electron microscopy (SEM, JEOL JSM-6490LV, JEOL, Tokyo, Japan) were used to evaluate the roughness. Considering the importance of pores roughness on bacterial adhesion, X-ray micro-computed tomography (M-CT) was also performed due to its capability to provide three-dimensional (3D) surface roughness data, among other parameters. These measurements were performed using a custom made X-ray scanner composed mainly of a micron focused X-ray source L8121-01 (with a W-target) from Hamamatsu, Japan [22]. Scans were performed at 100 kV and 100 μA, with a spot size of 5 μm and were recorded with a flat panel detector C7943 (120 × 120 mm, 2240 × 2368 pixel), also from Hamamatsu. A 3D reconstruction of the specimen was obtained by acquiring a certain number of X-ray projections during sample rotation over 360◦, followed by software reconstruction of these projections. This method allowed for the qualitative and quantitative exploration of the interior structure of the porous Ti samples, with a pixel size down to 6.4 μm at a 7.8-fold magnification. Parameters such as equivalent diameter (Deq), total porosity (Pt), and interconnected porosity (Pi) was determined by M-CT. All these analytical techniques have been implemented, following the procedures described in previous works [3–5,14].

On the other hand, the macro-mechanical behavior of Ti substrates was evaluated by ultrasound technique (dynamic Young's modulus, Ed) and uniaxial compression test (Young's modulus, Ec; yield strength, σy). For mechanical compression testing, the specimen dimensions were fixed to standard recommendation (height/diameter = 0.8). The tests were carried out with a universal electromechanical Instron machine 5505 by applying a strain rate of 0.005 mm/mm/min. All tests finished for a strain of 50 pct. The Young's modulus estimation from the compression stress–strain curves was corrected with the testing machine stiffness (87.9 kN/mm). On the other hand, conventional microhardness studies (Vickers indenter Shimadzu, model HMV-G, Japan) at two different applied loads (HV0.3 and HV1) were performed. All the measures were taken three times per substrate type and applied load.

#### *2.3. In Vitro Cellular Experiments*

#### 2.3.1. Analysis of Bacterial Behavior of Porous Substrates

*E. coli* strain Bacterial behavior on porous and fully-dense c.p. Ti substrates is studied, following the protocol summarized in Figure 1. The used *E. coli* was a lyophilized form from Ielab, Spain (*E. coli*; ATCC 25922). The culture solution was prepared with Peptone Water (dehydrated culture media, from Panread AppliChem, Spain) and *E. coli*-Coliforms Chromogenic Medium from Conda Laboratories S.A. The initial concentration of colony-forming units per 100 mL (CFU/100 mL), was determined by membrane filtration, followed by incubation of the membrane on a Petri plate containing *E. coli*-Coliforms Chromogenic Medium at 37 ◦C for 24 h. Specifically, the study includes three initial concentrations: 7.8·104 (C1), 3.3·103 (C2), and 4.3·102 (C 3) CFU/100 mL, respectively.

**Figure 1.** Analytical procedures employed to *E. coli* bacteria analyses.

For *E. coli* bacterial behavior studies, fully-dense and porous substrates were placed in 10 mL test tubes containing each bacterial concentration and these were incubated at 37 ◦C for 1 h. One tube containing only bacteria was used as reference (R).

Afterwards, the bacterial growth was analyzed by three protocols (Figure 1). Protocol 1 (P1) quantified the amount of *E. coli* per tube and per initial concentration by membrane filtration (pore size of the membrane: 0.45 μm) of the solution of each tube [23]. Protocol 2 (P2) evaluated the *E. coli* density on the Ti substrate using a sterile swab to smear the total surface of the Ti samples. P2 was applied on the flat surface, meaning the areas of the samples without pores. Then, the swab was applied following a zigzag pattern on a Petri dish containing *E. coli*-Coliforms Chromogenic Medium and incubated at 37 ◦C for 24 h. Protocol 3 (P3) analyzed the *E. coli* proliferation inside the pores by re-immersing the swabbed substrates in test tubes with 10 mL of sterile Peptone water. This solution was mixed using a vortex mixer for 1 min and cooled at 4 ◦C for 2 h. Tube contents were again mixed with the vortex mixer to detach the bacteria from the cavities and to quantify them using membrane filtration. To guarantee the complete detachment of bacteria from the cavities, the process was repeated three times instead of doing it only once as recommended by ISO 11737:1:2007 standard [24]. The bacterial growth was expressed in total number of CFU for each substrate and concentration. Several measurements were made for each condition (protocol, substrate, and initial concentration).

#### 2.3.2. Evaluation of Cell Adhesion and Proliferation of Eukaryotic Murine C2c12-Gfp Premioblast Cells

Attachment and proliferation of C2C12-GFP were analyzed by inverted fluorescence microscopy (Olympus IX51) and CellD Software (Olympus). To that end, C2C12 murine premioblastic cell line was purchased to America Type Culture Collection (ATCC®CRL-1772™) and transfected, via lentivirus, to self-express constitutively green fluorescent protein (GFP). The presence of this GFP group in cells membrane allowed us to detect the cells and follow the adhesion and proliferation process. The routing passage was performed using DMEM (Dulbecco's modified Eagle medium, Sigma Aldrich), completed with 10% fetal bovine serum (FBS) and 1% of penicillin/streptomycin (100 U/mL, Invitrogen, Germany). For adhesion and proliferation studies, cells were seeded at an initial density of 10,000 cells/cm2,

and after 24 h samples were transferred to new 24-well plates and inverted fluorescent images were taken after day 1, 4, 7, and 10 of cell incubation. The complete medium was changed every 2–3 days. This protocol is summarized in Figure 2.

**Figure 2.** Analytical protocols for the cell studies.

#### 2.3.3. Cell adhesion and Proliferation Studies of Murine MC3T3E1 Osteoblast

A bone cell line, MC3T3-E1 (ATCC CRL-2593), was employed to evaluate the effect of porosity in cell viability, cell proliferation, and cell morphology of osteoblast. The subculture routing was performed in MEM (minimum essential medium, Sigma Aldrich, Germany), completed with 10% fetal bovine serum (FBS) and 1% of penicillin/streptomycin (100U/mL, Invitrogen, Germany). 75 cm2 culture flasks were placed in a humid atmosphere with 5% CO2 at 37 ◦C. Ti substrates were seeded at 10,000 cells/cm2 and after 24 h samples were transferred to new 24-well plates. Figure 2 summaries the protocol for this study.

#### Cell Viability of Murine MC3T3E1 Osteoblast

Cell metabolic analysis was measured in order to evaluate the cell viability at day 4 of cell incubation. To that end, AlamarBlue assay (DAL1100, Thermofisher, USA) was performed following the manufacturer's instructions. For this, samples previously seeded with cells were transferred to new 24-well plates and fresh media containing AlamarBlue dye added (10% of culture media volume). After 1.5 h of incubation in darkness at 37 ◦C, samples were removed and fluorescence signal was read in a microplate reader (Synergy HT, Biotek) using an excitation wavelength of 530 nm and emission wavelength of 590 nm. Fully dense surface was considered control surface as 100% viability. Assays were carried out by triplicate for each sample condition.

#### Cellular Morphology Evaluation of Murine MC3T3E1 Osteoblast

After 10 days of cell incubation, osteoblast cells attached to the substrates were washed twice with PBS, and subsequently, fixed using 4% paraformaldehyde (PFA) solution. Fully-dense and porous Ti samples were rinsed carefully with PBS twice, permeabilized with 0.1% (v/v) Triton X-100 and, finally, washed again with PBS. The actin cytoskeleton was stained with Texas Red®-X phalloidin (Molecular Probes). To that end, the commercial solution was prepared in PBS and left in dark conditions for 20 min at room temperature (RT). Hoechst (Invitrogen, Molecular Probes) was employed as a contrast marker, to detect cell nuclei. Fully-dense and porous substrates were analyzed using an inverted fluorescence microscope (Olympus IX51) with a TRICT filter (λex/λem = 550/600 nm) for Actin, and DAPI filter for Hoechst (λex/λem = 380/455 nm) and images were treated by CellD analysis software (Olympus).

#### Statistical Analysis

In general, each experiment was performed in triplicate with each bacteria or cell line, media, and sample. Mean values and standard deviation are reported. A *p*-value of < 0.01 was deemed to be statistically significant. For in vitro cells, mean and standard deviation are presented and a one-way ANOVA followed by a post Tukey analysis was performed in which values were be statistically significant when *p*-value was < 0.05.

#### **3. Results**

#### *3.1. Microstructural Characterization*

Results concerning to the microstructural characterization of the porous substrates can be observed in Figures 3 and 4. Particularly, Figure 3 (top) shows micro-images of the fabricated samples while Figure 3 (bottom) displays details about the aspect of the pores, size, and distribution as obtained by IA. So, by the micro-image we have an overview of the fabricated samples, while by IA details about the aspect of the pores, size and distribution can be determined. Figure 4 shows data obtained by M-CT which provides three-dimensional information, quite useful to determine the distribution of inner pores as well as the roughness volume percentage of the eroded pore compared to the virtually smoothed pore in dependence of the pore equivalent diameter. The presence of micro-scale roughness pattern within the pores was revealed by M-CT, so total inner surface was calculated analyzing a M-CT representative fraction of total sample volume, taking into account the local deviation of the gray values perpendicular to the contour line within a search distance of 20 μm. Data obtained by the analyses of the images contained in Figures 3 and 4 are summarized in Tables 1 and 2, such as porosity and parameters to characterize the pores (equivalent diameter of pores, Deq, shape factor, and roughness).

**Table 1.** Total, interconnected, and isolated porosity evaluated by different characterization techniques.

**Figure 3.** Macro-photography of the samples and optical microscopy images of the porous substrates for the different ranges of spacer particles size: 100−200 μm, 250−355 μm, and 355−500 μm.

**Figure 4.** Equivalent diameter determined and interconnected and isolated porosity analyzed by M-CT for two different ranges of spacer particles size: 100−200 μm and 355−500 μm.

#### *3.2. Mechanical Behavior*

The macro-mechanical (stiffness, yield strength) and micro-mechanical (Vickers hardness) behavior of porous c.p. Ti is summarized in Table 3. Young's modulus has been measured by ultrasound and uniaxial compression test. These results have been discussed in terms of the advantages and disadvantages of both techniques. On the other hand, the influence of the indentation load and its relationship with the size of the remaining Ti matrix among pores is evaluated. In this context, it should be noted that, for a fixed percentage of total porosity, an increase in pore size implies an increase in the mean distance between pores (the size of the titanium matrix increases).

**Table 3.** Macro and micro-mechanical behavior of porous substrates.

#### *3.3. Bacteria Behavior*

The number of CFU of *E. coli* on titanium substrates as a function of initial bacteria concentration following Protocol 1 is depicted in Figure 5a. It was observed that *E. coli* initial concentration affected the number of CFU on the substrate i.e., the higher the initial concentration the higher the number of CFU. However, as porosity increased (fully-dense vs. porous c.p. Ti) the number of CFU was significantly increased, especially at the lowest initial bacteria concentration (C3). *E. coli* bacteria concentration of the medium could influence the medium-substrate interaction. After incubating the samples for 1 h at 37 ◦C, the number of CFU in the corresponding solutions was determined by membrane filtration. Protocol 2 revealed no bacteria attached onto the flat surface. By Protocol 3, bacteria inside the pores was quantified and its normalization by the surface is shown in Figure 5b,c, respectively. As the size of the spacer increased the number of CFU also increased at both initial bacteria concentrations (Figure 5b).

**Figure 5.** Bacteria quantification following protocols described (P1, P2, and P3). (**a**) CFU of *E. coli* bacteria after membrane filtration (Protocol 1), CFU of *E. coli* bacteria inside the pores measured following Protocol 3: (**b**) absolute values and (**c**) normalized values by the inner pores surface. Inset: magnification view of the concentration view for the normalized values.

#### *3.4. In Vitro Cell Studies*

The different studied substrates were tested in vitro conditions using two murine cell lines. In first place, C2C12-GFP premioblastic line was used to evaluate the process of cell adhesion and proliferation after 1, 4, 7, and 10 days of cell culture compiling the results in Figure 6. C2C12-GFP cells were capable to attach and proliferate over the surface and inside the pores, increasing the cell density presented on the samples over time. In the first 24 h, the inverted fluorescence images showed similar cell density between fully-dense and porous substrates. After 4 days, several small green dots which correspond to small alive cell clusters were observed totally dispersed on the surface. Furthermore, at longer incubation periods, cells were able to proliferate, showing an increased cell density over the entire surface. Indeed, at day 10 C2C12 GFP cells were covering the entire porous structure, even filling the pores attaching on pore's walls and arriving to the edge of the samples (see Figure 7 after 10 days, 255–350 μm sample). Figure 7 showed a higher magnification of C2C12-GFP cells growing at 4 days to analyze the cell morphology on c.p. Ti samples. A clear different cell shape was observed between cells growing on flat surfaces and inside pores. Premioblastic cells growing on fully-dense samples showed fusiform morphology, with an elongated cell shape. This cell morphology, most frequent of fibroblastic phenotype, was also observed in flat surfaces of porous substrates independently of pore size (see Figure 7).

**Figure 6.** Inverted fluorescence microphotographs of cell adhesion and proliferation of C2C12-GFP cells after 1, 4, 7, and 10 days of cell culture. Common scale bar, 1mm, except for 7 days and pore size 355–500 μm, which a specific scale bar is indicated (200 μm).

**Figure 7.** Higher magnification of microphotographs of C2C12-GFP cells growing after day 4 of cell incubation. Common scale bar.

Cell viability determination was performed after 4 days using osteoblast cells and the results were compiled in Figure 8. Cell viability was expressed as percentage of fully dense samples, showing an increased cell viability of osteoblast growing on the three porous structures compared to the control surface (fully-dense). Although cell viability results are slightly higher for porous samples respect to fully-dense substrates, this observed trend was not statistically significant and, therefore, no differences in cell viability could be observed between surfaces (*p* > 0.05).

**Figure 8.** Cell viability quantification of preosteoblast MC3T3 growing on fully-dense and porous substrates after 4 days of cell incubation. Percentages of cell viability are coupled with representative images from a macroscopic point of view. No statistics differences were observed between surfaces (*p* > 0.05), however, porous substrates reached higher percentages than the control fully-dense surface. Common scale bar for all the images.

Figure 9 presents the images of preosteoblast MC3T3E1 cell cytoskeleton morphology, showing actin fibers (in red) and cell nuclei of stained osteoblast (in blue) after 10 days of cell culture. Firstly, it was observed that, the entire surface of fully-dense and porous substrates was completely covered by osteoblast. Some flat areas in samples prepared with higher space-holder particles size (355–500 μm) were not totally covered and a heterogeneous random distribution was observed compared to the homogeneous cell spread of osteoblast in the other three surfaces. Images on the middle and bottom depict porous structures in Figure 9 corresponds to a higher magnification area (using 10× and 20× lenses). These micro-photographs showed the same area with focusing in two different points, on flat surface and inside pores. As it is shown, the cell nuclei of MC3T3E1 cells in the three porous substrates were covering the total flat surface exposed and the elongated actin fibers connecting between cells almost producing a cell sheet. In contrast, fully-dense surfaces revealed more than one cell monolayer growing on top of another cell sheet. This situation might reduce the corresponding cell viability of the cell culture and promote the detachment of the cell sheet.

**Figure 9.** Immunofluorescence images of cell actin cytoskeleton (in red) and cell nuclei (in blue) of preosteoblast MC3T3 growing for 10 days. Medium and lower images correspond to higher magnification of areas focusing in same area on flat and inside pores of the studied specimens. Bottom images show the cells attached inside pores and medium the same cells well focused on flat areas.

#### **4. Discussion**

As it is shown in the optical microscopy images (Figure 3), all porous c.p. Ti substrates present two different ranges of pore sizes; pores with a size below 50 μm were generated during the sintering process. In addition, they are not critical in mechanical and biological behavior of the substrates. Otherwise, the larger pores correspond to the former spacer particles, showing a *Deq* in the range of the size of these particles (161.1 ± 28.5 μm for 100–200, 261.5 ± 9.0 μm for 250–355, and 293.4 ± 28.2 μm for the 355–500 μm of the spacer particles size, as it was determined by IA). Porosity results obtained by the different techniques (Table 1) differ due to the inherent characteristics of them: Archimedes' is a volumetric method, IA analysis is superficial while M-CT only takes into account a volumetric fraction of the sample. The total porosity is slightly below the expected 50%. Nevertheless, different factors can influence. For instance, remaining small amount of undissolved NH4HCO3, which are enclosed in the isolated pores, and the slight reduction of the volume associated with the space-holder during sintering would contribute to reduce the porosity while an increase on the porosity could be also expected because of the additional micro-pores generated during the sintering.

Total inner surface has been calculated analyzing a M-CT representative fraction of total sample volume. The lowest value was obtained for the intermediate size 255–350 μm (260.1 ± 29.6 mm2), while similar values were calculated for the other samples (321.0 ± 44.8 and 311.9 ± 29.9 for 100–200 μm and 350–500 μm, respectively). Similarly, roughness quantification by the analysis of confocal microscopy images revealed an increased roughness in porous surfaces compared to flat surface, as it has been described above.

Concerning the pores, *Deq*, analyzed by IA and M-CT (see Table 2), a similar tendency was observed for both techniques. In the case of IA, also information about the shape factor has been obtained, revealing high symmetry of the pores, since it is close to one for all the substrates. By M-CT, the roughness has been analyzed (Table 2), the higher pore size (355–500 μm), the smoother surface. Roughness improves cellular adhesion [3,5], so it is desirable for osseointegration; however, depending on the roughness values it can affect bacteria adhesion.

Additionally, the mechanical behavior of the substrate is an important feature since the ideal implant should be mechanically similar to the bone. The mechanical properties of the fully-dense substrates are far from those of cortical bone tissues. However, all porous substrates present a Young's modulus similar to the one of the cortical bone [25]. The yield strength values of the porous samples are lower than that of the bone, but they remain closer to the ideal values compared with fully-dense substrates [26].

Microhardness has been evaluated by two load levels that reveal different results: HV0.3 and HV1 (Table 3). Comparing the results, the differences in terms of applied load for one substrate could be related to the localized microplasticity phenomena and indentation size effects. The microhardness and the local stiffness of the porous c.p. Ti substrates depend on the remaining material and the surrounding pores (size and distribution). For low load (HV0.3), the microhardness of porous substrates and fully-dense samples is similar due to the Ti matrix remaining sufficiently large between the pores (i.e., no porosity influence). However, for higher load (HV1), an effect of the pore and porosity on the micro-mechanical behavior (material collapse, buckling, etc.) can be appreciated. The pore size influences the microhardness in an opposite way, it means, the samples with a higher pore size present a lower microhardness value due to the larger area of the Ti-matrix among the pores. However, a detailed study may be required to determine the main influential parameters to the results as the mean free path between pores and the number of indentations.

Once the mechanical properties have evaluated, next the bio-functional behavior is studied. Following Protocol 1, the bacteria concentration of the medium was measured. As it is shown in Figure 5a, both fully-dense and porous substrates inhibited the bacterial growth for the three studied concentrations compared to the reference solution (bacteria only). This reduction was particularly significant for fully-dense c.p. Ti substrates and lower initial bacteria concentrations (C3) while for bacteria solution (C1), the substrate influence was irrelevant. This difference between the fully-dense and the porous samples can be explained by two possible factors: the growth rate is higher for lower bacteria concentrations in the media, achieving a stationary state and/or higher concentrations of bacteria mean that the bacteria thronged the cavities, hampering penetration into the pore.

Measurements obtained by Protocol 2 for the *E. coli* strain showed no bacteria growth and adhesion onto the flat surfaces. For the porous samples, the flat surface corresponds to the polished c.p. Ti matrix surface that remains among all the pores. After incubating the Petri dish where the swab was smeared, no bacteria grew. For this same strain, Protocol 3 revealed bacterial presence inside the pores (Figure 5). This difference in terms of behavior between the two surfaces could be related to the micro-roughness pattern of the porous surface. A similar pattern has been reported for osteoblasts [4,5]. As it is shown in Figure 5b, the number of CFU inside the pores was larger for porous substrates with larger pore size and at higher concentrations. A higher initial bacteria concentration implies higher probability of bacteria attaching inside the pores. The influence of the pore size could be ignored compared to the effect of the concentration for absolute values. Nevertheless, once the number of CFU was adjusted in relation to the area (Figure 5c), it was relatively insignificant.

To compare the bacteria behavior and the studied substrates different factors have been considered. On the one hand, the presence of pores and, on the other hand, the real surface, which is related to pores inner roughness. In this context, it has been observed that the presence of pores promote bacteria proliferation. Otherwise, normalized values by real surface area reveals that the value was not higher than expected. Pore number, size, shape, and surface roughness inside the pores are also relevant. At higher pore size (368.4 ± 9.1 μm) as measured by M-CT, more bacteria are adhered into them. In contrast, bigger pores present smoother walls (3.3 ± 0.6 μm), which would make more difficult the bacteria adhesion. However, the obtained results for this bacteria strain show that pore size is more significant than roughness.

Regarding the behavior of mammalian cells such as myoblast or osteoblast cell lines, several studies have been performed. Cell viability, adhesion, proliferation, and differentiation of osteoblast are key factors for bone tissue reconstruction purposes [27] and these processes can be influenced by tailoring the surface properties of the medical device such as hydrophilicity, roughness, pore size, morphology, and distribution. Pore size is one key parameter that has been deeply reported in several studies suggesting pore size between 20 μm to 500 μm are appropriate for bone regeneration [28]. These values covered a wide range of pore sizes, however, if the pores are quite small, the blood vessel network cannot exchange nutrients and oxygen to new bone tissue. Conversely, if the pore is too large, the mechanical properties may compromise the internal stability of the implant. Although the most suitable pore size for bone regeneration process is still controversial, for Ti implants a pore range from 150 to 500 μm has been established, taking into account the role of the interconnected pores in the proper development of the new vascularization system [29]. During the course of this study, all porous surfaces showed a non-toxic behavior and a good cell adhesion for C2C12GFP achieving similar cell density after 24 h compared to the fully-dense control surface (Figure 6). In fact, the presence of cells attached on the surface on all c.p. Ti substrates serves as an indicative factor of biocompatibility of the manufacturing process of porous structures by powder metallurgy and space-holder techniques which have been previously described as safety technologies to develop non-toxic 3D porous structures [30].

Even though the cell density attached on the surface after day 1 is low, C2C12GFP cells could proliferate and grew further after 4, 7, and 10 days, increasing the cell density, which was observed as a higher fluorescence intensity signal (Figure 6). After 4 days, C2C12-GFP density increased compared to day 1. However, premioblastic cells grew in clusters randomly dispersed, leaving flat areas empty without cells attached on the surface. This behavior evolved, and after day 7, a continuous cell monolayer was developed covering the center and, in some cases, reaching the edges of the c.p. Ti samples, and thus, reducing the flat areas with no cells attached. Some pores were covered by osteoblast after 10 days which appeared as intense green fluorescence signal. Here, many areas were totally covered, including flat and inside pores, in which cells were expanded and connected forming a cell monolayer. At this time, the presence of pores offers a different surface (increased roughness values and surface contact area) which may allow the proper growth of both myofibroblast/osteoblast and *E. coli* as we have previously described. In case of fully-dense surfaces, the presence of more than one cell monolayer could reduce the cell viability, but bacteria may take advantage and spread and migrate to new sites over the medical implant.

Several authors have reported that the presence of pores (and their size, morphology, roughness, and interconnectivity), the surface chemistry, and surface free energy of metallic based biomaterials have a strong influence on the initial adhesion and early differentiation of osteoblast lineage cells [31–33]. Although the aim of this study was to evaluate the bacteria behavior on porous c.p. Ti structure compared to cell adhesion and proliferation of in vitro cell lines, both C2C12GFP and MC3T3 cells shared some similar points. Firstly, both cell lines were able to attach and proliferate confirming the cell viability values of osteoblast which were higher on porous surfaces over the samples after 4 days (see Figure 8). Even where there were not statistically differences on cell viability, the three porous c.p. Ti substrates achieved higher percentages than control fully-dense surface (113.89%, 112.28%, 117.74%, and 100%, respectively). The larger pore size, 355–500 μm, offered higher cell viability percentage

which was also correlated to the higher attached *E. coli*, confirming the higher pore surfaces structure as an advance surface to promote cellular and bacteria adhesion.

Secondly, the specific tailored microstructure of porous scaffolds by space-holder have a strong impact on cell adhesion which one can observe in the differential cell morphology showed in Figures 7 and 9. The cell morphology of osteoblasts is usually a response to the surface properties of the inner material. Inside the pores, osteoblast cells showed a more cuboidal cell cytoskeleton compared to a more fusiform cell shape on flat surfaces (either fully-dense or flat surface of porous c.p. Ti samples). This distinction in cell morphology can be correlated to an advanced cell adhesion state which also can affect other cellular processes such as osteoblast cell differentiation growing inside pores as it was described in the literature [33].

Some reports have described the influence of surface properties on cell morphology and how the shape of the cells is also connected with other cellular functions such as cell adhesion, proliferation, differentiation, and mineralization processes [34,35]. Furthermore, the increased inner roughness values of pores compared to flat surface (of fully-dense or even flat surface of porous samples) has been described as a key factor for osteoblast adhesion and differentiation studied [5,33]. Inside the pores, osteoblast cells attached on the wall perceive higher load transfer and this local stress will promote cell differentiation and mineralization, and therefore, the need for bone ingrowth [33]. However, these cells growing inside pores presented a different cell morphology, revealing a more cuboidal structure, round shape, and less elongation.

The cell morphology observed inside the pores could be related to an advance cell adhesion, being the 100–200 μm porous substrate the best pore size to promote cellular adhesion and reduced bacteria attachment compared to larger pore size 355–500 μm and fully-dense substrates.

#### **5. Conclusions**

The microstructural parameters of titanium porous samples fabricated by space-holder technique are consistent with the initial design criteria (characteristics of the spacer holder particles: volume fraction and size range), supporting the viability and reliability of this economic route of manufacturing substrates with controlled porosity. The porous substrates obtained with this technique, in the range of 100–200 μm, presented the best mechanical balance (higher Young's modulus, yield strength, and microhardness). Besides, this substrate also revealed the highest interconnected porosity, potentially improving the bone in-growth. It is worth to highlight that both, spacer particle size and pores inner surface topography, have a strong influence on adhesion of osteoblasts and bacteria proliferation (*E. coli* strain). However, the results revealed that the predominant parameter in bacteria behavior is pore morphology (equivalent diameter). Therefore, the bigger pore size showed higher bacteria attachment. To summarize, substrates with smaller pore sizes can be considered the most suitable, since this range of pore sizes can assure the best biomechanical, cellular, and antibacterial behavior.

**Author Contributions:** Conceptualization, project administration, supervision J.A.R.-O. and Y.T; Investigation, formal analysis, validation A.C., A.M.B., C.D.-T., M.D.G, J.L., F.G.M., J.V.C.-R., and J.J.G.; Discussion and writing—original draft preparation, all authors.

**Funding:** This research was funded by the Ministry of Economy and Competitiveness of Spain under the grant MAT2015-71284-P and of the Junta de Andalucía—FEDER (Spain) through the project ref. P12-TEP-1401.

**Acknowledgments:** The authors dedicate this paper to the memory of Juan J. Pavón Palacio (University of Antioquia, Colombia). Authors thank to A. S., C. L., and M.V.S. for experimental work for bacteria studies. The authors would like to thank J. Pinto for assistance in micro-mechanical testing. AMB also thanks the financial support of the University of Seville for the research stay at Institute of Applied Materials, Helmholtz-Zentrum Berlin für Materialien und Energie, Berlin (Germany) (Grant: VI Plan Propio de Investigación y Transferencia—US 2018, I.3A2).

**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* **Microstructure of a V-Containing Cobalt Based Alloy Prepared by Mechanical Alloying and Hot Pressed Sintering**

**Niannian Li 1, Fengshi Yin 2,\* and Liu Feng 3,\***

Received: 8 March 2019; Accepted: 16 April 2019; Published: 22 April 2019

**Abstract:** In this paper, a bulk V-containing cobalt-based alloy with high chromium and tungsten contents was prepared by mechanical alloying and hot pressed sintering using Co, Cr, W, Ni, V and C pure element powders. XRD, SEM, TEM and Vickers hardness tests were employed to characterize the microstructure and mechanical properties of the mechanical alloyed powders and hot pressed bulk cobalt-based alloy. The results show that all elements can be mixed uniformly and that the Co, Cr, and Ni elements were made into an amorphous state after 10 h ball milling in a high energy ball miller. The microstructure of the prepared bulk alloy was composed of a γ-Co matrix with a large number of nano-twins and fine M23C6 and M12C carbide particles well-distributed in the alloy. The V element was mainly distributed in M23C6-type carbide and no V-rich MC-type carbide was found in the microstructure. The prepared alloy had a high hardness of 960 ± 9.2 HV and good a fracture toughness *KIc* of about 10.5 ± 0.46 MPa·m1/2. The microstructure formation and strengthening mechanisms of the prepared cobalt-based alloy are discussed.

**Keywords:** cobalt-based alloy; mechanical alloying; hot pressed sintering; microstructure

#### **1. Introduction**

Stellite alloys, a kind of Co-based alloys, are very important in the development of science, technology, and the progress of industry. They are mainly composed of carbon, cobalt, chromium, tungsten, molybdenum, etc. Metal carbides formed in the alloy are used as hard phases and Co as the matrix bonding phase [1,2]. Stellite alloys are widely used industries such as aerospace, nuclear, mining machinery, and material machining due to their characteristics of wear-resistance, high temperature resistance, high strength, high hardness, and a certain toughness. In order to further increase the wear-resistant properties, Tribaloy alloy, another kind of Co-based alloys, which contains a large volume fraction of intermetallic Laves phases in γ-Co matrix, has been developed [3,4]. It is the presence of this large volume fraction of Laves phases that enables these materials to resist wear under poor or unlubricated conditions.

Typically, cobalt-based alloys are prepared by casting or powder metallurgy technology. The casting process is simple in terms of the technology needed and low in cost, but the size and distribution of carbides are strongly dependent on the casting temperature of the melt, cooling rate, and thermal treatment condition, which must be strictly controlled to avoid defects such as macro-segregation, micro-segregation, porosity, solidification shrinkage, and second phase inclusions [5,6]. Cobalt-based alloys can be also manufactured by powder metallurgy process using pre-alloyed powders [7,8]. However, chemical element segregation and oxidation will occur on the surface of the pre-alloyed powders, which introduces the prior particle boundary (PPB), and therefore, decreases the mechanical properties of the final powder metallurgy product [9]. Mechanical alloying (MA) [10–12] and hot pressed sintering, as another powder metallurgy technology, have the potential to shorten the whole preparing period and produce bulk materials with a uniform fine-grained microstructure. The as-prepared materials exhibit excellent properties such as high density, good hot workability and improved mechanical properties.

It is well known that vanadium is a strong carbide-forming element and vanadium-containing high speed steel W18Cr4V has a good wear-resistant property. However, there is no report about the effect of vanadium on cobalt-based alloys. Berthod [13] reported a TiC carbide containing cobalt-based alloy. In this work, based on the Stellite alloys, we designed a novel vanadium-containing cobalt-based alloy with high chromium and tungsten contents. The alloy was prepared using MA and vacuum hot pressed sintering methods aiming to further increase the wear-resistant property of cobalt-based alloy. The microstructure and mechanical properties of the mechanically alloyed powders and hot pressed bulk cobalt-based alloy was characterized.

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

The cobalt-based alloy with a nominal composition of Co-32.4%, Cr-32%, W-25%, Ni-5%, V-5%, and C-0.6% (in wt. %) was prepared by mechanical alloying and hot pressing using pure element powders of Co, Cr, W, Ni, V, and C. The purity of these elemental powders was greater than 99.5%. The process of mechanical alloying was carried out in a high energy ball mill (SPEX8000D, SPEX SamplePrep, Metuchen, NJ, USA) using a stainless-steel vial and stainless-steel balls of 7 mm in diameter. The ratio of ball to powder was 10:1 in weight. The 5 wt. % ethanol was used as the process control agent during the ball milling process. The volume of the vial was 56 mL. A 10 g mixture of the powders with the stainless-steel balls was charged into the stainless-steel vial under high pure argon atmosphere. The ball milled powders were put into a graphite mold and sintered in a vacuum hot press sintering furnace (ZRY-30L, Jinzhou Hangxing Vacuum Equipment Co., Ltd., Jinzhou, China). After the furnace was evacuated using a vacuum pump to 10 Pa and then heated to 1200 ◦C at a heating rate of 4 ◦C/min, the ball milled powders were sintered at a pressure of 60 MPa for 20 min and cooled in the furnace. The ball milled powders and hot pressed bulk sample were examined by X-ray diffraction with Cu Kα radiation (XRD, Bruker AXS D8 Advance, Bruker AXS GmbH, Karlsruhe, Germany), a scanning electron microscope (SEM, FEI Quanta 250, FEI, Hillsboro, OR, USA), and a transmission electron microscope (TEM, FEI Tecnai G2 F20, FEI, Hillsboro, OR, USA) equipped with energy dispersive X-ray spectrometry (EDS, EDAX Apollo) and a scanning transmission electron microscope mode (STEM). Quantitative metallography was used to measure the volume fraction and particle size of various phases in the microstructures of the obtained alloy. The TEM samples were prepared by an ion thinning instrument (Gantan 695.C, Gatan, Inc., Pleasanton, CA, USA). The hardness of the hot pressed bulk sample was measured by a Vickers hardness tester (Laizhou Huayin Testing Instrument Co., Ltd., Laizhou, China) using a 5 kgf load. Fracture toughness was determined according to the hardness indentation method. Five measurements were used to calculate average value of hardness and toughness.

#### **3. Results**

Figure 1 shows the XRD patterns of the powders ball milled with various time. It can be seen that at the initial state, the diffraction peaks of Co, Cr, W, and Ni are clearly shown in the pattern, while the corresponding diffraction peaks of C and V were not present due to their low content. The diffraction intensity of all elements decreased gradually with the milling time. The diffraction peaks of Co, Cr, and Ni elements almost disappear except for the W element after 10 h milling. No obvious change was seen upon further milling for up to 30 h, but peak broadening of W elements was observed. The disappearance of the diffraction peaks of Co, Cr, and Ni elements indicates that these element

powders were transformed into an amorphous state. The XRD peak broadening reveals that the W powders had been transformed into a nanocrystalline state. The crystallite size of the W powders after milling 10 and 30 h, estimated using the Scherrer formula [14], was found to be approximately 25 and 17 nm, respectively.

**Figure 1.** X-ray diffraction patterns of the powders ball milled for different time.

The morphology and micro-structural details of the milled powders were investigated by SEM. Figure 2 shows the SEM images for the powders ball milled for 10 h. The milled powders had an irregular morphology with a size below 15 μm. A local amplified back-scattered electron (BSE) image (Figure 2b) of one particle shows that some white particles about 30 nm in size are distributed uniformly in the powders. According to the XRD results of the milled powders (Figure 1), the white particles were deduced to be W phase. The element mapping image (Figure 2c) of EDS reveals that all elements had been mixed uniformly after 10 h ball milling.

**Figure 2.** SEM images of the cobalt-based alloy powders after milling for 10 h, (**a**) morphology of the milled powders, (**b**) back-scattered electron (BSE) image of one particle, and (**c**) element mapping image of energy dispersive X-ray spectrometry (EDS) results.

*Metals* **2019**, *9*, 464

The cobalt-based alloy powders milled for 10 h were selected for vacuum hot pressed sintering in the subsequent experiment. Figure 3 is the SEM-BSE image of the obtained bulk cobalt-based alloy showing that there were three kinds of phases in the alloy according to the contrast, i.e., bright, grey, and dark phases. XRD results (Figure 4) of the obtained bulk cobalt-based alloy show that these three phases were M12C (Co6W6C, JCPDS 22-0597), γ-Co (JCPDS 15-0806) matrix, and M23C6 (Cr23C6, JCPDS 85-01281), respectively, all of which had a cubic structure. According to the XRD results, the lattice constants of M12C, γ-Co and M23C6 were 1.0922, 0.3552, and 1.0657 nm, respectively. No PPBs were found (Figure 3) in the prepared alloy by hot pressed sintering using mechanical alloyed powders. Figure 5 is STEM high-angle annular dark field (HAADF) image and selected area electron diffraction (SAED) patterns of the phases with different morphology and contrast.  Table 1 summarizes the STEM-EDS results and identified phases by SAED patterns for the phases labeled in Figure 5a. Fe in EDS results would have been introduced from stainless-steel vial and balls during mechanical alloying process. According to the results mentioned above, we can conclude that the bright phase was W-containing M12C carbide in which lots of Cr and V atoms dissolved, the grey phase is the γ-Co matrix, and the dark phase was Cr and V-rich M23C6 carbide. The V element was mainly distributed in M23C6-type carbide and no V-rich MC-type (M indicates metal elements) carbide was found in the obtained alloy. The M12C-type carbide particles had a size of about 1 μm dispersed in the γ-Co matrix. Some coarse M23C6-type carbide particles have a strip-like shape with about 1 μm in length dispersed in the γ-Co matrix and some fine M23C6-type carbide particles with about 200 nm in length present within both the γ-Co matrix and M12C-type carbide. The volume fraction of M12C and M23C6 phases were determined to be about 19.4% and 11.95%, respectively, by using quantitative metallography.

**Figure 3.** SEM back-scattered electron image of the obtained bulk cobalt-based alloy.

**Figure 4.** XRD result of the obtained bulk cobalt-based alloy.

**Figure 5.** STEM high-angle annular dark field (HAADF) image of the obtained bulk cobalt-based alloy (**a**) and selected area electron diffraction (SAED) patterns of the different phases labeled in (**a**), (**b**) point 1, bright, (**c**) point 2, dark, (**d**) point 3, dark, and (**e**) point 4, grey.

**Table 1.** STEM-EDS results of the phases in the prepared bulk cobalt-based alloy (at. %).

Interestingly, the TEM bright field image of the prepared bulk cobalt-based alloy (Figure 6) showed a large number of nano-twins in the γ-Co matrix. The thickness of twin lamellae varied between 10–60 nm.

The alloy had a high hardness of the 960 ± 9.2 HV5. Figure 7 shows the indentation morphology of the Vickers hardness test and the crack produced by the hardness indentation showing the fracture of M12C and M23C6-type carbide particles as well as the plastic deformation of γ-Co matrix, marked by arrows A, B, and C in the Figure 7b, respectively. Only a very short crack was found revealing that the prepared bulk alloy has a good toughness. According to the indentation methods for determining toughness, the fracture toughness *KIc* of the alloy was about 10.5 ± 0.46 MPa·m1/2 calculated by the Equation (1) [15]:

$$K\_{l\varepsilon} = \delta(E/H)^{1/2}(P/a^{3/2}),\tag{1}$$

where, δ is a geometrical factor, for a Vickers indentation δ = 0.016; *E* is the Young's modulus, according to Ahmed's report [7], *E* = 320 GPa was selected in Equation (1); *H* is the hardness of the sample, *H* = 960 HV = 9.4 GPa; 2*a* is the crack size along the diagonal of the indentation, 2*a* = 115.2 ± 3.4 μm; *P* is the load for Vickers indentation, *P* = 49 N.

**Figure 6.** TEM bright field image of the sintered cobalt-based alloy showing a large number of nano-twins in the γ-Co matrix. Inlet is the SAED pattern for the marked area.

**Figure 7.** The indentation morphology of the Vickers indentation (**a**) and the crack generated by the indentation (**b**).

#### **4. Discussion**

The common cobalt-based alloy carbides are M7C3, M23C6, M6C, MC, etc. Service or aging at elevated temperature causes a large amount of secondary carbides, such as M6C and M23C6 [16]. In the present study, the prepared alloy had M6C and M23C6-type carbides (Figure 4). The vanadium was mainly distributed in M23C6-type carbide in the prepared cobalt-based alloy, as shown in Table 1, instead of forming V-rich MC-type carbide. Even though vanadium is a strong MC-type carbide-forming or MX-type (M indicates metal elements, X indicates carbon or nitrogen) carbonitride-forming element, V-rich MC-type carbide is not found in the prepared bulk cobalt-based alloy with 5 wt. % V. However, MX-type precipitate is usually found in 9-12Cr-type martensitic/ferritic heat resistant steel with much less V content of 0.2 wt. % [17,18]. The microstructural characteristics of the prepared alloy is that the M12C-type carbide particles and coarse M23C6-type carbide particles are distributed in the γ-Co matrix in which there are a large number of nano-twins, and fine M23C6-type carbide particles present within both the γ-Co matrix and M12C-type carbide particles (Figure 3).

According to the results mentioned above, the microstructure formation mechanism of the prepared bulk cobalt-based alloy can be illustrated as shown in Figure 8. During the heating and hot pressed sintering process, the mechanically alloyed powders transform from an amorphous state with some nano-sized W particles to γ-Co supersaturated solid solution, then the M12C-type and coarse M23C6-type carbides are formed as follows: 23(Cr, V) + 6C → (Cr, V)23C6 and 12(W, Cr, Co, V) + C → (W, Cr, Co, V)12C [19]. Finally, the fine M23C6-type carbide particles within both the γ-Co matrix and M12C-type carbide particles are precipitated from the γ-Co matrix and M12C-type carbide during the furnace cooling process after hot pressed sintering because their size is much smaller than that of the striplike M23C6 carbide. M23C6 carbide usually precipitates along grain boundaries during the cooling process in conventional alloys, such as austenitic steels and Ni-based superalloys. While the fine M23C6 carbide precipitates within γ-Co matrix in the present V-containing alloy, we think the reason might be that the M23C6 carbide contained a lot of V which has a larger atom size and lower diffusion coefficient than Cr resulting in nucleation of M23C6 within γ-Co matrix and growing up slowly.

The nominal composition of the alloy in atomic fraction was 35.8% Co, 40.1% Cr, 8.9% W, 5.5% Ni, 6.4% V, 3.2% C. From the STEM-EDS data as shown in Table 1, it can be seen that almost all W was distributed in M12C, almost all V in M23C6 and the atomic fraction of Cr in M23C6 is about two times the atomic fraction of V. Based on the STEM-EDS data, the formation of M12C and M23C6 may consume about 6.1%C in atomic fraction, much more than the added 3.2% C. In fact, the carbon content in the prepared alloy was determined to be 1.9% in weight fraction, i.e., 9.8% in atomic fraction by a carbon and sulfur analyzer. The stainless-steel vial and stainless-steel balls were used to prepare the Co-based alloy powders in this work. According to the STEM-EDS results, only a small amount of Fe was indeed introduced into the sintered alloy, it is impossible that the carbon introduced from the stainless-steel vial and stainless-steel balls have much influence. The excess carbon may come from graphite mold during used for hot pressed sintering and causes an increase in amount of carbides.

The mechanism of nano-twins formation in the prepared cobalt-based alloy is not fully understood yet. Twins are common in metals with a FCC structure because of low stacking fault energy [20,21]. Twins are usually formed by sliding a single crystallographic dislocation on a continuous {111} plane driven by a Peach-Koehler force [22]. However, there is no report about the existence of large amounts of nano-twins in a cobalt-based alloy. From a thermodynamic point of view, the formation of twins decreases the total interfacial energy, because the excess energy for coherent twin boundaries (TB) is much smaller than that for conventional high angle grain boundaries (GB). Twins prefer to nucleate at GBs or triple junctions (TJ) to reduce the GB energies by means of the twinning-induced orientation change. Although an extra TB is formed, the sum of the interfacial energies (including GBs and TBs) will be reduced by twinning [23]. In the prepared alloy, the lattice constants of M12C and M23C6 carbides were 1.0922 nm and 1.0657 nm, respectively, about three times the γ-Co lattice constant (0.3552 nm) calculated using XRD data. Therefore, the interface between γ-Co and carbides are prone to form coherent interfaces, which causes a large lattice stress and increases the interfacial energy. In order to reduce the interfacial energies, a large number of nano-twins are formed in the γ-Co matrix.

**Figure 8.** Schematic illustration of the microstructure evolution of cobalt-based alloy during hot pressing sintering and cooling.

The hardness of the prepared bulk cobalt-based alloy was up to 960 ± 9.2 HV5. Casas et al. [11] reported that the prepared ternary and quinary cobalt-based superalloys by field assisted sintering techniques have a hardness of about 650 and 900 HV, respectively. Ahmed [7] used Re-Hot Isostatic Pressing to prepare Stellite 20 alloy with a hardness of 650 HV. Compared with these reported results, the hardness of the prepared cobalt-based alloy in this work was increased significantly. It is reported that a Tribaloy T800 alloy coating has a hardness up to 1000 HV, however, the brittle nature of Laves phases causes a low fracture toughness and reduces the performance of T800 alloy in a wide range of applications where other properties such as certain ductility and fracture toughness are also needed [24]. The high hardness depends on the microstructure. A large amount of fine carbide dispersed uniformly in the prepared cobalt-based alloy, which served as a reinforcement agent. Carbides distributed at the γ-Co grain boundaries prevent slipping and migration of grain boundaries and inhibit grain growth. Carbides inside the γ-Co grains block dislocation movements to enhance the matrix. The W alloying element added in the cobalt based alloy was as high as 25%, which means that more W-containing M12C carbide in the prepared alloy can effectively improve the hardness, more than in a common cobalt based alloy [25]. In addition, many nano-twins are also found in the γ-Co matrix, which also play an important role in strengthening the matrix by blocking the dislocation motion [26,27], further enhancing the hardness. This is also the main difference in the reinforcement mechanism compared with the previously studied cobalt-based alloys [28]. Therefore, under the synergistic effect of solid solution strengthening, dispersion strengthening and nano-twins, the as-prepared cobalt-based alloy presents high hardness. Moreover, the matrix of the prepared Cobalt-based alloy is γ-Co with a FCC structure and has a fine grain structure. The crack propagation can be restrained by the plasticity of γ-Co matrix as shown in Figure 7 marked with arrow C. Therefore, the prepared cobalt-based alloy had not only high hardness but also good toughness, and a good wear-resistant property of the alloy can be expected.

#### **5. Conclusions**

In summary, a bulk V-containing cobalt-based alloy with high chromium and tungsten contents was prepared by mechanical alloying and hot pressed sintering using Co, Cr, W, Ni, V, and C pure element powders, the following can be concluded:

(1) All elements can be mixed uniformly while Co, Cr, and Ni elements are made into an amorphous state after 10h ball milling in a high energy ball miller.

(2) The microstructure of the prepared alloy was composed of γ-Co matrix with a large number of nano-twins and fine M23C6 and M12C carbide particles which were well-distributed in the alloy. The V element was mainly distributed in M23C6-type carbide and no V-rich MC-type carbide is found in the microstructure.

(3) The prepared alloy had a high hardness up to 960 ± 9.2 HV5 and good fracture toughness *KIc* of about 10.5 ± 0.46 MPa·m1/2.

(4) The high hardness was mainly attributed to the synergistic effect of solid solution strengthening, dispersion strengthening, and the large number of nano-twins.

**Author Contributions:** Conceptualization, N.L.; methodology, N.L. and F.Y.; validation, N.L.; formal analysis, F.Y. and L.F.; writing—original draft preparation, N.L.; writing—review and editing, L.F.

**Funding:** This research was funded by the natural science foundation of Shandong province (Grant Number: ZR2016EMM01).

**Acknowledgments:** This work was supported by the natural science foundation of Shandong province (Grant Number: ZR2016EMM01). The authors are grateful for the assistance from Jinye Niu.

**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* **Production of Ultrafine Grained Hardmetals by Electrical Resistance Sintering**

#### **Jesús Cintas 1 , Raquel Astacio 1, Francisco G. Cuevas 2,\* , Juan Manuel Montes 1, Thomas Weissgaerber 3, Miguel Ángel Lagos 4 , Yadir Torres 1 and José María Gallardo 1**

Received: 22 December 2018; Accepted: 28 January 2019; Published: 1 February 2019

**Abstract:** In this work, powders of cemented ultrafine WC-6 wt.% Co were consolidated. The feasibility of the medium frequency electrical resistance sintering (MF-ERS) technique were studied to prevent WC grain growth during consolidation. Porosity and hardness were measured at different zones of the MF-ERS compacts. The compacts showed a slight inhomogeneity in their properties across their section, but it was controlled by choosing suitable values of the processing parameters. The optimal values for the material studied were current intensities between 7 and 8 kA and sintering times between 600 and 800 ms. The main achievement using this consolidation method was that sintered compacts essentially maintained the initial WC grain size. This was attained to processing times of less than 2 s, and without the need for using protective atmospheres.

**Keywords:** electrical resistance sintering; hardmetal; ultrafine grain

#### **1. Introduction**

Conventional consolidation powder metallurgical (PM) techniques [1,2], mainly consisting of cold pressing and furnace sintering, are the most widely used, although the objective of finding a faster route has been a constant in researchers' mind.

New techniques are constantly emerging in an attempt to overcome the shortcomings of conventional processing. The direct use of electricity as a means to sinter powders has been suggested numerous times and with different approaches [3–11]. One of these techniques is electrical resistance sintering (ERS), in which heating is achieved by passing an electrical current through the powders.

The ERS process was already described in 1933 by Taylor [12] and later, in 1944, modified by Cremer [13], but its systematic study was not carried out until a few years later by Lenel [14]. Later, the ERS technique has incorporated technical variations of different importance [15,16]. Essentially, ERS involves the use of low voltage and high electrical current through a powder mass contained in an insulating die, while pressure is simultaneously applied. The medium frequency (about 1000 Hz) electrical resistance sintering (MF-ERS) technology has undergone an important evolution, making it possible to generate more heat for the same welding time.

The typical duration of the MF-ERS process is about one second. During this time, the applied current and pressure cause the densification of softened powders. This softening occurs by the temperature increase caused by the current passing through the powders, and the energy released by the Joule effect.

Among the advantages of this technique, as compared to the conventional PM route, the following should be considered: the use of relatively low pressures (around 100 MPa) to reach high densification rates, the need for very short processing times, and the possibility of sintering in air, without protective atmospheres.

WC, used in the productions of hardmetals or cemented carbides, has a high hardness and a low ductility, so that pieces of this material are usually made from powders by adding between 6% and 10% of Co as a binding agent [17]. WC–Co is mainly used in the production of tools for milling, drilling, and pressing/punching processes, although it is also used in jewelry or to make surgical instruments. Very good properties are attained by maintaining the size of WC grain size in the limit of 1 to 5 μm, although limiting the size to 0.5 μm allows achieving a unique combination of properties.

This paper details the production of hardmetal parts by electrical resistance sintering. The main aim of this work was to study the effect of current intensity and sintering time on the properties and grain size reached in different zones of the produced compacts.

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

The starting material was a commercially ultrafine WC-6 wt.% Co powder from Kyocera Unimerco (Sunds, Denmark). WC-6 wt.% Co powder was degassed (840 ◦C/0.75 h/N2H2) to avoid problems due to gases expulsion during processing, and pre-sintered (1100 ◦C/2 h/Ar) in order to increase its particle size and facilitate further processing. The carbon and oxygen content was 5.78 and 0.13 wt.%, respectively. Figure 1 shows the achieved granulated powder where the WC particle size was 262 ± 77 nm. Mean particle size of this degassed and pre-sintered powder (WC6Co) was 141 μm (Figure 1c), as measured by laser diffraction in a Mastersizer 2000 (Malvern Panalytical, Malvern, UK).

**Figure 1.** Degassed and pre-sintered WC-6 wt.% Co powder (WC6Co): Secondary electron SEM images of (**a**) the particles, (**b**) internal structure and (**c**) particle size distribution.

Compacts were prepared using a resistance welding machine (Beta 214 MF, Serra Soldadura, Barcelona, Spain), which was modified to implement the MF-ERS process. Figure 2a shows the scheme of the assembly used for MF-ERS, consisting of a 12 mm inner diameter ceramic tube (modified sodium zirconium phosphate-based ceramic, NZP) inserted in a metallic hoop. The powder mass (6.5 g in all experiences) was contained between two electrodes (thermal wafers and punches), all inside the ceramic tube. The temperature-resistant Cu electrodes (98.9% Cu, 1% Cr, 0.1% Zr) were in contact with heavy metal wafers (75.4% W, 24.6% Cu). These replaceable wafers were placed in direct contact with the powders to prevent electroerosion of the electrodes, powders sticking to the electrodes, and heat dissipation to the water-cooled electrodes due to their low thermal conductivity. In addition, to help the compacting process, the inner die wall was lubricated with a graphite–acetone suspension.

**Figure 2.** (**a**) Sketch of punches and die used in medium frequency electrical resistance sintering (MF-ERS) experiences, and (**b**) resistivity vs. compaction pressure of WC6Co powder.

The consolidation process consisted of three main stages: Cold pressing, sintering, and cooling. Compaction pressure, acting during the whole process, was 100 MPa for all the experiments. The cold-pressing and the cooling periods had a fixed duration of 1000 and 300 ms, respectively. The electrical current was only applied during the sintering stage. MF-ERS experiments were carried out with different current intensities and sintering times, always producing 12 mm diameter samples. Selected intensities were 6, 7, 8 and 9 kA, and heating times of 400, 600, 800 and 1000 ms were chosen. Intensities lower than 6 kA produced highly porous sintered samples. Intensities of 10 kA and above, combined with prolonged sintering times (1000 ms and longer), caused compacts welding to the wafers.

The previous description of the processing steps makes evident the importance of the relationship between the electrical resistivity and the compaction pressure, Figure 2b, because an adequate resistivity is required to allow the current passing at the time that causes the powders heating. Electrical resistance measurements were carried out by using a four-point probe with electrodes separated by 2 mm and a Kelvin bridge (micro-ohmmeter), by performing two measurements with inverted polarities to minimize the parasitic effects. Resistivity was determined from the measured resistance value. This powder showed a sharp drop in resistivity for low applied pressure values. The drop was lower as pressure increased. The resistivity values were between 3.76 × 10−4 and 6.87 × 10−6 Ωm in the range of 0 to 100 MPa.

Quantitative metallography on scanning electron microscopy (SEM) images (software Image Pro Plus 6, Media Cybernetics Inc., Rockville, MD, USA) was used to determine the porosity in specific regions of the compacts. The hardness of MF-ERS compacts was measured with a micro durometer Duramin-A300 (Struers GmbH, Willich, Germany), with the scale HV30, as recommended by the standard for hardmetals of this composition [18].

Samples were axially sectioned and subsequently prepared by appropriate grinding up to 400 grit SiC paper, and polishing on 6 and 3 μm diamond paste. Compact macrographs were taken with a D90 camera (Canon, Tokyo, Japan), and optical microscopy was performed with an EPIPHOT 200 microscope (Nikon, Tokyo, Japan). SEM studies of powders and compacts were carried out in a Dual Beam field emission gun scanning electron focused ion beam (FEGSEM-FIB) Auriga (Carl Zeiss, Oberkochen, Germany). In the case of sintered compacts, the same FEGSEM-FIB was required to resolve the exact morphology of these fine grained hardmetals. In order to improve the image quality, MF-ERS samples were, after being manually polished, ion etched into the microscope.

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

As compared to the relatively uniform porosity distribution in conventionally sintered compacts, MF-ERS compacts were characterized by a heterogeneous porosity distribution. Produced samples had a denser core than peripheral areas, as shown in Figure 3. This non-uniformity is the consequence of a

heterogeneous temperature distribution during processing, as shown by the process simulation [19]. In MF-ERS, the heating is produced by the Joule effect in the internal part of the sample. The temperature in the exterior is lower due to the heat transfer to the punches and die (the former are refrigerated and the latter acts as a heat sink). Therefore, MF-ERS compacts densify more in the inner, usually presenting the external porous layer (the brighter area) shown in Figure 3b.

**Figure 3.** (**a**) Scheme and (**b**) metallographic section of a sintered sample showing the dot pattern where properties were studied.

In order to evaluate the homogeneity of the produced samples, properties at different zones of the specimen section have been studied. Taking into account the revolution symmetry of sintered compacts, a total of five zones distributed in a quadrant of an axial section of the compact were selected (Figure 3). In any case, measuring locations 1 to 5 were inside the zone limited by the porous outer layer surrounding the densified core of the specimen.

With the purpose of simplifying the graphical representation of the results, only the most representative zones (zones 1, 3 and 5) were shown. Zone 1 was chosen because it was located in the most extreme area of the compact, while zone 5 was in the center. Finally, zone 3 was at an intermediate distance from the previous ones.

Figure 4a shows the attained porosity, as a function of the current intensity, for 600 ms sintered compacts. As can be seen, the measured porosity values were always very low (less than 0.55 % and in general under 0.10 %). Even the porosity measured at zone 1, in the limit with the outer porous layer shown in Figure 3, was maintained at a very low value. Results reveal that the compact core (zone 5) was slightly more porous than the intermediate layer (zone 3). The reason for this was the air trapped inside the powder mass, which cannot be completely evacuated before closing the interconnected porosity, due to the very short sintering times. The porosity increases in zones 3 and 5 for the highest current intensity (9 kA) must be highlighted. The much quicker process under these circumstances accounts for this behavior.

Regarding the porosity evolution with sintering time (Figure 4b), a greater influence than with current intensity was observed. Increasing the sintering time increased the thermal transfer by Joule effect [20], causing the porosity to decrease as sintering time is extended. The same aforementioned phenomena of porosity increase for the most energetic conditions was observed for 1000 ms.

Therefore, it can be concluded that the best results, from the point of view of porosity, were achieved by sintering at not very high intensities (Figure 4a) and not very high heating times (Figure 4b), i.e., probably a good combination could result from 600 to 800 ms at 7 to 8 kA. These processing parameters also allowed a reduced outer porous layer, with a thickness lower than 100 μm. Mean density with the best conditions were therefore very near to that of commercial samples, around 14.75 g/cm3 for the studied composition.

**Figure 4.** MF-ERS compacts porosity at zones 1, 3 and 5 as a function of (**a**) current intensity (experiences carried out at 600 ms), and (**b**) sintering time (experiences with 9 kA). Dotted lines are the trend lines of the different series.

In terms of hardness (Figure 5), most of the measured values are high, between 1800 and 2100 HV30. A typical value for commercial samples with this composition is around 1860 HV. The lowest values correspond to zones 3 and 5, according to the higher porosity at these zones, although possible different microstructural characteristics in both regions should also be considered.

**Figure 5.** MF-ERS compacts hardness at zones 1, 3 and 5 as a function of (**a**) current intensity (experiences carried out with 600 ms), and (**b**) sintering time (experiences with 9 kA). Dotted lines are the trend lines of the different series.

The sintering conditions resulting in compacts with the best hardness values were between 600 and 800 ms, for intensities between 7 and 8 kA; that is, intervals coincident with those previously indicated for porosity.

Grain size is one of the most important factors in controlling materials properties. Mechanical strength, toughness, and hardness are important engineering properties that are strongly influenced by this parameter. For this reason, the tungsten carbide size distribution, along with Co interparticle free path, have been characterized in detail.

Measurement of carbide size was conducted according to the corresponding ISO standard [21]. As stated in the standard, a minimum of 200 grains of WC were measured to obtain a value as accurate as possible. This method was based on the average grain intercepts along a line across the material surface. In ultrafine grained hardmetals, as used in this study, the average grain size was so small that good quality images were difficult to obtain in conventional SEMs. However, it has been shown that a high resolution instrument which uses a field emission electron beam can provide acceptable images (Figure 6). The measurements of phase sizes were done after an ion etching (using the ion gun of the microscope) of the manually polished sample for improving the image quality.

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

**Figure 6.** Secondary electron field emission gun SEM (FEGSEM) images of WC6Co compacts sintered at 9 kA–600 ms. Images correspond to positions (**a**) 3 and (**b**) 5 of the axial section of the compact.

Figure 6 shows representative FEGSEM images of the microstructure of a 9 kA–600 ms MF-ERS compact in zones 3 and 5. Similar images have been used for the full characterization of the carbides after the MF-ERS processes carried out with different processing conditions.

The results of WC grain size measurements are shown in Table 1. Only representative values from zones 3 and 5, of samples sintered for 600 ms are shown. Mean WC grain size slightly varied with sintering current intensity, ranging between about 260 and 325 nm, although not clearly modifying carbide sizes in a particular trend. The mean value for zone 3 results was 289 nm, whereas for zone 5 it was 297 nm. Carbide size distributions are very flat, with differences in general below the standard deviation for these measurements, which as shown in Table 1 are below 50 nm. In addition, due to the smaller size of Co regions as compared to WC, the accuracy of the measurements is lower, resulting in mean values about 39 nm for zone 3 and 49 nm for zone 5.

**Table 1.** Effect of current intensity on WC mean grain size and Co mean free path at zones 3 and 5, measured along a line perpendicular to the pressing direction. Sintering time was 600 ms.