Titanium alloys developed on the basis of the addition of cheap strong eutectoid β-stabilisers

Historically, titanium alloys have been developed disregarding the addition of eutectoid β-stabilisers as they generally lead to the formation of brittle intermetallic phases upon solidification of the alloy. However, such phenomenon can be prevented using powder metallurgy. Thus, this study considered the concurrent addition of cheap strong eutectoid β-stabilisers, namely Mn and Fe, for the development of new ternary Ti–Mn–Fe alloys, reducing the intrinsic cost of Ti alloys. It is found that the progressive addition of Mn and Fe in equal concentration enhances the densification of Ti during sintering, leading to lower amount of residual porosity, the transformation of the microstructure from purely lamellar to metastable, and the associated refinement of the microstructural features (grains and lamellae), as well as the stabilisation of a greater amount of β phase, and the formation of the metastable α″ phase. Such microstructural changes result in the strengthening (higher yield and ultimate tensile strength and hardness) and embrittlement of the alloy by changing the fundamental strain hardening mechanism of the ternary Ti–Mn–Fe alloys.


Introduction
Titanium (Ti) is the ninth most-abundant element on the planet and the fourth most-abundant structural metal, where the main mineral sources of Ti are rutile and ilmenite [1,2]. Ti is classed as a light nonferrous metal and Ti alloys primarily stand out due to two properties: high specific strength and excellent corrosion resistance. These features make Ti alloys suitable for a wide range of industrial applications, which could take advantage of their light-weighting effect and durability. Currently, the industrial adoption of Ti alloy is however restricted by its high cost derived by both its physical metallurgy and its manufacturing [3,4].
Ti crystallises at low temperatures in a modified hexagonal close packed (hcp) structure denoted as a Ti and the body-centred cubic (bcc) structure is stable at high temperatures, and it is denoted as b Ti [5,6]. The b transus temperature, 882 ± 2°C for pure Ti, is very important because processing and heat treatments are often carried out with reference to some incremental temperature above or below the b transus [7,8]. The alloying elements added to Ti are classified as neutral, a-stabilisers or b-stabilisers. The a-stabilising elements extend the a-phase field to higher temperatures, while the b-stabilising elements shift the b-phase field to lower temperatures, and neural elements have minor influence on the b transus temperature [9,10]. b-stabilising elements are subdivided into isomorphous b and eutectoid b elements. Eutectoid b elements are used in b-rich a ? b alloys or in b alloys because they are strong b-stabilisers and improve hardenability and response to heat treatment [11]. Moreover, most eutectoid b elements are cheaper than Ti and, therefore, they can be used to design now low-cost Ti alloys that are cheaper with respect to currently available ones.
Amongst the available b-stabilising elements, iron (Fe) is the cheapest because is readily available. Fe is a strong b-stabilising element, providing strong solidsolution strengthening effect, has higher density than Ti, and has high diffusivity in Ti [12]. The last two aspects justify why Fe was mainly ignored as alloying element in the production of Ti alloys via casting but has been considered to develop low-cost powder metallurgy Ti alloys [13]). For example, Chen et al. [14] analysed the effects of different post-sintering cooling conditions on the properties of binary Ti-xFe (x = 3, 5 and 7 wt%) alloys fabricated via cold press and vacuum sintering at 1250°C. Fe has also been considered in ternary and quaternary Ti alloys. For example, Kuroda et al. [15] investigated the mechanical behaviour and corrosion response of the newly developed Ti-8Fe-8Ta and Ti-8Fe-8Ta-4Zr alloys for biomedical applications, and Ozkan Gü lsoy et al. investigated the production of Ti-Fe-Zr alloys via injection moulding [16] and die compaction [17].
Manganese (Mn) has been considered as an alloying element in Ti alloys because it is a strong b-stabilising element and because of its lower cost/cytotoxicity compared to V, strong solid-solution strengthening effect, and higher availability compared to other alloying elements [18,19]. Therefore, Ti-Mn alloys are being developed for biomedical devices. For examples, Fernandes Santos et al. [20] analysed the fabrication of binary Ti-xMn (x = 8, 9, 12, 13, 15 and 17 wt%) alloys by means of injection moulding plus vacuum sintering at 1100°C. The addition of a third alloying element such as Mo [21] or Cu [22] has also be analysed targeting to enhance the mechanical behaviour and functionalise the material with antibacterial activity, respectively.
From literature, Mn and Fe were individually considered to manufacture both a ? b and b Ti alloys, but no major studies understanding their effect on the properties of a ? b ternary Ti-Mn-Fe alloys is available. To the best knowledge of the authors, the only study considering ternary Ti-Mn-Fe alloys for structural applications is the work of Ikeda et al. [23], who analysed the influence of Fe content of Ti-Mn-Fe b Ti alloys. Consequently, the aim of this work is, thus, to expand the knowledge on the manufacturability and properties of a ? b ternary Ti-Mn-Fe alloys obtained through the blending of elemental powders, cold pressing and vacuum sintering. This study analysed the level of relative density achievable, the evolution of the microstructure with the progressive addition of a higher amount of alloying elements, and the relationship between the previously mentioned factors and the mechanical properties of sintered ternary Ti-Mn-Fe alloys.

Experimental procedure
The powders purchased for the current study included a hydride-dehydride Ti powder with particle size lower than 75 lm (GoodFellow Ltd., oxygen = 0.23 wt.%, carbon \ 100 ppm), a comminuted Mn powder with particle size lower than 45 lm (Sigma-Aldrich Ltd., purity [ 99.0%), and an atomised Fe powder with particle size lower than 10 lm (GoodFellow Ltd., carbon \ 200 ppm). Figure 1 shows that the shape of the powder particles is irregular, angular, and spherical, respectively, for the Ti, Mn and Fe powders.
The ternary Ti-Mn-Fe alloys compositions analysed in this study were: Ti-0.5Mn-0.5Fe, Ti-1Mn-1Fe, Ti-3.5Mn-3.5Fe, and Ti-5Mn-5Fe. For that, the correct amount of each powder was loaded into a mixer and blended altogether using a rotational speed of 45 rpm for 30 min. From the literature about the manufacturing of Ti-based alloys through powder metallurgy [24][25][26], it was decided to press the powder blends by applying a uniaxial pressure of 600 MPa for their subsequent vacuum sintering. The latter was done heating the samples at 10°C/min until reaching the maximum temperature of 1300°C, which was kept for 120 min, under a vacuum level of approximately 10 -5 mbar. The samples were finally furnace cooled under vacuum.
The density of the 40 mm cylindrical pressed billets was calculated by means of the mass/volume ratio for which the mass was measured through a high precision analytical scale and the dimensions were obtained using a digital calliper. The density of the sintered samples was obtained by means of Archimedes' method measurements. To calculate the relative density values, and therefore the amount of porosity, the rule of mixture was used to compute the theoretical density of the alloys. The hardness of the sintered samples was quantified considering triplicate HRA measurements.
X-ray diffraction (XRD) 30°-80°patterns of the sintered alloys were obtained scanning the materials at a rate of 0.013°with dwell time of 0.5 s by means of Philips X'pert with Cu K a radiation. For microstructural analysis, the samples were cut by means of electrical discharge machining (EDM), metallographically prepared (emery papers grinding plus OPS polishing) and, subsequently, etched by means of a Kroll reactant (2 vol% HF ? 4 vol% HNO 3 in distilled water) to analyse the microstructure via an Olympus BX-60 optical microscope and a Hitachi S-4700 SEM.
Tensile test pieces with dog-bone geometry were also cut through EDM and their surfaces were ground to standardise the surface finishing and eliminate any potential effect from the preparation of the samples. Tensile tests were performed by means of an Instron 33-R-4204 machine. The strain rate used was 5Á10 -3 1/s, a minimum of five samples were tested, and the elongation was measured by means of a mechanical extensometer. The obtained stressstrain curves were then processed in order to quantify the strain hardening rate of the alloys on the basis of the true stress and true plastic strain. Figure 2 shows the variation of the relative density of the ternary Ti-Mn-Fe alloys as pressed and after sintering as well as the variation of the hardness as a function of the amount of alloying elements added. The green density (85.7 ± 0.25) initially slightly decreases and then increases as more alloying elements are added, whereas the sintered density (94.1-95.7%) continuously increases meaning that there is an increasing gain in density (8.5-9.6%) as the  total amount of alloying elements of the ternary Ti-Mn-Fe alloys increases. The hardness of the sintered ternary Ti-Mn-Fe alloys also increases continuously with the amount of alloying elements.

Results
The XRD patterns of the sintered ternary Ti-Mn-Fe alloys shown in Fig. 3 indicate that only the a phase was detected in the Ti-0.5Mn-0.5Fe alloy whilst both the a and b phases were present when the amount of Mn and Fe was increased. Furthermore, it can be noticed that the b phase is the predominant phase in the Ti-5Mn-5Fe alloy, the metastable a 00 phase was also detected, and no peaks of the elemental Mn and Fe powder were found indicating the complete dissolution of the alloying elements and the achievement of a homogenous chemical composition. Figure 4 shows representative optical and SEM of the sintered ternary Ti-Mn-Fe alloys where it can be seen that the alloys are characterised by the classical lamellar microstructure composed of a-grain boundaries (or prior b grains) and a ? b lamellae [27]. The only exception is the Ti-5Mn-5Fe alloy, whose microstructure is composed of fully stabilised b grains and a 00 needle-shaped grains primarily found at the grain boundaries, even if a small proportion is also present with the b grains. As the amount of alloying elements added is increased, the size of the a-grain boundaries decreases, the interlamellar spacing between the a ? b lamellae decreases, resulting in a significant refinement of the microstructure. Residual pores are clearly visible in the microstructure of the sintered ternary Ti-Mn-Fe alloys, where the majority of the pores are spherical in shape, are isolated, and are situated at the grain boundaries.
Representative engineering stress-strain curves and the average tensile properties of the ternary Ti-Mn-Fe alloys are presented in Fig. 5, where it can be seen that the materials show both elastic and plastic deformation prior to failure with the exception of the Ti-5Mn-5Fe alloy, which is characterised by only elastic deformation and fracture without undergoing any plastic deformation. As the amount of alloying elements increases, the average yield and ultimate tensile strength increase reaching the maximum values of 818 ± 5 MPa and 842 ± 12 MPa for the Ti-3.5Mn-3.5Fe alloy and, consequently, the strain decreases. However, despite the higher amount of alloying elements, the Ti-5Mn-5Fe alloy has low strength due to its brittle failure. Figure 6 shows representative micrographs of the fracture surface of the sintered ternary Ti-Mn-Fe alloys, where it can be seen that in the case of the Ti-0.5Mn-0.5Fe and Ti-1Mn-1Fe alloys failure occurred intergranularly at the lamellae, which withstood a significant amount of plastic deformation. Large dimples, of the size of the a lamellae, as well small dimples typical of ductile failure are found along the a-grain boundaries. The fracture surface of the Ti-3.5Mn-3.5Fe alloy is composed of large dimples primarily originated from the residual pores and small dimples as well as intergranular fracture at the agrain boundaries and a ? b lamellae boundaries, and tear ridges along the a-grain boundaries. In the case Ti-5Mn-5Fe Ti-3.5Mn-3.5Fe Ti-1Mn-1Fe Ti-0.5Mn-0.5Fe of the Ti-5Mn-5Fe alloy, failure of the material occurred transgranularly within the fully stabilised b grains and river patterns are found inside the grains, especially where a 00 needle-shaped grains were present. Although brittle in nature, a small amount of plastic deformation is found in the fracture surface, along the b grains, of the Ti-5Mn-5Fe alloy. Representative optical and SEM micrographs, respectively, of the sintered ternary Ti-Mn-Fe alloys: ab Ti-0.5Mn-0.5Fe, c-d Ti-1Mn-1Fe, e-f Ti-3.5Mn-3.5Fe, and g-h Ti-5Mn-5Fe.

Discussion
Ternary Ti-Mn-Fe alloys were formulated and fabricated via the blended elemental approach because the composition can easily be changed as the powder blend is made out from the elemental powders. From  Fig. 2a), which shows the variation of the relative green density of the samples, it can be seen that, even though there are small variations, the relative green density is not significantly affected by the total amount of alloying elements added as the variation is fairly limited, max 0.7% between the lowest and highest value. The green density values, which are related to the compressibility of the specific powder blend, thus indicate that the simultaneous addition of the Mn powder with angular morphology and particle size lower than 45 lm and of the Fe powder with spherical morphology and particle size lower than 10 lm does not significantly affect the response of the elemental Ti powder during the application of the uniaxial pressure for compaction. This means that the combination of powder morphology, particle size, and intrinsic hardness of the material counterbalance themselves leading to a fairly constant, or slightly increased, compressibility. Sintering of the ternary Ti-Mn-Fe alloys powder compacts leads to an increase in relative density which monotonically increases with the amount of alloying elements, therefore, also resulting in a constant increase in the gain of density between the green and the sintered samples (Fig. 2a). Such increase indicates that the addition of Mn and Fe actually favours the densification of the ternary Ti-Mn-Fe alloys. This is primarily due to the high diffusivity of Fe in Ti and the

Strain [%]
Ti-1Mn-1Fe Ti-5Mn-5Fe Ti-3.5Mn-3.5Fe Ti-0.5Mn-0.5Fe small size of the Fe powder particles, which enhances the sintering driving force, as it is related to the surface areas; although Mn has also been reported to increase the densification of Ti alloys [22]. The relative density values reported in Fig. 2a) are comparable to those of other Ti alloys processed via the uniaxial cold pressing plus sintering powder metallurgy route [28][29][30][31], resulting in the decrease of the total amount of porosity (4.3-5.9%) left by the sintering process. The reduction in the residual porosity is undoubtedly favourable to improve the hardness of the sintered ternary Ti-Mn-Fe alloys; however, the very steep increase in hardness within a small relative density range shown in Fig. 7a) indicates that the addition of the alloying elements, with the resulting changes in microstructural features, is the main factor responsible for the significant increase in hardness shown in Fig. 2b). From the results of the XRD analysis performed on the sintered ternary Ti-Mn-Fe alloys (Fig. 3), the b phase was not detected in the Ti-0.5Mn-0.5Fe alloy due to its small amount, which is most likely below the detection limit of the equipment. The addition of the alloying elements stabilises the b phase, which is actually detected in the other ternary Ti-Mn-Fe alloys and is the primary phase composing the Ti-5Mn-5Fe alloy. Therefore, as the amount of Mn and Fe added to Ti increases, more b phase is stabilised in the microstructure and the metastable a 00 is also formed, indicating the transition from a a ? b alloy to a metastable b alloy. Apart minor discrepancies, the phases identified in the XRD patterns concord with the results of the microstructural analysis. Specifically, from Fig. 4a), the initial addition of 0.5% of Mn and Fe to Ti leads to the stabilisation of a small amount of b phase, thus, leading to the creation of a near-a Ti alloy with coarse prior b grains (Fig. 4b), whose amount of stabilised b phase cannot be detected by means of XRD analysis. Further dissolution of the Mn and Fe powder particles within the Ti matrix leads to a greater amount of stabilised b phase (Fig. 4c), reducing the prior b grains size and the interlamellar spacing (Fig. 4d), resulting in the Ti-1Mn-1Fe alloy being an a ? b Ti alloy characterised by the typical lamellar structure composed of a-grain boundaries and a ? b lamellae. This microstructure, which is also found in the Ti-0.5Mn-0.5Fe alloy, forms during the slow cooling of Ti alloys comprising a relatively small amount of b-stabiliser elements upon crossing the b transus of the alloy. The addition of 3.5 wt% of Mn and Fe also results in the formation of a lamellar structure (Fig. 4e), but with slightly smaller prior b grains and significantly finer interlamellar spacing (Fig. 4f) due to the increased amount of stabilised b phase. The combination of microstructural and XRD analyses indicates that the Ti-3.5Mn-3.5Fe alloy is rather a metastable b Ti alloy, as the metastable a 00 phase is present in the XRD pattern, although of the resemble with a lamellar microstructure. In the case of the Ti-5Mn-5Fe alloy, the microstructural analysis confirms that this is a metastable b Ti alloy composed of b grains and needle-like grains of the metastable a 00 phase (Fig. 4h), which remains as residue of the incomplete stabilisation of the b phase from the total amount of bstabilisers added (Fig. 4g). Further from the microstructural analysis results shown in Fig. 4, residual porosity is present in the microstructure of the sintered ternary Ti-Mn-Fe alloys, in agreement with the relative density results (Fig. 2). The total amount of residual pores decreases but their average size seems to slightly increases with the amount of alloying elements, where most of the pores are located at the grain boundaries. The location and coarsening of the pores by coalescence are expected once the alloys reach the last stage of sintering [5].
Ti alloys can be developed and classified using different criteria including: (1) the identification of the phases composing the material manufactured using a process that guarantees near-equilibrium cooling conditions [32,33], (2) through the quantification of the molybdenum equivalent (MoE) parameters, which takes into account the stabilisation strength of each alloying elements added [9,34], or (3) through the concept of ''bond order/d-orbital energy'' (Bo-Md) maps [35,36]. Aiming to clarify the nature of the sintered ternary Ti-Mn-Fe alloys, Table 1 and Fig. 8, respectively, report a summary of the phases found and of their features as well as the predicted alloy type using the previously mentioned criteria. It can be seen that there is a fair good agreement between the different criteria in the case of the Ti-1Mn-1Fe and Ti-5Mn-5Fe alloys whist there is less certainty for the Ti-0.5Mn-0.5Fe and Ti-3.5Mn-3.5Fe alloys. For the former, this is due to the low amount of alloying elements, which leads to the stabilisation of a small amount of b phase that cannot be detected via XRD analysis. Therefore, this pushes the alloy to be classified as an a ? b Ti alloy, rather than a neara Ti alloy, if the MoE or Bo-Md criteria are used. For the Ti-3.5Mn-3.5Fe alloy, the relatively high amount of Mn and Fe, which are both strong b-stabilisers, still results in the formation of a lamellar microstructure combined with the formation of the metastable a 00 phase typical of metastable b Ti alloys.
In terms of mechanical behaviour (Fig. 5), the sintered ternary Ti-Mn-Fe alloys are characterised by an elastoplastic response when subjected to a quasi-static uniaxial load, with the exception of the Ti-5Mn-5Fe alloy, which shows a purely elastic response before failure. Therefore, the progressive simultaneous addition of Mn and Fe to Ti leads to a continuous increases in the yield and ultimate tensile strength, reaching the highest strength for the Ti-3.5Mn-3.5Fe alloy, whilst the ability to withstand plastic deformation decreases with the amount of alloying elements added (Fig. 5b). The mechanical behaviour of powder metallurgy materials is dependent on the residual porosity present in the material after sintering, as the residual pores act as stress concentrators [37]. At sufficiently high relative density levels (C 92-94%), a higher impact is expected in the case of the strain rather than for the strength. Specifically, as for the hardness (Fig. 7a), in the case of the sintered ternary Ti-Mn-Fe alloys the residual porosity is not the main factor determining their mechanical response as Fig. 7b) shows a very sharp increase in strength and a decrease in strain over a small relative density range. The enhancement of the strength, and the associated reduction in the strain of the sintered ternary Ti-Mn-Fe alloys, is therefore primarily controlled by the effects that the addition of strong bstabilising elements have on the phases and their features. Specifically, putting aside the Ti-5Mn-5Fe alloy, which has a microstructure composed of b grains, the reduction in the size of the a-grain boundary grains combined with the refinement of the   interlamellar spacing of the a ? b lamellae, the solid solution strengthening effect of the presence of a higher amount of atoms of the alloying elements dissolved within the Ti lattice, and the stabilisation of a greater amount of the b phase, are all responsible for the continuous increase of the resistance to plastic deformation, resulting in the increase of the strength and the decrease of the strain. This is evidently supported by the phases formed during the slow cooling from the sintering temperature (Fig. 3), where the formation of the metastable a 00 phase within a lamellar microstructure (i.e. Ti-3.5Mn-3.5Fe alloy) is still beneficial to enhance the strength of the material. However, its presence once the alloy has a b grains microstructure (i.e. Ti-5Mn-5Fe alloy) is detrimental and leads to the premature brittle failure of the alloy. Because of these effects, the failure mode of the sintered ternary Ti-Mn-Fe alloys switches from intergranular ductile to transgranular brittle (Fig. 6). The associated fracture surface, therefore, changes from being composed of dimples of the size of the a lamellae in the Ti-0.5Mn-0.5Fe and Ti-1Mn-1Fe alloys, to the presence of tear ridges in the Ti-3.5Mn-3.5Fe alloy, and the formation of river patterns within the b grains in the Ti-5Mn-5Fe alloy. Due to the limited amount of literature, it was not possible to compare the tensile behaviour of the sintered ternary Ti-Mn-Fe alloys with other Ti-Mn-Fe alloys; however, the strength of the sintered ternary Ti-Mn-Fe alloys is comparable, for an equivalent amount of alloying elements, to that of binary Ti-xFe alloys obtained via powder metallurgy [14] or casting [38], and that of binary Ti-Mn alloys obtained via powder metallurgy but with higher Mn concentrations [20]. The variation of the strain hardening rate as a function of the true plastic strain of the sintered ternary Ti-Mn-Fe alloys is shown in Fig. 9, where it can be seen that the shape of the strain hardening curves of the ternary Ti-Mn-Fe alloys with low addition rates of Mn and Fe (i.e. Ti-0.5Mn-0.5Fe and Ti-1Mn-1Fe) is similar, whereas is different for higher addition rates. Moreover, it can be noticed that it was not possible to obtain the strain hardening rate curve of the Ti-5Mn-5Fe alloy due to its brittle nature. A similar shape of the strain hardening rate curves means that the deformation mechanism is the same, which agrees with the fact that microstructure changes from purely lamellar for the Ti-0.5Mn-0.5Fe and Ti-1Mn-1Fe alloys (Fig. 4) to metastable for the Ti-3.5Mn-3.5Fe alloy. More in detail, in the initial stage of the curve, known as Stage II hardening [39] (i.e. below approximately 0.01 true plastic strain), the Ti-0.5Mn-0.5Fe alloy shows lower strain hardening rate than the Ti-1Mn-1Fe alloy up to approximately 0.25 true plastic strain, but subsequently the trend is reversed. This means that during the Stage III hardening [39], where dislocations generation and annihilation occur, the hindering of the movement of the dislocations becomes progressively more significant in the Ti-1Mn-1Fe alloy than in the Ti-0.5Mn-0.5Fe alloy. The former is, therefore, characterised by a lower strain hardening rate at higher true plastic deformation.

Conclusions
From this study about the development of ternary Ti-Mn-Fe alloys via the powder metallurgy blended elemental approach it can be concluded that the manufacturing route used to obtain the alloys is suitable for their development. The addition of Mn and Fe powders to Ti marginally affects the compressibility of the powder blends and leads to a continuously increasing trend of the relative density with the amount of alloying elements. In particular, both Mn and Fe promote the densification of the alloy due to their high diffusivity in Ti, therefore, reducing the amount of residual porosity. The progressive simultaneous addition of Mn and Fe changes the microstructure of the alloy from purely lamellar, typical of a ? b Ti alloys, to metastable b, reduces the size of the prior b grains, refines the interlamellar spacing between the a ? b lamellae and, eventually, leads to the formation of the metastable a 00 phase. The consequence of these microstructural changes is the

PlasƟc strain
Ti-3.5Mn-3.5Fe Ti-1Mn-1Fe Ti-0.5Mn-0.5Fe Figure 9 Variation of the strain hardening rate as a function of the true plastic strain of the sintered ternary Ti-Mn-Fe alloys.
monotonic strengthening and embrittlement of the material, up to the point that the Ti-5Mn-5Fe alloy is characterised by a purely elastic tensile behaviour. The same microstructural changes also determine the shape of the strain hardening rate curve of the alloy and, therefore, the actual strain hardening mechanism derived by the balance of dislocations generational and annihilation. As the amount of alloying elements in the ternary Ti-Mn-Fe alloys increases, the failure mode switches from intergranular to transgranular, reflecting the progressive more brittle nature of the ternary Ti-Mn-Fe alloys.

Acknowledgements
This work was supported by the New Zealand Ministry of Business, Innovation and Employment (MBIE) through the UOWX1402 research contract.

Funding
Open Access funding enabled and organized by CAUL and its Member Institutions.

Data and code availability
All metadata pertaining to this work will be made available on reasonable requests.

Declarations
Conflict of interest The authors declare no conflict of interest.
Ethical approval The work does not require ethical approval as no experiments involving human tissue were performed.
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