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
] 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
]. 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
]. 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 Ti2
Cu, 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 Ti2
] or Ti3
]. The thermodynamic modelling done in this work is based on the first description [15
], which excludes Ti3
Cu 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 1
a,b the phase fractions in the alloys with 1 wt.% Cu and 5 wt.% Cu addition are shown. Given the prerequisites mentioned, the Ti2
Cu forms at 656 °C in the 1 wt.% Cu alloy (Figure 1
a) 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 Ti2
Cu increases with Cu addition, already at 3 wt.% Cu and here at 5 wt.% Cu (Figure 1
b), 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.
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 (
), 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.
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 Ti2
Cu. 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
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 Diffraction
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 (Ti2
Cu), 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.
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 (Ti2
Cu) 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 Ti3
] 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 Ti2
Cu-transus temperatures are within 1 °C of each other and could lead to precipitation of phases according to
. For Cu concentrations below 5 wt.% Cu, the phase development will likely proceed according to according to the reaction
. When the Cu concentration is below the 5 wt.% Cu, the phase development will likely proceed according to the reaction
. 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 Ti2
Cu 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 Ti2
Cu 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 Ti2
Cu phase only forms at temperatures lower than 747 °C (Figure 1
a). 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 3
d,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
, which is in line with kinetics for active eutectoid transformations described in studies on Ti-Cu [18
]. The Τi2
Cu 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 Τi2
Cu, which is the detection limit for the X-ray diffraction technique [20
]. The 5 wt.% Cu however showed the Τi2
Cu 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 Τi2
], 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 (Τi2
Cu) 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 Τi2
Cu however were indexed and matched well to the standard crystals. EDX-maps of the crystals in SEM (Figure 5
and Figure 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 3
e). 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 4
b). 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 4
a) with Ta surrounding the crystals. These Cu-rich crystals could be the Ti2
Cu 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, Τi2
Cu 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 Τi2
Cu 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
]. This microstructure contains “globular” precipitates of a Cu-rich phase (Ti2
Cu) that seems to form between the α and β phase (Figure 3
a 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
] according to the reaction
. 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
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.