Elastoplastic and Electrochemical Characterization of xTiB₂ Strengthened Ti Porous Composites for Their Potential Biomedical Applications

Abstract

The microstructure, elastoplastic properties, and corrosive response of induced porous Ti-TiH₂ materials reinforced with TiB₂ particles were investigated. Samples were fabricated using CP-Ti Grade1, Titanium Hydride (TiH₂), TiB₂ powders (0, 3, 10, and 30 vol.%), and ammonium bicarbonate salt (40 vol.%) as a space holder. Composites were fabricated using the Powder Metallurgy technique under high-vacuum conditions (HVS) at 1100 °C. Scanning electron microscopy, X-ray diffraction, nanoindentation tests, and electrochemical assays were used to investigate the pore formation, pore distribution, phase formation, elastoplastic properties, and electrochemical behavior of the compounds, respectively. With a mean pore diameter of 50–900 µm and Young’s modulus of less than 100 GPa, which is close to the properties of human bone, the pore structures of the compounds processed here are shown to be a potential biomaterial for osseointegration. In addition, their $H/E_r$ and $H^3/{E_r}^2$ ratios for the reinforced samples are higher than those of the unreinforced sample (1.5 and 4 times higher than the unreinforced sample, respectively), suggesting a better wear resistance of the Ti-TiH₂/xTiB₂ composites. Electrochemical experiments demonstrated that the Ti-TiH₂/xTiB₂ composites exhibited superior passivation properties compared to the Ti-TiH₂ sample. Additionally, the corrosion rates exhibited by the 3 and 10 vol.% of TiB₂ samples were found to be within an acceptable range for potential biomedical applications (29.26 and 185.82 E-3 mm·y⁻¹). The elastoplastic properties combined with the electrochemical behavior place the Ti-TiH₂/3-10TiB₂ composites as potential candidates for the biomedical application of CP-Ti.

1. Introduction

Porous Ti and its alloys are commonly used in orthopedic and dental applications because of their remarkable properties and biocompatibility [1,2,3,4]. Commercial metallic biomaterials have stronger properties than human bone, causing the so-called stress-protecting effect in the interaction between them. Ti-based alloys have a lower elastic modulus compared to commercial metal implants, but these properties are still superior to those of human bones [3,4,5,6]. This has led to the development of porous, instead of totally dense, bulk materials [3,6,7,8]. Using space-holder technology seems to be a viable approach with a simple method to obtain porous compounds with appropriate physical and mechanical properties without compromising their biological response [4,6,9]. The application affects mean pore size and porosity percentage. The ideal porosity of medical implants, for instance, is between 20 and 50% with pore sizes between 100 and 400 µm [4], and it is preferred that the porosity is of the open type rather than closed porosity [4].

In addition, the cost of high-purity of starting powders is often a limiting handicap to the development of titanium alloys by powder metallurgy techniques; hence, researchers suggested the use of TiH₂ as raw powders to reduce the manufacturing costs [9,10,11,12]. The employment of a proportion higher than 3 wt.% of TiH₂ in substitution of traditional Ti powder considerably improves the sintering, resulting in a near-total density of the final product and a higher homogeneity of the manufactured alloy parts [10,11,12].

On the other hand, Ti-alloys exhibit poor tribological properties that limit their usage in applications where the wear phenomena are present. Several means have been suggested for relieving wear issues; those techniques include the reinforcement of the Ti-matrix with ceramic particles, such as nitrides, borides, or carbides [13,14]. TiB as a reinforcement particle provides excellent properties over other Ti compounds because it exhibits good mechanical and thermodynamic stability [13,15,16].

In the literature, there are numerous reports on porous titanium materials [4,6,17]. Most of them focus on biocompatibility, mechanical and physicochemical properties [2,4,17]. The extant literature on the controlled porosity of titanium-based composites and their electrochemical corrosion characteristics is, at this time, not particularly extensive. Corrosion test results of porous titanium samples with total porosities ranging from 10 to 75% have been documented in the literature [4,17]. Based on the available literature, other factors influence the corrosion parameters of porous titanium materials such as the surface area, which increases with the material’s porosity. Pore interconnectivity and pore shape also play an important role in determining the corrosion resistance of porous materials [4,6,17]. In the present study, the effects of the induced porosity and the TiB₂ reinforcement particle additions on the microstructure, elastoplastic, and electrochemical response of the Ti-TiH₂ matrix manufactured by the HVS process were investigated.

2. Materials and Methods

2.1. Materials Preparation

CP-Ti Grade1 (<45 µm particle size, Raymor AP&C), TiH₂ grade VM (<8 µm particle size, Chemetall) and TiB₂ (<10 µm particle size, Sigma-Aldrich) powders, and ammonium bicarbonate salt (NH₄HCO₃, Sigma-Aldrich, St. Louis, MI, USA, ReagentPlus^®, ≥99.0%, 7–8.5 Ph, melting point 60 °C) were used in this study. Ti and TiH₂ powders for the matrix were mixed in a 1:1 ratio. Then, amounts of 0, 3, 10, and 30 vol.% of TiB₂ particles were added as reinforcement, as described in the previous study [12]. Subsequently, the NH₄HCO₃ was used as a space holder, mixing 40 vol.% with the Ti-TiH₂/xTiB₂ to induce the porosity in the composites. The samples were formed using powder metallurgy processing, which involved three main steps:

○

Mixing. The starting powders were mixed in a sealed polyethylene container within a glove chamber. The mixture was subjected to a vacuum and argon flooding process to displace the oxidizing atmosphere of the air. This process was repeated three times. Raw powders were mixed in a turbula (WAB, Muttenz, Switzerland) at a speed of 75 rotations per minute for a period of five hours; the objective was to achieve homogeneous mixtures.

○

Compacting. To compact the previously mixed powders that were poured into a die to be compacted, a polyvinyl alcohol binder was added. Zinc stearate was used as a lubricant on the parts of the die that come into contact with the powder mixture to facilitate compaction and extraction of the green samples. A series of cylindrical green samples measuring 10 mm in height and 12 mm in diameter were obtained by subjecting the powders to a process of compaction with a constant pressure of 442 MPa. The powders were compacted using a Physical Test Solutions model FMCC-200 universal testing machine.

○

Sintering. The consolidation of green compacts was achieved through the application of sintering in a high-vacuum environment (10⁻⁶ mbar), into a BREW furnace. Samples were sintered following three steps: first, the NH₄HCO₃ was eliminated at 100 °C during 8 h; second, the dehydrogenation of TiH₂ was performed at 550 and 700 °C for 1 h, respectively; finally, composites were sintered at 1100 °C with a dwell time of 2 h. The cooling of the samples was carried out inside the furnace.

For better identification of the composites, these were called “U” for the unreinforced sample (Ti-TiH₂ sample) and “RX” for the reinforced samples (where X = 3, 10, and 30 vol.% of TiB₂).

2.2. Material Characterization

After the sintering process, samples were cut in cross-section and prepared superficially following conventional metallographic characterization to reveal their microstructure [18]. All samples were mechanically grinded using SiC sandpaper from 180 to 2500 grade and polished to a mirror surface using an alumina solution (3, 1 and 0.05 µm of particle size). After that, some of the non-etched samples were employed to investigate the structure (by XRD), porosity, nanoindentation and electrochemical essays, respectively. The etched samples were utilized for microstructural characterization. An etchant composed of 3 mL of HCl, 5 mL of HNO₃, 2 mL of HF, and 190 mL of distiller water was used [19]. The microstructures of the compounds were observed and analyzed by field emission scanning electron microscopy and energy-dispersive X-ray spectroscopy (FE-SEM EDS, TESCAN MIRA 3LMU instrument) with a voltage of 5 and 20 kV, respectively. The structural characterization of the compounds was conducted by the Empyream, PANalytical Diffractometer employing CuKα radiation with a scanning rate of 0.026° and 2Ɵ ranging from 25 to 80° at room temperature. The collected data were indexed with the aid of Jade 6.0 software (Materials Data, Inc., Livermore, CA, USA) according to the ICDD/JCPDS database to identify the phase developed on the sintering composites. Moreover, the percentages of each phase in the composites were calculated using the PearsonVII function in the Jade software, integrating the peak areas of each phase on the corresponding X-ray diffraction pattern according to Equation (1) [20,21].

$$V_{f,i}=\frac{A_i}{\sum A_i}$$

(1)

where V_f,i is the volume fraction of a specific phase, A_i is the total integral area of this phase and ∑A_i is the total integral area for all phases detected on an XRD pattern.

2.3. Relative Density

The relative density of the composites was determined by estimating the dimensions and mass of the sintered samples. Five measurements of the diameter (d), height (h) and mass (m) of each sample were made. These results were used to calculate the measured density of the samples (${\rho }_m$) using Equation (2). This density is used to estimate the relative density (${\rho }_R$), Equation (4), together with the theoretical density of the mixtures (${\rho }_T$), Equation (3), where ($f_i$) is the volume fraction of the materials used and (${\rho }_i$) is the respective density of each element.

$${\rho }_m=\frac{m}{\pi ·{\left(\frac{\sum d_n}{2n}\right)}^2·\left(\frac{\sum h_n}{n}\right)}$$

(2)

$${\rho }_T=\sum \left(f_i·{\rho }_i\right)$$

(3)

$${\rho }_R=\frac{{\rho }_m}{{\rho }_T}$$

(4)

2.4. Nanoindentation Tests

Instrumented indentation tests were performed at room temperature using a NANOVEA CB-500 instrument and a Berkovich tip to determine the elastoplastic properties of the compounds, following a part of the methodology described in the ASTM E2546 standard [22]. To obtain data on the elastoplastic properties, the recorded load-depth plots of the composites were analyzed. Prior to each test, the instrument was calibrated according to the manufacturer’s instructions using a fused silica standard at 300 mN applied load, 600 mN/min loading-unloading rate, 2 µm/min approach rate and 0.08 mN contact load. Five different areas of the samples were then tested. The specimens had already been polished to a mirror finish. In mapping mode, a 5 × 5 indentation matrix was used per zone with 60 µm spacing. The maximum load was set at 50 mN and applied at a loading-unloading rate of 100 mN/min and a speed of 2 µm/min, with a contact load of 0.08 mN. The method described by Oliver and Pharr was used to estimate the elastoplastic properties [23].

2.5. Electrochemical Assays

The electrochemical measurements of the composites, for the U and R samples, immersed in simulated body fluid (SBF [24]) at 36 ± 1 °C was studied. The evaluation was conducted in three stages. Initially, the open circuit potential (OCP) was analyzed, with a consideration of the variation in the corrosion potential over time. This was followed by the study of potentiodynamic polarization curves. Finally, electrochemical impedance spectroscopy (EIS) tests were carried out, in accordance with part of the methodology described in the ASTM G3 standard [25]. Before electrochemical testing, the sample was sealed in epoxy on copper wire, leaving an exposed surface area of 1.06 cm². The sample surfaces were mechanically grinded and polished to a mirror-like finish. The edges were sealed with silica gel after cleaning the samples. In this study, a three-electrode electrochemical cell was constructed, comprising a Saturated Calomel Reference Electrode (SCE) serving as a reference electrode, graphite acting as a counter electrode, and a sample acting as a working electrode for the electrochemical measurements employing a commercial potentiostat ZRA Reference 600 Gamry Instruments, Warminster, PA, USA with the aid of the Gamry Echem Analyst software, V6.24 for the control and data acquisition.

The corrosion potential, E_corr, was measured over time (3600 s) from the initial immersion of the electrodes in the SBF solution. The potentiodynamic polarization and the EIS assays were conducted at 2 and 1 h, respectively, following immersion.

The EIS measurements presented were tested in the frequency range from 1 × 10⁵ Hz to 1 × 10⁻² Hz. The scanning range of the potentiodynamic polarization tests was carried out in the range of −500 mV to +1500 mV at 1 mV/s. The EIS and potentiodynamic polarization data were analyzed using the Gamry Echem Analyst software. To ensure the accuracy of the measurements and their replicability, at least five electrochemical tests were conducted on clean samples for each condition, in parallel.

With the measured values of corrosion potential (E_corr) and current densities (J_corr), the polarization resistance (R_P) was graphically estimated from the respective polarization potential versus current density [25]. The slope of the linear function was obtained according to the theory of Stern and Geary, representing the value of the R_P [26]. Furthermore, from the intersection of anodic and cathodic Tafel regions using the extrapolation method, it was determined and compared with the value of the polarization resistance obtained by the graphical method, utilizing Equation (5) [25].

$$R_P=\frac{b_a{b}_c}{2.303\left(b_a+b_c\right)i_{corr}}$$

(5)

where $b_a$ is the anodic Tafel slope, $b_c$ is the cathodic Tafel slope, and $R_P$ is the polarization resistance.

The exposed area and surface morphologies of the samples following the electrochemical measurements were analyzed using FEI-ESEM QUANTA FEG-250, a scanning electron microscope with energy-dispersive X-ray spectroscopy (FE-SEM-EDS).

3. Results and Discussion

Detailed backscattered electrons images for the cross-sectional surface topographies of the porous compounds, U and R samples, are shown in Figure 1. Sample U shows small porosities due to porosity induction and microporosities in denser regions due to the processing method. In addition, it has a microstructure that is typical of equiaxial α-Ti phase. Furthermore, the reinforced samples (R3 to R30) showed an increase in porosity due to the presence of the predominant α-Ti phase microstructure in combination with TiB and TiB₂ phases [12]. Through the space-holder technique used, porosities were generated effectively in the composites, providing a complex pore morphology.

Figure 2 shows the analysis of surface porosity in the samples. The micrographs reveal that the formation of superficial porosity presents a hemispherical morphology with pore diameters ranging from 50 to 900 µm, close to the optimum values recommended in the literature [4]. The U sample exhibits a pore density with a mean diameter of 153 µm. As the concentration of reinforcing particles increases, the mean pore diameter density increases to 225.01 µm for the R3 sample, 274.24 µm for the R10 sample, and 284.95 µm for the R30 sample. Therefore, the combination of the space-holder method and the addition of TiB₂ reinforcing particles leads to an increase in both porosity and pore size in the materials studied here.

Figure 3 displays the behavior of relative density in the composites, estimated by Equation (4). The relative density of the composites follows the same pattern as the surface porosity analysis, with the addition of reinforcement particles increasing the porosity of the composites. The relative densities of the U and R3-10 composites have optimal densification values (20%–50%) [4] for effective stimulation of osseointegration. It is evident that the U sample, with 0 vol.% of TiB₂, exhibits induced porosity of approximately 40% and reaches maximum densification compared to the R samples, which was expected in this study, demonstrating the effectiveness of the process used (the dehydrogenation and the sintering process). The R3 sample displays the highest relative density (approximately 57.5%) among the R samples due to the small concentration of reinforced TiB₂ particles, temperature, TiB phase formation by a vacancy mechanism, and the dehydrogenation process [12].

X-ray diffraction patterns of the composites (U and R samples) are shown in Figure 4. The as-received titanium powder (AR-Ti) was included in this study to enable a comparison of the evolution of the α-Ti phase. In comparison with the AR Ti raw powder, the U sample displays the characteristic α-Ti phase peaks. In regard to the R samples, the heightened intensity of α-Ti peaks suggests that the matrix material represents the primary component of the samples. Additionally, the TiB and TiB₂ phases are observed in the patterns of R samples. These phases were identified according to the crystallographic database of the ICDD/JCPDS (PDF-cards #05-0700 and #35-0741 for TiB and TiB₂ phases, respectively). It can be clearly observed that the reinforcement particles and sintering temperature employed induce the precipitation of the TiB intermetallic compound from the saturated α-Ti solid solution. Furthermore, the intensity of the peaks increases with the amount of TiB₂ reinforcement particles. According to the phase diagram, Ti-B [26], the formation of the TiB phase can be attributed to the reaction Ti + TiB₂ → 2TiB. This is because the boron atoms diffuse into the Ti crystal lattice by vacancy mechanisms which are strengthened in combination with the vacancies generated by the dehydrogenation process [12]. The α-Ti peaks shifted slightly to high 2Ɵ angles compared to the Ti raw powder (AR Ti pattern). The absence of TiH₂ peaks in the composite patterns is indicative of the effectiveness of the dehydrogenation process, which resulted in the removal of hydrogen and the formation of only α-Ti peaks.

The volume fractions of the phases found in the compounds can be calculated by means of Equation (1) and the results are given in

Table 1. Volume fractions of each phase of the compounds estimated by integrated areas from the XRD resultant patterns shown on Figure 4.

Sample	α-Ti	TiB	TiB2
U	100 ±1.35	-	-
R3	93.74 ±2.30	2 ±0.33	4.26 ±0.43
R10	87.83 ±2	6.27 ±1	5.9 ±0.85
R30	74.37 ±1.80	19.73 ±1.53	5.85 ±1.69

, according to the pattern of the XRD results. The data indicate that the α-Ti phase is dominant, but its concentration decreases with an increase in the amount of reinforcing particles. As stated in the XRD results, the strengthening particles underwent a phase transformation resulting in the formation of the intermetallic TiB phase. The volume fraction of the TiB₂ phase remained between 4.26% and 5.9% for all R samples despite the decrease in particle concentration.

Likewise, EDS results reveal the predominant presence of α-Ti phase for all samples, U and R compounds, confirming the XRD results, and the complete removal of hydrogen to obtain the predominant α-Ti phase, as shown in Figure 5. In addition, it is observed that as the concentration of B increases, the size of the precipitates also increases, as reported in our previous work [12]. Finally, the R samples are characterized by the presence of boron concentration, as shown in the SEM images. This denotes the presence of unreacted TiB₂ particles, generating the formation of TiB coronas around them, for all R samples. For the maximum concentration of reinforcing particles, there is a higher presence of TiB phase due to the high concentration of TiB₂ particles, as presented in our previous work [12]. From the semi-quantitative analyses, it can be seen that the U sample contains only Ti, without any foreign element. On the other hand, the presence of the main element Ti of the compounds can be seen in the specimens with reinforcement (R3-30), together with boron, resulting from the addition of the reinforcing particles, and finally oxygen, resulting from the passivation of the specimens due to the etching agent [27].

Nanohardness (H) and reduced elastic modulus (E_r) are presented in Figure 6. It was noticed that H exhibits an inverse trend compared to that of E_r. When the TiB₂ particles were added, H was achieved linearly by 4.5 GPa for the R30 sample. On the other hand, E_r was reduced from 108 GPa for the U sample to 77 GPa for the R10 and R30 samples. Hence, the highest proportion of TiB₂ particles exhibit the highest H values and the induced porosities produced the reduction in the E_r values. The modulus of elasticity of the R10 and R30 samples has a modulus of elasticity below the modulus of elasticity of current implants and closer to that of human bone (4–25 GPa) [8], allowing these materials to be established for potential use in bone replacement implants. Moreover, in addition to the parameters H and E_r in the literature, there are a number of other factors that can be utilized to forecast the lifespan of a component [15,28], the elastic of strain to failure (H/E_r) and the yield pressure (H³/E_r²). Furthermore, the higher value of hardness and reduced elastic modulus (E_r²) give the composites better wear and resistance to plastic deformation [15,28]. As can be seen in Figure 6b, the resultant values were achieved 1.70% higher for the H/E_r ratio and 3.70% for the H³/E_r² ratio for the R30 sample in comparison with the U sample, indicating a higher wear resistance for the R samples than the U sample. Nonetheless, the ratios obtained in this investigation were found to be higher than the values reported for a CP-Ti produced by Selective Laser Melting (around 0.025 and 0.002 GPa, for the H/E_r and the H³/E_r² ratios, respectively) [29], which suggests that the U and R samples fabricated by HVS in this study exhibit superior wear behavior.

The open-circuit potential (OCP) and potentiodynamic polarization plots are presented in Figure 7a,b. The OCP was recorded by 3600 s in the SBF solution at 37 °C. As can be seen in Figure 7a when the amount of TiB₂ increased, the OCP for the R samples increased in comparison to the U sample, exhibiting a more passive response. Similarly, it was observed that the R samples demonstrate a stabilization of the potential in relatively short timescales, at around 500 s. This is a shorter timescale than that observed for the U sample, which required a longer time to stabilize (above 2500 s). Furthermore, Figure 7b illustrates that the reinforced samples exhibit enhanced corrosion potentials (E_corr) due to the anodic passivation mechanism of the TiB₂ particles. This phenomenon can be represented by the chemical equation: TiB₂ + O⁻² → Ti (BO₃) + O⁻² → TiO⁻² + SBF. However, the TiB₂ particles are a conductive ceramic material. Thus, it is not possible to reduce the current corrosion density (J_corr) demand in the electrochemical process; therefore, the R samples had higher Jcorr values, as can be seen in Figure 7b. At the same time, the polarization resistance (R_P) and corrosion rate (CR) were influenced by the porosity as a result of the TiB₂ reinforced content. The R_P decreased and CR increased when the amount of TiB₂ particles rose up to 30 vol.%, due to the increment in the current corrosion density (J_corr) and changed in the porosity percent, while with 3 vol.%, a better performance was observed. Similar behavior has been reported in the literature [17], where the CP-Ti with porosity (45–75%) exhibited a low corrosion resistance compared with the better corrosion performance in the dense CP-Ti. The electrochemical behavior properties were calculated by means of Tafel extrapolation. The results are summarized in

Table 2. Tafel extrapolation of the U and R samples, immersed in SBF solution.

Sample	RP (kΩ)	Jcorr (µA·cm−2)	Ecorr (mV/SCE)	ba (mV/Decade)	bc (mV/Decade)	CR (mm/Y) × 10−3
U	83.43 ±7.18	0.161 ±0.013	−347.61 ±7.83	63.17 ±0.93	52.37 ±7.65	1.25 ±0.15
R3	5.55 ±0.83	2.44 ±0.04	−128.58 ±8.64	74.91 ±2.95	55.32 ±3.60	29.26 ±1.79
R10	1.26 ±0.19	14.48 ±2.42	−80.83 ±3.50	67.09 ±6.45	55.70 ±4.43	185.82 ±6.93
R30	0.79 ±0.07	17.07 ±1.41	−84.66 ±2.39	65.97 ±7.04	69.37 ±5.80	240.84 ±19.82

.

The Nyquist plot of the U sample exhibited only one capacitive response, while the R samples showed two capacitive contributions represented by imperfect semicircles, as can be seen in Figure 8a, owing to the presence of the TiB and TiB₂ phases, as shown by the XRD patterns. Also, the diameter of the semicircles decreased with the amount of TiB₂ particles, as evidenced by the better electrochemical corrosion behavior of the R3 sample with respect to the other samples, which it can be appreciated in the potentiodynamic polarization plots. A similar performance was presented for [27], and for the equivalent circuit for the U sample used for the EIS fitting consisting of a single constant phase element (CPE), described in Figure 8d. On the other hand, to fit the EIS results to the R samples, another equivalent circuit was used [30] (Figure 8e) considering the outer and the inner pore’s formation, with its respective CPE, and almost the constant phase element of the native barrier of the passive film. From the Bode-phase plot (Figure 8c), the maximal phase angle in the U sample is 70°, while in the R samples, it shifts around to 45–35°, characteristic of heterogeneous interphase; then, the reinforcement and the sintered process affect the surface as a result of the formation of porous compounds, as can be observed in the SEM images. From the Bode-impedance plot Figure 8b, the reduction in the polarization resistance of the samples when the reinforced particles are greater than 3 vol.% is noted. In this study, the EIS parameters exhibit the expected behavior for the presence of a passive film generated by the TiB₂ additions; they display the conventional shape observed in such instances. The course of these parameters indicates that the R3 sample exhibits a higher impedance. The electrochemical parameters estimated from the fitting EIS experimental data are presented in

Table 3. EIS data for the U and R samples, immersed in the SBF solution.

Sample	Rs (kΩ·cm2)	Rp (kΩ·cm2)	Rsip (kΩ·cm2)	RSop (kΩ·cm2)	CPE (nF·cm−2)	n	CPE1/2wp (nF·cm−2)	n	CPEwp (nF·cm−2)	n
U	8.52 ±0.40	4.03 ±0.07	--	--	0.169 ±0.003	0.602 ±0.01	--	--	--	--
R3	3.19 ±0.22	18.78 ±2.02	1.52 ±0.14	0.014 ±0.0006	3.870 ±0.063	0.110 ±0.006	1.790 ±0.05	0.564 ±0.04	0.098 ±0.003	0.766 ±0.025
R10	2.87 ±0.04	36.88 ±4.62	2.47 ±0.17	0.049 ±0.003	13.590 ±0.044	0.662 ±0.031	6.493 ±0.06	0.552 ±0.05	0.107 ±0.001	0.815 ±0.026
R30	2.70 ±0.27	43.76 ±1.59	1.85 ±0.25	0.037 ±0.003	27.361 ±0.156	0.942 ±0.044	9.414 ±0.05	0.623 ±0.03	0.533 ±0.002	0.875 ±0.019

.

Figure 9 shows the surfaces and semiquantitative of the EDS analysis of the U and R samples by SEM under BSE mode. For all the samples, there is no evidence of severe corrosion mechanism, which indicates that the dissolution process was achieved during the anodic polarization process. Similar behaviors were found after the corrosion test in SBF medium of spark-plasma-sintered Ti-Zr-Si-B alloys [31]. The EDS analysis of the non-reinforced sample (U) shows that compared to the passivation layer formed in the internal pores, which has a higher oxygen content (20.74 wt%), the passivation layer formed on the whole surface has a lower oxygen content (5.2 and 6.03 wt%). This may be the reason why the non-reinforced sample has a higher corrosion resistance. Due to the increased porosity of the reinforced samples, the residual components of the SBF [24] medium used in this study were mainly concentrated in the pores of these samples. For this reason, chlorine and sodium in the samples R3 and R10 can be found inside the pores. In addition to chlorine and sodium, the presence of calcium was detected in sample R30. Furthermore, the oxygen content was located at the boundary of the reinforcing particles. In addition, the close-up images of the reinforcing particles show the presence of boron in the reinforced samples. This percentage decreases with increasing particle concentration due to the diffusion of the reinforcing particles to form the TiB phase from the TiB₂ phase, besides the predominant presence of the alpha phase [12]. This confirms the results obtained by XRD and EDS in the present study.

4. Conclusions

Porous TixTiB₂ (x = 0.3, 10, and 30 vol.%) composites were successfully prepared by powder metallurgy and space-holder sintering methods under high vacuum conditions with significant use of Ti raw powder. X-ray diffraction and field emission scanning electron microscopy analyses revealed the predominant presence of the α-Ti, TiB, and TiB₂ phases, with no residual TiH₂ phase, confirming the effectiveness of the dehydrogenation process. The microstructure observed was primarily α-Ti for the U sample, with the presence of residual TiB₂ particles with TiB coronas around them for the R samples.

The microporosity present in the compounds would result in an increase in the total porosity, which would consequently increase the relative density. Furthermore, the morphology and dimensions of the pores generated in the composites, in conjunction with the addition of the TiB₂ reinforcing particles, resulted in a pronounced impact on the elasto-plastic properties, nanohardness, and reduced elastic modulus. As the concentration of TiB₂ reinforcing particles increases, the nanohardness number rises and the reduced elastic modulus value falls, which could facilitate osseointegration and, as a consequence, a robust bond between the metal implant and the bone tissue. Moreover, the ratios of H/E_r and H³/E_r² increase as the concentration of reinforcing particles increases, suggesting an enhanced trend performance in wear applications.

Finally, depending on the degree of porosity, pores of different diameters predominate in the composites. This has a decisive influence on the change in corrosion resistance and the presence of corrosion residues in the materials. The electrochemical performance was found to be negatively influenced at a content of reinforced TiB₂ higher than 10 vol.% of in comparison to CP-Ti, as evidenced by the electrochemical evaluation. Consequently, it was proposed that lower contents of TiB₂ (approximately and below 10 vol.%) reinforced particles be utilized.