*Article* **E**ff**ect of Thermomechanical Treatments on the Phases, Microstructure, Microhardness and Young's Modulus of Ti-25Ta-Zr Alloys**

**Pedro Akira Bazaglia Kuroda 1,2, Fernanda de Freitas Quadros 1,2, Raul Oliveira de Araújo 2,3, Conrado Ramos Moreira Afonso <sup>4</sup> and Carlos Roberto Grandini 1,2,\***


Received: 10 September 2019; Accepted: 23 September 2019; Published: 30 September 2019 -

**Abstract:** Titanium and its alloys currently are used as implants, possessing excellent mechanical properties (more suited than stainless steel and Co-Cr alloys), good corrosion resistance and good biocompatibility. The titanium alloy used for most biomedical applications is Ti-6Al-4V, however, studies showed that vanadium and aluminum cause allergic reactions in human tissues and neurological disorders. New titanium alloys without the presence of these elements are being studied. The objective of this study was to analyze the influence of thermomechanical treatments, such as hot-rolling, annealing and solution treatment in the structure, microstructure and mechanical properties of the Ti-25Ta-Zr ternary alloy system. The structural and microstructural analyses were performed using X-ray diffraction, as well as optical, scanning and transmission electron microscopy. The mechanical properties were analyzed using microhardness and Young's modulus measurements. The results showed that the structure of the materials and the mechanical properties are influenced by the different thermal treatments: rapid cooling treatments (hot-rolling and solubilization) induced the formation of α" and β phases, while the treatments with slow cooling (annealing) induced the formation of martensite phases. Alloys in the hot-rolled and solubilized conditions have better mechanical properties results, such as low elastic modulus, due to retention of the β phase in these alloys.

**Keywords:** titanium alloys; biomaterial; heat treatment

#### **1. Introduction**

The first alloy developed specifically for use in humans is known as "vanadium steel (1900)", a biomaterial produced for repairing bone fractures and for fixation screws used in orthopedics [1]. Currently, elements such as chromium, tantalum, zirconium, cobalt, titanium, niobium, tungsten, iron and nickel are also used as biomaterials, although the human body tolerates these elements only in small quantities [2]. Materials developed for use as biomaterials must possess certain favorable physical and chemical characteristics and must not induce inflammation, toxic reactions or allergenic symptoms in the host. They must be biocompatible, biofunctional, bioactive, bioinert and sterilizable [3].

Unfortunately, metal implants have been known to fail after prolonged use due to their high elasticity modulus compared to bone, and to their poor resistance to wear and corrosion in body fluids [4,5]. The bonding force between atoms determines Young's modulus. This bond strength is related to the crystalline structure of the material and to the distances between atoms [6,7]. It is possible to change the bond strength between atoms by thermomechanical treatment [8]. Plastic deformations are another treatment that can alter the structure of titanium, thus changing the atomic bond strength and Young's modulus [9]. The ω phase of titanium is known to have the highest modulus of elasticity, followed by the α and the α" phases; the β phase has the lowest modulus, due to its low atomic bond strength [10]. New alloys are being developed and investigated whose microstructure and mechanical properties can be altered with the addition of various elements, usually β-stabilizers [11–16].

Ti-Ta-Zr alloys are being developed for biomedical applications. Zhou et al. [17] indicate that the most promising of the Ti-Ta binary system alloys for use as a biomaterial is Ti-25Ta. Some authors assert that zirconium is a neutral element in titanium alloys, that is, its addition does not modify the structure of the material [18]. Recently, new research has shown that zirconium helps stabilize the β phase in α + β-type alloys, acting as a β-stabilizer [19–24]. Therefore, alloys of the Ti-25Ta-Zr system may be more promising as biomaterials, hence the addition of zirconium facilitates the β phase stabilization, in reason of the fact that β-type alloys tend to have a lower modulus of elasticity than alpha-type alloys. Besides aiding in the stabilization of the β phase, zirconium contributes to cost savings in the manufacture and commercialization of these alloys [25], because tantalum has a high market value and a high melting point, making production in high ladder more difficult.

The literature does not have great information about the structure, microstructure, phases and mechanical properties of Ti-Ta-Zr alloys. Yan at al [26] developed Ti-15Ta-xZr system alloys (x = 1.5, 5.5, 10.5, and 15.5 wt.%) manufactured by selective laser melting. Structural and microstructural analyses show that the alloys have the α and β phases and excellent elastic modulus values, highlighting the Ti-15Ta-10.5Zr alloy with a value of approximately 43 GPa. Biesiekierski et al. [27] developed new Ti-Ta-Zr system alloys with high tantalum and zirconium contents (Ti-45Zr-10Ta, Ti40Zr-14Ta, Ti-35Zr-18Ta, and Ti-30Zr -22Ta). The results showed that the alloys produced are of β-type and have elastic modulus like commercially pure titanium and Ti-6Al-4V (109-130 GPa). All alloys have good biocompatibility.

Ti-25Ta-10Zr alloy is the only alloy of type α'+ α" + β reported in the literature. Hardness values are above commercially pure titanium, being hardening by the solid solution. Cytotoxicity assays showed that the alloy does not have a cytotoxic character; on the contrary, the alloy showed perfect cell division, maintaining morphology, indicating good integration between material and cell [28].

The objective of this work was to analyze the influence of thermomechanical treatments, such as hot-rolling and annealing heat treatments, on the structure, microstructure, microhardness and elastic modulus of ternary alloys of the Ti-25Ta-Zr system, with zirconium content varying from 0 to 40% in weight.

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

The alloys were melted in an arc-melting furnace [22,28–32]. After melting, the ingots were submitted to a mechanical hot-rolling process at 1273 K, followed by air-cooling.

Because the microstructure of the metal undergoes changes during rolling, an annealing heat treatment was used to reduce the internal stress caused by the mechanical rolling process. The purpose of the annealing treatment was to eliminate residual internal stress and most of the imperfections present in the material and to promote the growth of the grains [33].

Two types of heat treatments were performed. After thermomechanical processing, the samples were submitted to recrystallization heat treatments at 1273 K, remaining at this temperature for 21.6 Ks, followed by slow cooling (SC) (−5 K/min), and another similar treatment was followed by water quenching (RC) (~273 K).

For structural analysis, X-ray diffraction (XRD) analysis (Rigaku model D/Max-2100PC) was used to obtain the diffractograms adopting the powder method, with Cu-Kα radiation, 20 mA current, potential 40 kV, for 3.2 s, range from 20◦ to 100◦, and step of 0.02◦, in the fixed time mode.

Before microstructural analysis, samples were submitted to a standard metallographic preparation. Images were obtained using an optical microscope (Olympus model BX51M), a scanning electron microscope (SEM) (Carl Zeiss model EVO-015) and a transmission electron microscope (TEM) (FEI Tecnai G<sup>2</sup> F20). TEM analysis was performed using a FEI Tecnai G2 F20 200 kV microscope with energy-dispersive X-ray spectroscopy (EDS).

Hardness measurements were performed in a microhardness tester (Shimadzu, model HMV-2). In each sample, 20 indentations were made with a 25-g load for 60 s [34]. For dynamic elastic modulus, E (GPa) measurements, the impulse excitation technique (ATCP, Sonelastic®) was used [35].

#### **3. Results**

#### *3.1. Phase Identification*

The X-ray diffractograms and optical and SEM micrographs for Ti-25Ta-Zr system alloys after hot-rolling are shown in Figure 1; after annealing and with slow cooling in Figure 2; and in Figure 3, the results for samples subjected to annealing with fast cooling, are presented. As shown, in system alloys Ti-25Ta-Zr the crystalline structure was sensitive to thermomechanical treatments.

#### *3.2. Mechanical Characterization*

Figure 4 shows the hardness values for Ti-25Ta-Zr system alloys, performed in the samples after each processing condition included in this study. An anomaly for the hardness of Ti-25Ta-30Zr alloy can be observed after annealing with rapid cooling. This anomaly can be associated with the retention of phase [36,37].

To verify the presence of ω phase, TEM measurements were made in the Ti-25Ta-30Zr alloy after annealing with rapid cooling, and the results, presented in Figure 5, show the bright (part (a)) and dark (part (b)) field images, the diffraction pattern (part (c)) and a high-resolution image of the selected area (part (d)). The presence of ω phase can be clearly observed [33].

Figure 6 shows the values of Young's modulus for Ti-25Ta-Zr system alloys, under all conditions investigated in this study. It can be observed that the values of elastic modulus decrease with the zirconium amount for the hot-rolled and annealed with rapid cooling samples. For the annealed with slow cooling, an increase of elastic modulus with the zirconium amount was observed.

**Figure 1.** X-ray diffraction patterns and optical and SEM micrographs for Ti-25Ta-xZr system alloys after hot-rolling.

**Figure 2.** X-ray diffraction patterns and optical and SEM micrographs for Ti-25Ta-xZr system alloys after annealing heat treatment with slow cooling.

**Figure 3.** X-ray diffraction patterns and optical and SEM micrographs for Ti-25Ta-xZr system alloys after annealing heat treatment with rapid cooling.

α

ω

α α

α

β

**Figure 4.** Vickers microhardness values of Ti-25Ta-Zr system alloys after hot-rolling and heat treatments, annealing with slow (SC) and rapid cooling (RC).

**Figure 5.** TEM images for Ti-25Ta-30Zr alloy after annealing with rapid cooling: bright-field (**a**), dark-field (**b**), diffraction of the selected area (**c**) and high-resolution image (**d**).

**Figure 6.** Young's modulus for Ti-25Ta-Zr system alloys, in all studied conditions.

#### **4. Discussion**

It can be observed, in Figure 1, that Ti-25Ta and Ti-25aTa-10Zr alloys show only peaks of α" phase in the hot-rolled condition. For the Ti-25Ta-20Zr alloy, the diffractograms show the coexistence of peaks associated with α" and β phases, but those of β phase have lower intensity. Ti-25Ta-30Zr and Ti-25Ta-40Zr alloys have peaks that are associated with α" and β phases, too. It is evident that the zirconium acted as a β-stabilizing element, hence its presence increases the β phase precipitation.

After the annealing heat treatment with slow cooling, the orthorhombic α" phase changed to the hexagonal close-packed structure, α phase, and the β phase remained in the structure of the material. It can be clearly observed that heat treatment modified the phases of the alloys. Since the annealing was carried out at a temperature above β-transus and cooled slowly, it was provided the necessary thermodynamic conditions for the stabilization of the phases at high temperatures, if they change to stable phases at room temperature, in the case α"→α [38].

The alloys submitted to the annealing treatment with rapid cooling presented the metastable orthorhombic α" phase. In addition, alloys with a zirconium content above 20 wt% showed a more intense β phase peak at θ ~ 68◦, indicating higher fraction and greater stabilization of bcc structure. It is stated in the literature that rapid cooling of titanium alloys from the β field can induce the formation of the metastable α" martensite since the fast cooling hinders the atomic rearrangement (retaining the β phase) and produces distortions in the crystalline structure [33]. In addition to detection of α" and β phases, in alloys with 0, 10, 20 and 30 weight percent zirconium, an ω phase peak can be visualized at θ ~ 78◦. In the Ti-25Ta-40Zr alloy it was not possible to observe ω phase peaks, showing that zirconium acted as an omega phase suppressor.

Combining the results of XRD, a typical martensitic structure can be seen in Figure 1 to Figure 3, α" phase, and grains deformed due to plastic deformation in the Ti-25Ta, Ti-25Ta-10Zr and Ti-25Ta-20Zr alloys, in the condition after hot-rolling. After the annealing, there was recrystallization of the material leading to homogenization of the structure and grain size. In these alloys, after recrystallization with slow cooling, an α-type lamellar structure can be seen, with needles emerging from the grain

boundaries and thinner needles parallel to the thick needles. In the annealed condition with rapid cooling, there was an increase in grain size, but the martensites phases are more refined, that is, the intra-grain needles are finer with this type of treatment. The images indicate that the cooling rate influences the microstructure in the material—slow cooling induces the formation of the α phase, and rapid cooling promotes the formation of the orthorhombic α" phase.

In Ti-25Ta-30Zr alloy after hot-rolling, the α" and β phases coexist; the micrographs exhibit an equiaxial structure of very fine needles distributed in some β-type distorted grains. After annealing with slow cooling, there was an increase in grain size and the α phase needles were within the grain boundaries. In annealing with water cooling, grains of the β-matrix can be observed, with small α" phase precipitates.

In the alloy with 40 wt% of zirconium, it is possible to observe small grains of the material with morphology characteristic of the β phase, in the after hot-rolling condition. The annealing treatments promoted the growth of these grains. Typical structures of the martensites are difficult to see in micrographs, although the x-ray diffractograms showed peaks of the α 'and α" phases.

A major problem in the manufacture of titanium prostheses is the mechanical conformation. A biomedical material must be easily shaped to facilitate its handling to produce screws, plates, and other shapes. The microhardness of a material indicates whether it conforms easily (low hardness value) or is a hard, brittle material. Further, biocompatible alloys must have a hardness appropriate to their specific application, because if they have very high hardness in relation to the implanted tissue, it can result in rapid tissue wear. Thus, many studies evaluate the microhardness in titanium alloys with respect to heat treatments, processing and the microstructure [39–43].

Figure 4 shows the hardness values for each process performed in the samples of Ti-25Ta-Zr alloys included in this study. According to the literature, among the phases formed in titanium alloys, the ω phase presents the greatest hardness value, followed by the α, α', α" and β [44–46]. Other studies report that β-type alloys may have a greater hardness level than mastensitic phases α' and α", because they have a stronger solid solution due to phase stability (ω > α' > α" > β > α) [47].

As can be observed, after the hot-rolling process the alloys have an equal or lower hardness value, compared to the values of annealed alloys with slow cooling. From the X-ray diffractograms, it is possible to visualize the effect of hot-rolling, heat treatment and zirconium substitution in the structure of the material. After the hot-rolling, there is a formation of the α" phase for the alloys with zirconium up to 20%, and in those above 30% of zirconium, the β phase coexists with the α" phase. The annealing alleviated the internal stress from the mechanical process and modified the α" to the α phase in the alloys undergoing slow cooling. As mentioned above, α type titanium alloys have greater hardness compared to α" alloys. Consequently, the alloys in the hot-rolled condition tend to have a lower hardness value compared to the annealed alloys with a slow cooling rate [47].

Ti-25Ta-30Zr and Ti-25Ta-40Zr alloys annealed with slow cooling have high hardness values (~425 HV). This increase can also be explained by the presence of the α and β phases, which may be an obstacle to the dislocations motion.

From hardness measurements, it can be seen that for the lower (SC) cooling rate (5 K/s) condition, there is the formation of β + α phases for all the compositions. For water quenched (RC) condition after heat treatment, it might form precipitates of athermal ω phase due to higher hardness values (from 10 to 30%Zr) reached, but in a lower fraction than SC samples. Rapid cooling condition (RC) resulted in a greater stabilization of β phase for higher zirconium content (40%Zr), suppressing also the ω phase and decreasing the hardness value.

It was observed that the hardness of the hot-rolled alloys and the alloys subjected to the annealing treatment with fast cooling have similar values. The alloys in these two conditions have the same crystalline structure (α" and β), and therefore their mechanical characteristics are similar. Only in the Ti-25Ta-30Zr alloy (with α" and β phases), there is an "anomaly" with reference to the hardness value. This high hardness value can be an indication of the formation and highest precipitation of the ω phase

in this alloy, as the fast cooling can retain the ω phase, which has the characteristic of hardening and weakening the material.

Due to the high hardness of the Ti-25Ta-30Zr alloy subjected to rapid cooling, a more detailed analysis of the microstructure of the material was performed to see the ω phase as the enhancer. Figure 5 shows the bright and dark field images of TEM, a bright field (BF) image and the selected area diffraction pattern (SAD) showing fine martensite α" phase needles with a width of approximately 50 nm dispersed in a β phase (bcc) matrix. Nanoscale precipitates of athermal ω phase from 5 to 20 nm are visualized in the TEM micrographs; their nucleation and precipitation are located close to the α" phase structures. Such athermal ω phase formation occurs in α + β type alloys and when such alloys are subjected to rapid cooling from high temperatures. The omega phase recipients are located intra-grain in the titanium structure, in circular format, visualized in nanometer scale, according to similar results of Niinomi (2016) and Homma (2018) [48,49]. Due to the precipitation of ω phase, Ti-25Ta-30Zr alloy submitted to rapid cooling showed a high value of hardness, as indicated in Figure 4 [50,51].

Figure 6 shows the values of the dynamic elastic modulus of the samples of all Ti-25Ta-Zr alloys used in this study, under all investigated conditions. In this figure, it can be observed that Young's modulus of the alloys studied after hot-rolled condition was approximately 86 GPa for the non-zirconium alloy and reached a minimum value of 72 GPa for the alloy with the highest zirconium content (40 wt%). After hot-rolling, there is an increase in the percentage of the β phase, which decreases the elastic modulus, since the β phase has the smallest modulus value among the phases in titanium alloys. The hot-rolling process generates a high cooling rate when the sample is passed through the mill roller, retaining the β phase.

Annealing with slow cooling, by contrast, increased the value of Young's modulus with the addition of zirconium. The treatment decreased the percentage of the β phase and modified the α" to α phase due to the recrystallization of the sample. The alloys with the α phase have higher elastic modulus values than α" alloys. Annealing with rapid cooling decreased the modulus of these alloys, as this type of treatment at high temperatures and rapid cooling induces the formation of the β phase, which has the lowest modulus value in titanium alloys. The Ti-25Ta-40Zr alloy has the lowest modulus of the studied alloys of E = 60 GPa.

It can be noted by the hot-rolled condition, that increasing zirconium fraction leads to the decreasing of the elastic modulus from 86 down to 72 (GPa), due to a combination of stress-induced martensite formation and β phase stabilization with suppression of ω formation. Evaluating the elastic modulus E (GPa) for a lower cooling rate leads to a tendency like the hardness measurements, where the appearance of the α phase in high zirconium alloys leads to an increase in mechanical property values, increasing the modulus from 74 to 88 GPa. Compared with the rapid cooling condition, the modulus only just decreased significantly from 74 down to around 60 GPa for higher zirconium content (40%), due to greater stabilization of β phase.

The Ti-25Ta-xZr alloys produced in the annealed condition with fast cooling have a better modulus compared to Ti-cp and other commercial biomedical metals, such as Ti-6Al-4V, 316L and Co-Cr [12]. Among the produced alloys in this study, the Ti-25Ta-40Zr has the lowest elastic modulus value (60 ± 2) GPa, which represents twice the modulus of elasticity of human cortical bone [52].

Although the Ti25Ta-40Zr alloy is promising in the biomedical field, new mechanical property testing and corrosion analysis are still required, as the surface of metals in contact with body fluids may corrode, reducing implant longevity, and to perform biological biocompatibility to verify if the material has cytotoxic character in the human organism.

#### **5. Conclusions**

The results detailed above lead to the following conclusions.


**Author Contributions:** Investigation, P.A.B.K. and F.d.F.Q.; data analysis, P.A.B.K., F.d.F.Q., R.O.d.A., C.R.M.A. and C.R.G.; writing—original draft preparation, P.A.B.K., F.d.F.Q. and R.O.d.A.; writing—review and editing, C.R.M.A. and C.R.G.; supervision, C.R.G.

**Funding:** This research was supported by the funding agencies FAPESP (grants #2012/22.742-6 and #2015/09.480-0) and CNPq (grant #307.279/2013-8).

**Acknowledgments:** The authors would like to thank the Faculdade de Ciências de Bauru, UNESP, for the XRD and SEM measurements.

**Conflicts of Interest:** The authors declare no conflicts 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* **Investigation of Copper Alloying in a TNTZ-Cux Alloy**

#### **Lee Fowler 1, Arno Janse Van Vuuren 2, William Goosen 2, Håkan Engqvist 1, Caroline Öhman-Mägi 1,\* and Susanne Norgren 1,3**


Received: 16 October 2019; Accepted: 6 November 2019; Published: 8 November 2019

**Abstract:** Alloying copper into pure titanium has recently allowed the development of antibacterial alloys. The alloying of biocompatible elements (Nb, Ta and Zr) into pure titanium has also achieved higher strengths for a new alloy of Ti-1.6 wt.% Nb-10 wt.% Ta-1.7 wt.% Zr (TNTZ), where strength was closer to Ti-6Al-4V and higher than grade 4 titanium. In the present study, as a first step towards development of a novel antibacterial material with higher strength, the existing TNTZ was alloyed with copper to investigate the resultant microstructural changes and properties. The initial design and modelling of the alloy system was performed using the calculation of phase diagrams (CALPHAD) methods, to predict the phase transformations in the alloy. Following predictions, the alloys were produced using arc melting with appropriate heat treatments. The alloys were characterized using energy dispersive X-ray spectroscopy in scanning transmission electron microscopy (STEM-EDS) with transmission Kikuchi diffraction (TKD). The manufactured alloys had a three-phased crystal structure that was found in the alloys with 3 wt.% Cu and higher, in line with the modelled alloy predictions. The phases included the α-Ti (HCP-Ti) with some Ta present in the crystal, Ti2Cu, and a bright phase with Ti, Cu and Ta in the crystal. The Ti2Cu crystals tended to precipitate in the grain boundaries of the α-Ti phase and bright phase. The hardness of the alloys increased with increased Cu addition, as did the presence of the Ti2Cu phase. Further studies to optimize the alloy could result in a suitable material for dental implants.

**Keywords:** titanium alloy; microstructures; biomaterial; TNTZ

#### **1. Introduction**

Commercially pure titanium (CP-Ti) and the Ti-6Al-4V (Ti-64) alloy are standard materials for medical implants, but problems with both have come to light with in vivo use. The former has a lower yield strength, which for grades 1–4, ranges from 170–480 MPa, respectively [1]. The latter contains vanadium which could be toxic [2] and aluminium which has been linked to Alzheimer's disease [3]. These problems have inspired development of β-Ti [4–7] and α + β-Ti [8] alloys to replace these. Of the novel alloys to date, the Ti-Nb-Ta-Zr (TNTZ) system has shown promising properties for biocompatibility and strength comparable to Ti-64 [9–11]. Depending on alloying compositions, this alloy can be manipulated to have Young's moduli lower than Ti-64, which is an appropriate step towards achieving elasticity similar to cortical bone in future [12]. All these Ti-alloys are however vulnerable to biofilm formation and patients could require antibiotic treatments in the event of infection [13,14]. In lieu of the growing problem with antibiotic resistance [15], antibacterial alloys could be useful for the biomaterials field.

Excessive addition of Cu in Ti-Cux binary alloys could however lead to toxicity [13] or material embrittlement, which is a disadvantage for load bearing biomaterials [16]. Therefore careful microstructural design is required so that the mechanical properties can be optimized for the intended application. While these findings are descriptive of a binary alloy of Ti-Cux, and Cu alloying has been performed on Ti-13Nb-13Zr-10Cu [17], it is envisioned that a similar antibacterial ability may be engineered into other quaternary alloys, i.e. TNTZ [11].

The alloying of Cu to TNTZ could lead to a novel alloy with several advantages, but microstructural, mechanical and biological properties still require careful study for optimization. Binary systems such as Ti-Cu show clear microstructural dependence on crystal relationships, chemical phases present, chemical-migration and -ordering [18]. Donthula et al. [18] and Contieri et al. [19] in particular have described the actively driven eutectoid transformation of β-Ti to α-Ti and Ti2Cu, which elucidates why β-Ti is not found in rapidly quenched alloys of this variety [20]. In contrast, studies on TNTZ without Cu present, show β-Ti and α-Ti microstructure with metastable β-phases present [21]. These two alloy systems are micro-structurally dissimilar, which further motivates the investigation into Cu addition in the TNTZ alloy systems. For these reasons the present study aims to determine the effects of Cu addition to an existing alloy of TNTZ [11], and characterize the material.

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

#### *2.1. Computational Modelling of Alloys*

The impact of Cu to the Ti-Nb-Ta-Zr system was modelled using computational thermodynamic modelling based on the CALPHAD approach [22], using the Thermo-Calc software (Thermo-Calc software AB, Solna, Sweden) and the SSOL5 database, available from www.thermocalc.se.

Since the TNTZ alloy [11] has low additions of Zr and Nb, it is hypothesized that these will remain in the solid solution of the α and β phases, respectively. The Cu, being a, β-eutectoid stabilizer, is also expected to create a eutectoid microstructure of lathes, but with increasing Cu additions, a Cu-rich phase is predicted to precipitate. It is likely that this phase will precipitate preferentially at the grain boundaries, which could lead to embrittlement. In the literature there are two contradictory predictions of the Ti-Cu phase diagram where the first Cu-rich phase is either Ti2Cu [23] or Ti3Cu [24]. The thermodynamic modelling done in this work is based on the first description [15], which excludes Ti3Cu which is a metastable phase, based on the observations by Zhang et al. [25]. Therefore, the binary Ti-Cu system was taken from the 1996 description by Kumar et al. [23].

The results should be regarded as an initial prediction of the phases and transition temperatures. This is due to the fact that the Ta-Cu binary, the Ta-Nb-Cu ternary and the Ti-Ta-Cu ternary systems have not been thermodynamically assessed and thus are lacking in the SSOL5 database. Nevertheless, the predictions given by the calculations are useful as a starting point for alloy development and to guide the experimental work. The Ti-1.7 wt.% Nb-10.1 wt.% Ta-1.6 wt.% Zr (TNTZ) has been modelled previously [11] and gave only α and β phases. The equilibrium phases as a function of temperature were modelled for this alloy with increasing Cu additions (0 wt.% Cu, 1 wt.% Cu, 3 wt.% Cu, 5 wt.% Cu and 10 wt.% Cu). In Figure 1a,b the phase fractions in the alloys with 1 wt.% Cu and 5 wt.% Cu addition are shown. Given the prerequisites mentioned, the Ti2Cu forms at 656 ◦C in the 1 wt.% Cu alloy (Figure 1a) in thermodynamic equilibrium; However, since the mole fraction is very low it is not likely to nucleate due to kinetic reasons. Nevertheless, when the phase fraction of Ti2Cu increases with Cu addition, already at 3 wt.% Cu and here at 5 wt.% Cu (Figure 1b), the phase fraction is considerable. The modelling resulted in the phase transition temperatures given in Table 1, where transus in this case, is the temperature above which the phase is no longer stable.

**Figure 1.** Mole fraction of phases as a function of temperature for (**a**) the Ti-Nb-Ta-Zr-1 wt.% Cu alloy and (**b**) the Ti-Nb-Ta-Zr-5 wt.% Cu alloy. Composition of the alloys can be found in Table 3.


**Table 1.** Calculated β- and Ti2Cu-transus temperatures for the investigated alloys.

The wt.% Cu added to each alloy introduces a change in the Gibb's free energy for each of the predicted phases [26], and depending on the resultant driving force (Δ*Gv*), the phase development will proceed as specified according to phase reactions (Table 2). The cooling rate will furthermore determine the microstructure, where metastable martensitic phases (α' and ω) could form and change the resultant material properties during a rapid quenching.

**Table 2.** Reaction equations for stable phases and microstructural development.


#### *2.2. Production of Alloys*

Alloys of Ti-Nb-Ta-Zr-Cux were produced in the range from 0 to 10 wt.% Cu (Table 3). Pre-alloyed Ti-Nb-Ta-Zr (Sandvik AB, Stockholm, Sweden) and 99.9999% pure copper rods (365327-21.5G, Sigma Aldrich, MO, USA) were used to produce the investigated alloys. Alloys were re-melted 5 times in an arc furnace, then melted into rods in the same furnace (Series 5 Bell Jar, Centorr Vacuum industries, Nashua, NH, USA). Partial homogenisation was achieved by the turning-over of the melted alloys between each of the five melting events. Complete homogenisation was achieved by heat treatments of the alloys at 988 ◦C, which is above all the calculated β-transus temperatures, for 48 h, then 747 ◦C for 18 h followed by a rapid quench. The second temperature was chosen based on the solution treatment for the Ti2Cu. The annealing was done in vacuumed ampoules, at a pressure of 1.333 mbar, to reduce the oxygen content in the alloys. All alloys were quenched in salt brine water. Thereafter all samples were embedded in Bakelite resin (PolyFast, Stuers, Ballerup, Denmark) and cut into slices using an aluminium oxide disk (50A13, Struers) before further analysis. Metallographic preparation included grinding according to the three-step preparation developed by Vander Voort [27], which was appropriately adapted (Table 4).


**Table 3.** Nominal weight percentages of elements in the TNTZ-Cux alloys.

**Table 4.** The 3-step metallographic preparation of TNTZ-Cux. All products sourced from Struers, except H2O2, which was sourced from BASF SE (Ludwigshafen, Germany).


#### *2.3. Calorimetric Measurements of Phase Transformations*

The β-transus and phase transformation temperatures were measured by differential scanning calorimetry (DSC) using a Netzsch STA 409 CD (NETZSCH-Gerätebau GmbH, Selb, Germany). Aluminium oxide crucibles were used. The 1, 3, 5 and 10 wt.% Cu samples were chosen for these measurements, which have gone through the above-mentioned heat treatments with rapid cooling. The rate of the temperature change was 10 ◦C/min. The phase transformation temperatures were determined using the onset method.

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

X-ray diffraction patterns were recorded in the Bragg-Brentano geometry using a Bruker TWIN-TWIN diffractometer (D8 Advance, AXS GmbH, Karlsruhe, Germany) with Ni-filtered Cu Kα radiation (Kα1 = 1.540598 Å). Samples were polished to 6 μm using a diamond suspension (DiaDuo-2, Struers). Crystalline phases were studied in EVA software version 4.3 (Bruker, Billerica, MA, USA). The identified phases from the ICDD database PDF–4+ 2019 [28] included PDF# 04-003-1382 (Ti2Cu), PDF# 00-044-1294 (HCP-Ti) and PDF# 03-065-9616 (HCP Ti-Ta).

#### *2.5. Microstructural Studies*

The microstructure of the samples was studied in scanning electron microscopy (SEM), focused ion beam (FIB) and scanning transmission electron microscopy (STEM) using a Zeiss Merlin (Oberkochen, Germany), an FEI Helios Nano-Lab (Brno, Czech Republic), and a JEOL 2100 TEM/STEM (Tokyo, Japan), respectively.

The Zeiss SEM and JEOL TEM/STEM were equipped with INCA AZtec Energy Dispersive X-ray Spectroscopy systems (EDS, Oxford Instruments, High Wycombe, UK) while the SEM and FIB-SEM each were equipped with back scatter, in-lens and Everhart-Thornley detectors. The JEOL TEM/STEM was additionally equipped with an annular dark field and bright field detector (JEOL, Tokyo, Japan).

For crystallographic investigations, Transmission Kikuchi Diffraction (TKD) was done with a custom made sample holder, in a JEOL 7001F SEM instrument (JEOL, Tokyo, Japan) equipped with a Schottky FEG and electron backscatter diffraction system (EBSD, Oxford Instruments) coupled to and INCA Aztec system (EDS, Oxford Instruments).

#### *2.6. Hardness Studies*

The hardness of the alloys was measured using an EMCO Test Duravision Vickers Hardness tester (Prufmaschinen GmbH, Kuchl, Austria). The machine was calibrated with a standard Vickers sample, before testing the samples. The samples were polished to grit of P400 with silicon carbide grinding paper (Struers). The applied mass for the hardness tester was set to 9.8 centinewton for all the alloys.

#### **3. Results**

#### *3.1. Phase Calculations*

The phase transformation temperatures, as determined by differential scanning calorimetry, were 829 ◦C, 751 ◦C, 746 ◦C and 744 ◦C for the 1, 3, 5 and 10 wt.% Cu alloys, respectively. The measurements were in good agreement with the calculated values for the β-transus temperatures at 746 ◦C (10 wt.% Cu alloy) and 753 ◦C (5 wt.% Cu alloy). The discrepancy between the calculated and measured values increased at lower Cu additions.

The thermodynamic prediction of phases for the alloys is given in Figure 1a,b as a function of temperature, and predicted HCP-Ti (α) and BCC-Ti (β), and additionally Ti2Cu for Cu additions of 1% and higher. At the heat treatment temperature of 747 ◦C, the calculated mol% of the phases is given in Table 5, where the 10 wt.% Cu alloy was predicted to have no α-Ti phase present, while alloys below 5 wt.% Cu were predicted to have no Ti2Cu phase present.


**Table 5.** Calculated mol% of phases of α, β and Ti2Cu in the investigated alloys at 747 ◦C.

Given the prerequisites mentioned earlier, the Ti2Cu forms at 656 ◦C in the 1 wt.% Cu alloy (Figure 1a), which is below the annealing temperature. In addition, the predicted phase fraction at lower temperature is very small, thus the phase is not likely to nucleate on quenching due to kinetic reasons.

#### *3.2. XRD and Microstructure*

With X-ray diffraction studies on the 0 wt.% Cu and 1 wt.% Cu alloys, only the α-phase was determined to be present (Figure 2). However, SEM imaging for the same alloys (Figure 3b,c) indicated that the materials had two crystal phases present, as predicted (i.e. within the α + β region). The bright phase is most likely remaining β-phase since Ta, Nb and Cu (where Cu was given by the calculations) are β-stabilisers and the heat treatment temperatures were at 747 ◦C, thus within the α-β region.

**Figure 2.** X-ray diffraction on TNTZ-Cux alloys: (**a**) Diffraction for all alloys including references from Ti2Cu (04-003-1382), HCP- (Ti-Ta) (03-065-9616) and HCP-Ti (00-044-1294). (**b**) X-ray diffraction pattern showing the 2θ angular ranges for the alloys from 38◦–41◦ and 76◦–78◦. Note the Ti2Cu peak at 39.5◦.

In the 3 wt.% Cu alloy a third phase was predicted and observed in SEM imaging (Figure 3e). Diffraction peaks were not observed for this phase but it was assumed to be the predicted intermetallic Ti2Cu. However, a very low amount was observed at large grain boundaries (GBs), which was in line with the prediction. β-Ti was not detected in the X-ray diffractogram (Figure 2) for 3 wt.% Cu alloy either, but is present in SEM images (Figure 3e). When the Cu addition was increased, the microstructure became coarser grained with thicker and disrupted bright (β) phase lathes. This coarsening is observed when comparing the 3 wt.% Cu to the 5 wt.% Cu alloy (Figure 3e,d, respectively). Likewise the phase fraction of the bright Ta, Cu-rich phase increased from 3 to 5 wt.% Cu.

The 5 wt.% Cu alloy clearly had three phases, while only small precipitates were observed in the 3 wt.% Cu. The X-ray diffraction patterns displayed an increase in Ti2Cu phase with increase in Cu concentration from 5 wt.% Cu to 10 wt.% Cu (Figure 2b).

By comparing the 10 wt.% Cu and 5 wt.% Cu samples, it is clear that the microstructure is much coarser in the 10 wt.% Cu and that a higher volume fraction of the Ti2Cu phase is found in (Figure 3a and inset). The phase diagram for 10 wt.% Cu also predicts that the Ti2Cu phase exists, above the β-transus predicted to be at 747 ◦C and this is supported by the micrographs, which show large "globular" structures of the predicted Ti2Cu phase present (Figure 3a inset). These precipitates were found exclusively at the GBs between α-Ti and β-Ti, for all alloys (Figure 3).

**Figure 3.** SEM micrographs showing (**a**) 10 wt.% Cu alloy with Inset showing three crystal phases, (**b**) 1 wt.% Cu alloy, (**c**) 0 wt.% Cu alloy (**d**) 5 wt.% Cu alloy with 3 crystal phases shown (**e**) 3 wt.% Cu alloy with three crystal phases shown.

#### *3.3. Chemical and Crystal Phase Analysis*

A change in phase development with the precipitation of the Cu-rich phase (Ti2Cu) was observed in the 3 wt.% Cu and 5 wt.% Cu samples and thus they were the focus of further study.

The 3 wt.% Cu alloy had a microstructure similar to the lower Cu content alloys with thin lathes, but by using backscattered electron imaging, smaller precipitates were discovered at the GBs of the larger α-Ti grains (Figure 3e). These areas were studied further by preparation of focused ion beam (FIB) lamella, STEM-EDS and transmission Kikuchi diffraction (TKD). Regions of Cu-rich precipitates were observed, with adjacent crystals containing Ti and Ta (Figure 4a). The grains with the brightest contrast, which probably was β-Ti considering the heat treatment temperature, were also slightly coarser grained in the 3 wt.% Cu alloy compared to those with lower Cu content. The phases were assigned as a matrix phase of α-Ti, Ti2Cu and a bright phase, where the bright phase could not be assigned to a known crystal phase using TKD (Figure 5). The 5 wt.% Cu alloy had a microstructure of irregular lathes compared to those with lower Cu content (Figure 3). The lathes that formed were not straight-line structures as in the 3 wt.% Cu, but instead lathes disrupted by Cu-rich globules, formed along the length of the bright β-phase.

Using TKD, the Cu-rich phase and the matrix phase were designated as Ti2Cu and α (HCP-Ti), respectively (Figure 6). Assignment of a known crystal to the bright phase was challenging using TKD for this alloy as well (Figure 6). The formation of the Cu-rich phase- in 3 wt.% Cu and 5 wt.% Cu occurred selectively at the GBs between the α-Ti and the bright phase. Spectroscopic comparison of the 3 wt.% Cu and 5 wt.% Cu showed that the former alloy contained a crystal with more Cu in a "globular" shaped crystal, surrounded by a crystal with more Ta and Ti (Figure 4a). The 5 wt.% Cu contained a thin crystal enriched with Cu and Ta and surrounded by crystals of Ti with lower concentrations of Cu with Ta (Figure 4b). Using TKD coupled to EDX mapping on a different lamella, the 3 wt.% Cu alloy showed a bright phase crystal with Cu, Ta and Ti (Figure 5f–h).

**Figure 4.** STEM-EDS maps on the crystal boundary showing 3 crystal phases for (**a**) 3 wt.% Cu alloy and (**b**) 5 wt.% Cu alloy, with associated Annular dark field detector image, Cu K series map, Ta M series map and Ti K series map.

**Figure 5.** 3 wt.% Cu alloy studied using TKD and EDS. From the TKD study (**a**) displays band contrast, (**b**) a phase map and (**c**) IPF Z map of the same area with (**d**) associated pole figures. (**e**) An electron image of the area investigated with both techniques, (**f–h**) EDS maps of: (**f**) Cu K series (where X indicates Ti2Cu), (**g**) Ti K series and (**h**) Ta M series.

**Figure 6.** 5 wt.% Cu alloy studied using TKD and EDS. From the TKD study (**a**) displays band contrast, (**b**) a phase map and (**c**) IPF Z map of the same area with (**d**) associated pole figures. (**e**) An electron image of the area investigated with both techniques, (**f–h**) EDS maps of: (**f**) Cu L series (where X indicates Ti2Cu), (**g**) Ti K series and (**h**) Ta M series.

#### *3.4. Hardness*

The hardness (Figure 7) of the 0 wt.% Cu alloy (135 ± 3 Hv) was significantly lower than the 1 wt.% Cu alloy (198 ± 9 Hv, p = 0.0005), which in turn was significantly lower than the 3 wt.% Cu, 5 wt.% Cu and 10 wt.% Cu alloys (p < 0.035). No statistically significant difference in hardness was found among the 3 wt.% Cu (238 ± 15 Hv), 5 wt.% Cu (232 ± 7 Hv) and 10 wt.% Cu (238 ± 19 Hv) alloys (p > 0.975).

**Figure 7.** Vickers hardness of the TNTZ-Cux alloys ("\*" indicates statistical significance of p < 0.05, as per Tukey Anova test, while "NS" indicates no statistically significant difference).

#### **4. Discussion**

The present study investigated the addition of copper to a TNTZ alloy and its effect on the microstructure, with the scope of developing a biomedical alloy with potential antibacterial ability in future. Predictions of the stable phases in the alloy system - that are to be regarded as an initial approach to the development of TNTZ-Cux alloys-revealed 3-phases in equilibrium (Figure 1).

Comparison of the predicted β-transus temperatures to the experimentally observed values, revealed discrepancies in the data at 829 ◦C, 751 ◦C, 746 ◦C and 744 ◦C for the 1, 3, 5 and 10 wt.% Cu alloys, respectively. The measurements were in good agreement with the calculated values of 746 ◦C and 753 ◦C, for the 10 wt.% Cu and 5 wt.% Cu alloys. However, the discrepancy between the calculated and measured values increased as the Cu content decreased. The reason for the discrepancies could be the absence of the Ti-Ta-Cu system in the database. An additional cause for variance in the discrepancies could be due to the reduction in the effective Cu content, since Cu is bonded in the intermetallic (Ti2Cu) phase, which was identified by diffraction for the 5 and 10 wt.% Cu alloys. A further reason could be that the β-stabilizers of Ta and Nb are soluble in the intermetallic phases, in addition to Cu. It is also uncertain whether the metastable Ti3Cu [24] is present for the lower Cu compositions, thus further research is required.

The changing copper concentrations in the TNTZ materials also changed the phase development by affecting the "driving force" (change in Gibb's free energy), leading to one of three phase development scenarios (Table 2). For the 5 wt.% Cu alloy, the β- and Ti2Cu-transus temperatures are within 1 ◦C of each other and could lead to precipitation of phases according to β → *Ti*2*Cu* + α. For Cu concentrations below 5 wt.% Cu, the phase development will likely proceed according to according to the reaction Δ*GTi*<sup>2</sup>*Cu* > Δ*G*α. When the Cu concentration is below the 5 wt.% Cu, the phase development will likely proceed according to the reaction Δ*GTi*<sup>2</sup>*Cu* < Δ*G*α. Cu addition also affects the volume fraction of the phases that develop and thermodynamic modelling predicted that no α-phase would nucleate for the 10 wt.% Cu alloy quenched from 747 ◦C (Table 5). Predictions were also made that no Ti2Cu would nucleate in the quenched 3 wt.% Cu alloy (Table 5). A possible reason for this difference is that the quenching rates from 747 ◦C, might have been too slow, and thus the α-phase nucleated in the 10 wt.% Cu alloy. The same process could have caused the Ti2Cu to nucleate in the 3 wt.% Cu alloy. Alternatively, deviations in the furnace temperature prior to quenching might also be responsible for the nucleation of the phases. Further investigations in modelling and rapid quenching could elucidate the cause for the differences between predictions and experiments.

The 0 wt.% Cu TNTZ was heat treated at 747 ◦C, and predicted to have a microstructure consisting of α (76.4%) and β (23.6%), with a hardness of 135 ± 3 Hv. In a previous study [11] the alloy was found to be a α (50%) and β (50%) alloy with hardness of 340 HVN. The differences found were probably due to various forging treatments of the alloy in the previous study [11]. The addition of 1 wt.% Cu did not cause a third phase to precipitate, presumably due to the fact that the Ti2Cu phase only forms at temperatures lower than 747 ◦C (Figure 1a). Therefore the 0 wt.% Cu and 1 wt.% Cu alloys are confirmed as two-phased materials via diffraction (Figure 2) and microscopy studies (Figure 3).

The 3 and 5 wt.% Cu alloys both had a 3-phased (Figure 3d,e) crystal structure, even though calculation of phases at 747 ◦C predicts a 2-phase structure (α and β) for the 3 wt.% Cu (Table 5). This indicates that precipitation could have taken place below 747 ◦C prior to the quenching into salt water. Alternatively, since 3 wt.% Cu is lower than the 5 wt.% Cu, the precipitation of α might have occurred according to the reaction β + α → α + *Ti*2*Cu*, which is in line with kinetics for active eutectoid transformations described in studies on Ti-Cu [18,19]. The Ti2Cu was not observed in the XRD pattern of the 3 wt.% Cu but this could be due to the volume fraction being below 2 wt.% of Ti2Cu, which is the detection limit for the X-ray diffraction technique [20]. The 5 wt.% Cu however showed the Ti2Cu in the diffractogram (Figure 2) and these were present as "globular" and irregular crystals in the microstructure. The bright phase was observed as "globular" shaped in the 5 and 10 wt.% Cu alloys.

The standard crystal structures assigned using TKD were HCP-Ti [29] and Ti2Cu [30], since these crystals gave the most appropriate match to the experimental data. The indexing of 3- and 5 wt.% Cu in TKD gave similar chemical phases of the Ti-Ta (bright β phase), the α (HCP-Ti) phase as well as (Ti2Cu) phase. The β phase was nano-crystalline and therefore contained many internal GBs, resulting in difficulty for the computational phase assignment routines in TKD to assign these GB pixels to specific phases [31], resulting in an inability to index the crystals. The α (HCP-Ti) and Ti2Cu however were indexed and matched well to the standard crystals. EDX-maps of the crystals in SEM (Figures 5 and 6) and STEM (Figure 4) provide a view to the "globular" structure of the Cu-rich phases in the 5 wt.% Cu, while lathes of Cu-rich crystals were discerned in the 3 wt.% Cu (Figure 3e). This difference in the structure for the Cu-rich phase could be due to particle coarsening, at higher Cu concentrations, via diffusion in the matrix of the alloy [32]. The maps also indicated that the 5 wt.% Cu had thin crystals containing Ta and Cu with Ti in adjacent phases (Figure 4b). The thin crystal could be the bright phase enriched in Ta surrounded by the Ti-rich α-Ti matrix. The 3 wt.% Cu had crystals enriched in Cu (Figure 4a) with Ta surrounding the crystals. These Cu-rich crystals could be the Ti2Cu phase that was surrounded by the Ta-rich bright phase found in the study.

The 10 wt.% Cu alloy had a 3-phased microstructure (Figure 3) with α, Ti2Cu and β (bright) phase, which contradicts thermodynamic modelling predictions, where the α phase was calculated as absent (Table 5). The prediction of only β and Ti2Cu after a rapid quench is not experimentally observed in microstructural transformations of active β-eutectoid alloys, of which Ti-Cu is one such example [18,33]. This microstructure contains "globular" precipitates of a Cu-rich phase (Ti2Cu) that seems to form between the α and β phase (Figure 3a inset), similar to 3- and 5 wt.% Cu. The bright (β) phase seems to have changed to a "globular" shaped structure with the increase in wt.% Cu, similar in shape to the matrix (α) phases. While the 10 wt.% Cu was not studied in TKD, it is probable that the β (bright) phase would index the same as in the 3- and 5 wt.% Cu. All transformation reactions in these alloys containing the bright phase seem to preclude β phase formation, and it is possible that the reason could be martensitic phase transformations, on account of the rapid cooling [34].

The Cu addition to the TNTZ material drives the alloy to rapid transformations via the eutectoid reactions. For the 5 wt.% Cu and 3 wt.% Cu, the bright phase could not be indexed as α, and might be a martensitic phase. This is observed in Ti-Ta (1–10 at.% Ta) and the phase identified was orthorhombic α [34] according to the reaction β → α → α. This reaction is thought to occur exclusively at temperatures below 927 ◦C (1200 K) and involves mechanical shearing and shuffling of atoms to achieve the transformation [35]. The exact shearing planes are however presently being debated and could involve twinning shear in one or both of the planes of {332}113 [36] and 111{112} [37].

The hardness increased significantly with the addition of Cu (Figure 7). This could be a consequence of the Cu atoms dissolving into the α and β crystal phases as a solid solution mixture. The hardness of the 3 wt.% Cu, 5 wt.% Cu and 10 wt.% Cu was not significantly different but was significantly higher than 0 wt.% Cu and 1 wt.% Cu, which could indicate that the saturation limit of solid solution alloying had been reached.

#### **5. Conclusions**

In the present study, the effect of Cu addition to an existing TNTZ alloy was investigated. The alloys in the range of 3 to 10 wt.% Cu all had Ti2Cu present, while in alloys with less than 3 wt.% Cu, Ti2Cu was not observed. From 5 to 10 wt.% Cu, the alloys showed the presence of Ti2Cu in increasingly "globular" structures with increase in Cu concentration. An associated effect of increasing to Cu content from 0 to 3 wt.% Cu was that the hardness increased, but no additional increase was achieved from 3 to 10 wt.% Cu. The hardness could be a result of solid solution strengthening, but might also be affected by martensitic transformations. While the material has reasonable hardness, the potential antibacterial ability of the material requires assessment in future. Therefore further studies are envisioned for this alloy system to optimize the mechanical, antibacterial and corrosion properties for the purpose of producing a suitable antibacterial implant material.

**Author Contributions:** Conceptualization, C.-Ö.M. and S.N.; Formal analysis, L.F., A.J.V.V. and W.G.; Funding acquisition, H.E., C.-Ö.M. and S.N.; Investigation, L.F., A.J.V.V. and W.G.; Methodology, L.F., C.-Ö.M. and S.N.; Project administration, H.E. and C.-Ö.M.; Resources, H.E. and C.-Ö.M.; Supervision, H.E., C.-Ö.M. and S.N.; Validation, L.F., H.E., C.-Ö.M. and S.N.; Writing—original draft, L.F. and S.N.; Writing—review & editing, L.F., W.G., H.E., C.-Ö.M. and S.N. The supervision, planning, conceptualization, funding, project administration, resources, writing and reviewing was the responsibility of H.E, C.-Ö.M. and S.N. The initial and final writing of the manuscript, production of alloys, analysis and characterization was the responsibility of L.F. and A.J.V.V assisted with TEM lamellas, as well as STEM-EDS studies of the samples. The TKD studies and analysis of the samples was performed by W.G. All authors assisted in the writing and reviewing of the manuscript.

**Funding:** This research was funded by the Swedish Foundation for International Cooperation in Research and Higher Education grant number GA SA2017-7127 (STINT), the South African National Research Foundation grant number STNT170905261815 (NRF), the INSPIRE Scholarship of the ERASMUS Mundus program of the EU (no grant number available) as well as the Axel Hultgren Fund (no grant number available).

**Acknowledgments:** A special thanks is made to the Centre for High Resolution Transmission electron Microscopy at Nelson Mandela University and Johan Westraadt for fruitful discussions.

**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* **Development of Antibacterial Ti-Cu***x* **Alloys for Dental Applications: E**ff**ects of Ageing for Alloys with Up to 10 wt% Cu**

**Lee Fowler 1, Nomsombuluko Masia 2,3, Lesley A. Cornish 3, Lesley H. Chown 3, Håkan Engqvist 1, Susanne Norgren <sup>1</sup> and Caroline Öhman-Mägi 1,\***


Received: 29 October 2019; Accepted: 28 November 2019; Published: 3 December 2019

**Abstract:** Peri-implantitis, a disease caused by bacteria, affects dental implants in patients. It is widely treated with antibiotics, however, with growing antibiotic resistance new strategies are required. Titanium-copper alloys are prospective antibacterial biomaterials, with the potential to be a remedy against peri-implantitis and antibiotic resistance. The aim of this study was to investigate Ti-Cux alloys, exploring how Cu content (up to 10 wt%) and ageing affect the material properties. Electron microscopy, X-ray diffraction, hardness testing, bacteriological culture, and electrochemical testing were employed to characterize the materials. It was found that alloys with above 3 wt% Cu had two phases and ageing increased the volume fraction of Ti2Cu. An un-aged alloy of 5 wt% Cu showed what could be Ti3Cu, in addition to the α-Ti phase. The hardness gradually increased with increased Cu additions, while ageing only affected the alloy with 10 wt% Cu (due to changes in microstructure). Ageing resulted in faster passivation of the alloys. After two hours the aged 10 wt% Cu alloy was the only material with an antibacterial effect, while after six hours, bacteria killing occurred in all alloys with above 5 wt% Cu. In conclusion, it was possible to tune the material and antibacterial properties of Ti-Cux alloys by changing the Cu concentration and ageing, which makes further optimization towards an antibacterial material promising.

**Keywords:** titanium alloys; copper; Ti2Cu; Ti3Cu; biomaterial; antibacterial

#### **1. Introduction**

Peri-implantitis is a disease caused by bacterial population of an implant surface, which results in local inflammation of soft tissue and ultimately implant loosening and failure [1]. The loss of bone tissue is often associated with the inflammation surrounding the implants, leading to revision surgeries often being troublesome to successfully accomplish [1]. The prevalence of peri-implantitis in the general population is estimated to be between 19% [2] and 28% [3]. Despite the discrepancy in prevalence data, it is generally agreed that the prevalence is under-estimated at present [1,3]. Nevertheless, the importance of the problem is clear and increasing, considering that the dental implant industry is estimated to have reached USD 5 billion in 2018 [1]. The bacteria responsible for the disease include: *Peptostreptococcus micros, Porphyromonas ginigivalis and Treponema denticola*, among others [4,5]. The fact

that these bacteria have virulent characteristic behavior [6], coupled with the increasing problems with antibiotic resistance [7], corroborate that an effective antibacterial solution is needed.

The effectiveness of bacterial colonization lies in the creation of biofilms [8], which have the ability to survive antibiotic treatments due to a multifaceted strategy of slower growth in the biofilm, poor penetration of antibiotics, and genetic variations in the bacteria [9,10]. The gram-positive *Staphylococcus epidermidis* is a biofilm forming bacteria, and is one of the most common bacteria in implant infections [11], making the bacteria a suitable strain for testing the antibacterial ability of biomaterials.

To date, much research has been focused on various approaches to reduce the bacterial burden correlated to implants, which include: surface treatments [12], the use of photochemical reactions [13], alloying of silver to titanium [14], and recently, the alloying of copper to titanium [15–17]. The antibacterial effect of Cu additions was found to be superior compared to silver [18]. The Ag appears to be effectual only at elevated temperatures [19], whereas Cu was found—using in vitro tests—to have antibacterial ability at ambient conditions [20]. For this reason, Cu is being investigated as an antibacterial element with increased frequency in thin films [21,22], ionic form [23,24], and in alloys [17,25,26]. In particular, Zhang et al. [27] have investigated the addition of Cu in titanium and found that there are discrepancies in the literature regarding required copper content and the mechanisms causing bacterial reduction. While success in antibacterial applications has been achieved with copper, the molecular process of bacterial death is not clear. However, it is hypothesized to be caused by Cu ions creating reactive oxygen species (ROS) via the Fenton reaction (Cu<sup>+</sup> <sup>+</sup> H2O2 <sup>→</sup> Cu2<sup>+</sup> <sup>+</sup> OH<sup>−</sup> <sup>+</sup> ·OH) [24], where ROS plays a role in the prohibition of bacteria 16SrRNA replication, resulting in bacteria death [28].

While it is important to understand the biological interactions in these systems, the antibacterial effect is only one aspect of successfully functioning antibacterial biomaterials. The mechanical properties are also important and should be considered, since implants experience mechanical loading in vivo. When excessive Cu is alloyed to Ti, intermetallic compounds (Ti2Cu [29] and Ti3Cu [30]) form, which can lead to brittleness [31]. While heat treatments of Ti alloys does allow some degree of ductility to be achieved in initially brittle materials, the resultant effect on microstructure and antibacterial properties must be assessed and understood. This inter-relationship between the various choices of compositions and heat treatments, and their impacts on mechanical properties and antibacterial effects, provide an opportunity for optimization of the final alloy. Therefore, the aim of the present study was to investigate these aspects (microstructural, mechanical and antibacterial) for different Ti-Cu alloys, to contribute to the understanding of these materials, and facilitate their future use as antibacterial biomaterials. The investigation only focused on alloys in the range 0 to 10 wt% Cu, since Cu ions released in this range of alloys did not show notable toxicity effects in previous works [15,23].

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

#### *2.1. Production and Preparation of Alloys*

Alloys of titanium and copper were produced in a range from 0 to 10 wt% Cu (Table 1). Titanium grade 4 (Sandvik AB, Stockholm, Sweden), "CPTi", and 99.9999% pure copper rods (365327-21.5G, Sigma Aldrich, St. Louis, MO, USA) were used to produce the alloy mixtures. Alloys were re-melted five times in an arc furnace, then re-melted and cast into rods in the same furnace (Series 5 Bell Jar, Centorr Vacuum industries, Nashua, New Hampshire, USA). Partial homogenization was achieved by turning over the melted alloys between each of the five melting events. Complete homogenization was achieved by heat treatments of the rods at 900 ◦C for 18 h, then 798 ◦C for 24 h and then finally, for half of the produced alloys, ageing at 400 ◦C for six hours. The annealing was done in vacuumed ampoules to reduce the oxygen contamination in the alloys, at a pressure of 1.333 mbar. All alloys were quenched into salt brine water by breaking the ampoules, to allow faster cooling. The alloy samples were cut into slices for analysis using a silicon carbide disk (150s, Struers, Ballerup, Denmark), to minimize

cutting damage. Following sectioning, the samples were embedded in Bakelite resin (PolyFast, Stuers, Ballerup, Denmark) for metallographic preparation.


**Table 1.** Ti-Cu*x* samples used in the study.

The metallographic preparation of the samples included grinding and polishing for subsequent microstructural studies. The three-step metallographic procedure developed by Vander Voort [32] was adapted for the present study, further details are given in Table 2.

**Table 2.** Three-step metallographic preparation of Ti-Cu (All products sourced from Struers, except H2O2 sourced from BASF SE, Ludwigshafen, Germany).


#### *2.2. Microstructural Studies*

The microstructures of the materials were studied by scanning electron microscopy (SEM)—using a LEO1550 SEM (ZEISS group, Oberkochen, Germany) equipped with an INCA AZtec Energy Dispersive X-ray Spectroscopy (EDX) system (Oxford Instruments, High Wycombe, UK). Images were taken with an Everhart-Thornley (HE-SE2 detector, ZEISS group, Oberkochen, Germany) and backscattered electron detector (BSD). To ensure that the microstructure was discernible in the SEM, the samples were etched in Kroll's Etchant: 2 mL HF, 4 mL HNO3 and 100 mL distilled water, with etching times of 30 s using cotton swabbing [33].

The energy dispersive x-ray spectroscopy (EDX) analysis of the 5Cu798 sample was performed on a Merlin SEM (ZEISS group, Oberkochen, Germany) equipped with an Ultim-Max 100 mm2 Silicon Drift Detector (Oxford Instruments, High Wycombe, UK) at a voltage of 15 kV and 8.5 mm working distance.

#### *2.3. Hardness Studies*

The hardness of the Ti-Cu*<sup>x</sup>* alloys was measured using a Vickers Hardness tester (Duravision EMCO, Prufmaschinen GmbH, Kuchl, Austria) with a 9.8 centinewton load. The machine was calibrated before testing with a standard Vickers sample, and an average was taken from three indentations.

#### *2.4. Phase Identification*

All samples studied by X-ray diffraction were polished as specified by Vander Voort [32] (Table 2). X-ray diffraction (XRD) was performed on a D8 Advance TWIN-TWIN diffractometer (Bruker, AXS GmbH, Karlsruhe, Germany) with Cu Kα radiation (Kα<sup>1</sup> = 1.540598 Å with a Ni filter) and using the Bragg-Bretano experimental set-up. The detector system was a LynxEye XE PSD detector (Bruker, AXS GmbH, Karlsruhe, Germany). The EVA software suite (2015, Bruker) was used for phase analysis, while diffractogram plotting was done in Origin (2018b, OriginLab Corp., Northhampton, MA, USA). All crystallographic data were retrieved from the ICDD database PDF–4 + 2019 [34] and included PDF# 00-044-1294 (HCP-Ti), PDF# 00-015-0717 (Ti2Cu) [35] and PDF# 00-055-0296 (Ti3Cu).

#### *2.5. Bacterial Luminescence by Direct Contact Test*

The bacteria direct contact test has already been described [17], so only a summary is provided here. Staphylococcus epidermidis (XEN43) is a genetically modified bacterial strain that is bioluminescent due to the luxABCDE gene being bio-engineered into the bacterial genome [36,37]. Overnight bacterial inoculum of XEN43 was prepared a priori then seeded on the surface of sterile 5 mm diameter Ti-Cu*x* alloys (both aged and un-aged) and allowed to attach. Tryptic soy broth (TSB, Fluke-Sigma Aldrich, Stockholm, Sweden) was then carefully added to the test wells. Periodic measurements of luminescence were recorded for the samples on a Hidex plate CHAMELEON V (425-106, Multilabel counter, Turku, Finland), and the mean was used to determine the antibacterial rate (R) after two- and six-hours of exposure, using Equation 2 [38]:

$$\mathcal{R} = \frac{N\_{\text{control}} - N\_{\text{sample}}}{N\_{\text{control}}} \times 100\% \tag{1}$$

where *Ncontrol* = mean luminescence from the CPTi sample, and *Nsample* = mean luminescence for the Ti-Cu*x* alloys [38].

#### *2.6. Corrosion Testing*

The corrosion testing was performed on all the samples at 37 ± 1 ◦C in a phosphate buffered saline (PBS) solution (containing 8 g/L sodium chloride, 0.2 g/L potassium chloride, 1.44 g/L sodium hydrogen phosphate, and 0.24 g/L potassium di-hydrogen phosphate) maintained at a pH of 7.4.

The sample preparation for the corrosion studies included mounting the samples (CPTi and Ti–Cux) of approximately 1 cm<sup>2</sup> in area, in non-conductive epoxy resin with a copper wire soldered to the samples for electrical connection. The samples were ground to 120 grit surface finishes using a SiC paper, then rinsed with de-mineralized water and degreased in acetone.

The electro-chemical testing was performed using a computer-driven Potentiostat (AutoTafel, ACM Instruments, Cark, Cumbria, England). Two graphite rods acted as counter-electrodes and a Haber-Luggin capillary made the junction with a saturated calomel reference electrode (SCE). All potential values were with respect to the SCE. Nitrogen was bubbled continuously during testing in order to remove all the oxygen and maintain an anaerobic condition.

After each sample was immersed in the solution, the open circuit potential against time curve was recorded for up to 4 h to determine the open-circuit potential (OCP). When the potential had reached a sufficiently stable value at four hours, a cyclic polarization scan was recorded from −250 mV to 1500 mV versus the corrosion potential at a scanning speed of 10 mV/min. The scan direction was not reversed.

#### *2.7. Statistical Analysis*

The statistical analysis for the bacterial luminescence measurements was done in Origin software (2018b, OriginLab Corp., Northhampton, MA, USA) at two and six hours, using a One-way ANOVA with a Tukey HSD Post-Hoc test, and with Levene's homogeneity of variance test. The same statistical test was also performed for the hardness with the exception that a Brown-Forsythe test for homogeneity of variance was performed. All tests had a statistical significance setting of *p* = 0.05.

#### **3. Results**

#### *3.1. Microstructural Studies*

Microstructures for the alloys were compared to determine any differences due to ageing. The microstructures of the alloys (both aged and un-aged) with less than 3 wt% Cu had a single-phase structure of α-Ti (the hcp solid solution of titanium) (Figure 1). However, the 1Cu798 showed peaks for the 2θ diffraction angle of the Ti3Cu phase, at 20.9◦ and 23.4◦ (Figure 2). Those with equal to

and greater than 3 wt% Cu were all two-phased: Ti2Cu and α-Ti (Figure 1). X-ray diffraction studies confirmed these findings (Figure 2), where HCP-Ti and Ti2Cu peaks were identified, but peaks for the Ti3Cu crystals were also observed at the 2θ diffraction angles of 20.9◦ and 23.4◦ for the alloys: 1Cu798, 3Cu798, 5Cu798, and 3Cu400.

**Figure 1.** Secondary electron images of: (**a**) CPTi-798, (**b**) CPTi-400, and backscattered electron images of: (**c**) 1Cu798, (**d**) 1Cu400, (**e**) 3Cu798, (**f**) 3Cu400, (**g**) 5Cu798, (**h**) 5Cu400, (**i**) 10Cu798, and (**j**) 10Cu400.

Comparison of the 10 wt% Cu before and after ageing displayed that a larger volume fraction of Ti2Cu precipitated due to this heat treatment. While an increase in Ti2Cu with ageing was also found for the 5 and 3 wt% Cu, the volume fraction of precipitated Ti2Cu was lower. In addition, these alloys had lamellar microstructures of Ti2Cu and α-Ti (Figure 1).

#### *3.2. EDX Study of Precipitates*

Precipitates in the 5Cu798 sample were studied to determine the presence of variations in the Cu content for the individual phases in the alloy. Since it is known that the Ti2Cu and Ti3Cu crystals have an atomic % Cu concentration of 33% and 25%, respectively [29,30], point analysis of individual crystals was performed to ascertain which phase was present. For sample area 1 (Figure 3a), the atomic concentrations for precipitates 1 and 2 were 33.2 ± 1.1 and 34.1 ± 0.3, respectively (Table 3). For sample area 2 (Figure 3b), the atomic concentrations for precipitates 3 and 4 were 31.5 ± 2.6 and 24.6 ± 0.5, respectively (Table 3).

**Figure 2.** X-ray diffraction patterns of Ti-Cu*x* alloys heat-treated at 798 ◦C, some aged at 400 ◦C: blue triangles indicate HCP-Ti, red triangles indicate Ti2Cu and green triangles indicate Ti3Cu. Note the peaks for Ti3Cu at 20.9◦ and 23.4◦.

**Figure 3.** Positions of EDX spectral analyses of sample 5Cu798: (**a**) Sample area 1 with precipitate 1 (48, 49, 50) and precipitate 2 (51, 52, 53, 54, 55). (**b**) Sample area 2 with precipitate 3 (65, 66, 67) and precipitate 4 (68, 79, 80).


**Table 3.** EDX point spectral analyses of sample areas 1 and 2 (point spectral analysis correspond to points in Figure 3).

#### *3.3. Bacterial Luminescence*

At two hours, no significance in luminescence was found among the samples, except between CPTi-798 and the 1Cu400 and 3Cu400 samples, which had significantly higher values (Figure 4). However, the antibacterial rate of 10Cu400 (R = 12%) was higher than the other aged alloys at this time point. All other alloys had a negative R at this time point, indicating that more bacteria were found on these alloys than on the respective control.

For the un-aged alloys, it was found that after six-hours of exposure (Figure 5) the R for 10Cu798 (42%) was higher than that for 5Cu798 (7%). The alloys with lower Cu content still had negative R-values. The 10Cu798 alloy also had significantly lower luminescence than alloys 1Cu798, 3Cu798 and 1Cu400 (*p* < 0.022). At this time point, the 10Cu400 had lower luminescence than the samples CPTi-798, 1Cu798, 1Cu400, 3Cu798, and 3Cu400 (*p* < 0.036). Within the aged alloys, it was found that R for 10Cu400 (45%) was greater than that for the 5Cu400 (15%). As for the un-aged alloys, the 3Cu400 and 1Cu400 had negative R-values.

**Figure 4.** Mean luminescence counts (bars for standard deviation) of XEN 43 bacteria after two hours for all alloys, the negative controls were (CPTi), and the corresponding antibacterial rates are shown in black. Levene's Test: F (9,20) = 1.83, *p* = 0.123. One-way Anova: F = 3.176 (*p* = 0.015) with Tukey Test. Note: The same letters (e.g, "A" compared to "A") denotes being non-significantly different, while different letters (e.g. "A" compared to "B") denotes being statistically significantly different.

**Figure 5.** Mean luminescence counts (bars for standard deviation) of XEN 43 bacteria after six hours for all alloys, the negatives controls were CPTi, and the corresponding antibacterial rates are shown in black. Levene's Test: F (9,20) = 2.21, *p* = 0.066. One-way Anova: F = 5.697 (*p* < 0.001) with Tukey Test. Note: The same letters (e.g, "A" compared to "A") denotes being non-significantly different, while different letters (e.g. "A" compared to "B") denotes being statistically significantly different.

#### *3.4. Hardness Tests*

The effects of alloying Cu to Ti and ageing on hardness were investigated (Figure 6). The 10Cu798 alloy had the highest hardness (350 ± 12 Hv) of the alloys while the CPTi-400 alloy had the lowest (120 ± 3 HV). The increased Cu addition to the CPTi correlated with a gradual augmentation in the Vicker's hardness, except for the 10Cu798 alloy with a hardness twice that of the 5Cu798 alloy (171 ± 7 HV, *p* < 0.0001 in comparison with all other alloys). Ageing did not significantly affect the hardness of the alloys, except for 10Cu798, which was 1.9 times (*p* < 0.0001) harder than 10Cu400 (182 ± 1 HV), and CPTi798 which had a 14% (*p* < 0.021) increase in HV after ageing.

**Figure 6.** Mean Vickers hardness (bars for standard deviation) of all samples. Brown-Forsythe: F (9,20) = 1.413, *p* = 0.2476. One-way Anova: F = 368.92 (*p* < 0.001) with Tukey test. Note: The same letters (e.g, "A" compared to "A") denotes being non-significantly different, while different letters (e.g. "A" compared to "B") denotes being statistically significantly different.

#### *3.5. Corrosion Tests*

Corrosion studies were performed to determine the effect of the copper additions on the corrosion resistance, as well as the effect of ageing. The alloys were plotted in a single graph (Figure 7). The alloys aged at 400 ◦C showed a distinct corrosion profile after the anodic reaction took place. The aged alloys tended to passivate rapidly after the anodic reaction commenced. The un-aged alloys (quenched from 798 ◦C) clearly had a different corrosion profile trend, which indicated a gradual change in the passivation after the anodic reaction.

**Figure 7.** Corrosion plots for the aged (at 400 ◦C) and un-aged (quenched from 798 ◦C) alloys. The aged alloys are indicated with triangles to indicate the profile. Note: the 10Cu798 alloy was not included in the corrosion study.

#### **4. Discussion**

The present investigation focused on alloying up to 10 wt% of Cu into commercially pure titanium (a well-known implant material), where Cu has shown antibacterial potential in previous studies [17,20,38,39]. Its purpose was to determine the resultant microstructural and antibacterial characteristics of the materials, as a function of Cu addition and ageing heat treatments.

Microstructural and X-ray diffraction observations indicated no β-Ti in any of the alloys, despite a rapid quench (Figure 2). The reasons for this could be the active eutectoid transformation kinetics that drive transformation from β − Ti → α − Ti + Ti2Cu [40]. The two phases of α-Ti (HCP-Ti) and Ti2Cu were observed as predicted in phase calculations [17]. The CPTi and 1Cu–except the 1Cu798 alloy that could have Ti3Cu present−alloys were single-phase α-Ti alloys (Figures 1 and 2), while in alloys with higher Cu content (aged and un-aged) both α-Ti and Ti2Cu were found. The ageing had a major effect on the microstructure and the Ti2Cu precipitates had a larger volume fraction in the aged alloys. The ageing of the 10Cu798 alloy coarsened the Ti2Cu from small precipitates to larger lamellar microstructures (Figure 1). A similar change occurred with ageing in the other two-phased alloys (5Cu798 and 3Cu798), but with coarsening of uniform lamellae of Ti2Cu in the α-Ti phase.

α-Ti and Ti2Cu were observed by X-ray diffraction, in addition to Ti3Cu at 2θ angles of 20.9◦ and 23.4◦ for the 1Cu798, 3Cu798, 3Cu400, and 5Cu798 alloys. EDX measurements on a 5Cu798 sample confirmed that some of the precipitates had the expected 33 at% Cu for Ti2Cu [35], while at least one precipitate had compositions closer to 25 at% Cu, which is the expected composition for Ti3Cu [30]. However, the results might have been affected by analyzing small areas, which resulted in collecting the signal from the surrounding phases. Although this study did not make use of a standard reference material for the analysis, the results gave insight into the variation of composition for the precipitates with Cu in the Ti-Cu*<sup>x</sup>* alloys. It should however be noted that the precipitates with compositions closer to Ti3Cu were in lower volume fractions—for the 5Cu798 alloy—relative to the precipitates with compositions close to 33 at% Cu. This lower volume fraction could be the reason for the absence of Ti3Cu by X-ray diffraction in the 10 wt% Cu alloys, since it could have been below the detection limit for the technique. However, this was only a preliminary study on one alloy and further studies are needed to confirm the general presence of Ti3Cu in Ti-Cu*<sup>x</sup>* alloys.

As well as the microstructure, the hardness was also affected by Cu additions and ageing, and the 10Cu798 was twice as hard as the 10Cu400 alloy. The mechanism causing this could be the coarser lamellar structure in the 10Cu400 [41]. Ageing had a hardening effect on the CPTi-798 alloy as well, because the hardness was significantly higher for the CPTi-400 alloy.

The bacteria test gave an interesting result after two hours of exposure for the alloys, and 10Cu400 had an antibacterial rate R of 11%, while all the other Ti-Cu*<sup>x</sup>* alloys had higher mean luminescence than CPTi. The reason for the higher luminescence could be that the bacteria underwent a stress response, which caused a rapid increase in population [23]. Comparison of the aged alloys at two hours and at six hours gave a similar trend for antibacterial rate (R). The 10Cu400 alloys seemed to be effectual at bacteria reduction already from two hours and thereafter increased, and at six hours the R was 46%. The 5Cu798 and 5Cu400 alloys after six hours had R values of 7% and 15%, respectively, although this was not significantly different from alloys with lower Cu content. The greater R at six hours for both aged (45%) and un-aged (42%) 10 wt% Cu alloys agreed with previous findings where a higher Cu content has been reported to yield a stronger antibacterial material [17,38]. The difference in R for the 10 wt% Cu alloys at two hours could be due to the larger amounts of Ti2Cu in the aged alloy leading to higher Cu ion release rates as has been observed in studies on Ti-6Al-4V-5Cu alloys [28]. However, further studies on ion release from Ti-Cu*x* are needed to support this conclusion. Additionally, direct contact of bacteria with a Cu rich surface—such as the 10Cu400 with higher volume fraction Ti2Cu—could also have played a role in the antibacterial effect [24]. Thus, Cu ions and direct contact killing contributed to the antibacterial effect, and future studies could investigate how the amount and size of the Ti2Cu phase influences copper ion release and the contact killing of bacteria.

The corrosion results indicated that alloying with Cu and ageing affected the corrosion rates and passivation rates. Alloying and ageing induced faster passivation for higher Cu contents, and decreased the corrosion rates. These results are in agreement with the findings of Zhang et al. [38] for an aged 4 wt% Cu alloy.

A limitation to the present study is that only alloys in the range of 0 to 10 wt% Cu were investigated. It has been demonstrated that a higher Cu content yields a better antibacterial response [17], however, an excess of copper can be harmful to the human body [42]. Previous studies [15,23] have reported that the Cu ions released from alloys with up to 10 wt% Cu did not have any notable toxic effects and were therefore the focus of the present study. Moreover, only two heat treatments were studied due to a necessary delimitation of the study. There are other heat treatments that could be tested, and future studies should further optimize the alloying process. The investigation on the presence of Ti3Cu was a preliminary study and the results should not be viewed as conclusive due to only one alloy being characterized, the low number of precipitates characterized and the low resolution of the EDX technique. Another limitation is that the antibacterial ability of the alloys was only tested with one type of bacteria. However, the bacterium used, Staphylococcus epidermidis*,* is one of the most prevalent bacteria in implant infections and its use additionally allowed comparison with previous studies [23].

#### **5. Conclusions**

Both Ti2Cu and Ti3Cu were detected in Cu alloys using X-ray diffraction. A preliminary EDX study on the un-aged 5 wt% Cu alloy also indicated the presence of Ti3Cu, however with the Ti2Cu in majority. While the study of the Ti3Cu was not quantitative in this investigation, it is recommended that future works focus on determination of this phase.

The hardness was reduced after ageing of the 10Cu798 alloy, as the microstructure changed from small precipitates to coarser lamellar microstructures. The alloys had good antibacterial rates above 5 wt% Cu (aged and un-aged) after six-hours of exposure, and the two-hour exposure was probably too short a duration for bacterial reduction to occur in most of the alloys. Ageing the alloys only ensured faster antibacterial rates after 2 hours, when the concentration of Cu was greater than 10 wt% Cu. The addition of Cu to Ti and ageing increased the corrosion resistance of the alloys, which could protect the biomaterial in vivo. However, since the ageing also lowered the hardness of the higher Cu content alloys, care should be taken to avoid over-ageing, and hence softening.

In conclusion, the un-aged 10 wt% Cu alloy was considered a suitable candidate material, which provided a good antibacterial effect, with superior hardness and corrosion protection. The ideal composition and heat treatment of these materials will however depend on the specific application envisioned, and will require further optimization.

**Author Contributions:** Conceptualization, H.E., S.N. and C.Ö.-M; Formal analysis, L.F., N.M., S.N. and C.Ö.-M; Funding acquisition, L.A.C., L.H.C., S.N. and C.Ö.-M; Investigation, L.F. and N.M.; Methodology, S.N. and C.Ö.-M; Project Administration, L.A.C., L.H.C., H.E., S.N. and C.Ö.-M; Resources, L.A.C., H.E., S.N. and C.Ö.-M; Supervision, L.A.C., L.H.C., H.E., S.N. and C.Ö.-M; Validation, L.F. and N.M.; Visualization, L.F. and N.M.; Writing-original draft, L.F.; Writing-review & editing, L.F., N.M., L.A.C., L.H.C., H.E., S.N. and C.Ö.-M.

**Funding:** This research was funded by the South African-Sweden Bilateral Scientific Research Cooperation program funded by the Swedish Foundation for International Cooperation in Research and Higher Education (STINT, GA SA2017-7127) and the South African National Research Foundation (NRF funding through STNT170905261815).

**Acknowledgments:** The authors thank Victoria Sternhagen for help with SEM imaging.

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

#### **References**


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