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Titanium and its alloys are widely used for biomedical applications such as bone repair or reconstruction due to their excellent biocompatibility, high specific strength and high corrosion resistance . However, a number concern associated with the titanium and it alloy used as implant materials are still arise. Discrepancy of mechanical properties between natural bones and the implant titanium alloy are raised as the first concern. The Young's modulus of implanted materials, such as Ti-6Al-4V (110 GPa) and CP Ti (105 GPa) is much higher than that of compact bone (10 - 40 GPa) . This mismatch can lead to stress shielding effect. Alloying elements of the titanium alloys are the second concern. Ni and Cr are well known as metal allergen . Vanadium has been confirmed causing cytotoxic and unpleasant tissue reactions while aluminium is associated with potential neurological disorders . Hence, aluminium and vanadium free titanium alloys as implant materials are being explored. The use of ¢-stabilizing elements, such as CP-Ti, Nb, Ta and Zr, has been reported to result in alloys with good compatibility concerning their interaction with cells . Furthermore, these alloying elements would also be expected to enhance strength and decrease the elastic modulus of the resulting alloys . According to this, titanium alloys containing non-toxic elements (e.g., Nb, Ta, Zr, Mo) have been proposed as implant materials.
Another concern is that titanium alloys are categorized as bioinert materials. The attachment of bioinert materials into a bone may lead to encapsulation by dense fibrous tissues resulting in weak bone implant interface, loosening the implant and fixation failure . Efforts to enhance bioactivity of implant materials can be carried out through coating onto substrates such as titanium powder coating using PM techniques [8, 9], hydroxyapatite (HA) coating and titanium coating using deposition techniques  or plasma sprayed techniques . The last techniques, currently is widely used in manufacturing technology for ceramic coating on metallic substrate. However previous research has found that this high temperature and rapid cooling process convey some concerns on microstructural and structural transformation due to rapid solidification of feedstock particle and partial melting. The inherent roughness of the coatings provided by these techniques, or bioactivity of coating materials, demonstrated that these combinations in fact support the osseointegration, with good adhesion to the substrate. Yet the interface between the coating layer and substrate becomes a concern , while the elastic modulus of the implant may remain constant. However, this may be overcome through introducing of porosity which is known to reduce the stiffness. In addition porous surface may entrap specific protein such as vitronectin and fibronectin to promote osteoblastic cell attachment, proliferation and differentiation on implant surface . Further bone ingrowths into the pore and bonding with adjacent tissue occur providing mechanical interlocking between the bone and implant materials. Therefore, by adjusting into appropriate porosity of the implant materials, fixation of the hard tissue replacement may be improved coinciding with reducing stress shielding effects and eventually prolonging the devices life-times.
Currently, variety solid state processing route to produce porous titanium for use in orthopaedic application have been developed. Sintering lose powder with low pressure compaction produces low porosities . In space holder techniques using spacer particle such as ammonium hydrogen carbonate  or polymeric materials  or magnesium , pure titanium foam with high porosities has been achieved (70-80%). Yet, contamination of the spacer during removal processing is still a significant concern. Furthermore, inert gas such as argon can be used to develop porosity through expansion pressurized argon bubble in metal matrix at elevated temperature. Using this technique porosity can be controlled by varying the parameter such as argon back filled pressure, foaming temperature and foaming holding time [18, 19]]. Porous pure titanium and titanium alloy based on prealloyed starting powder [17, 20] have been developed using this technique. However, the prealloyed powder often contains lower level of biocompatible elements and relatively more expensive compared to the elementals powders.
In the present work, the pressurized bubble entrapment was used to develop porous titanium alloy base on non toxic elemental starting powder (Ti-Ta Nb Zr). The fabrication proses, microstructural characterisation, porosity measurement, alloying phase transformation and the potential of application of these metal foam are described.
2. Experimental Method
The starting elementals powder i.e Ti, Nb, Ta and Zr supplied by CERAC Inc. (USA). They have a wide particle size distribution below 44 um and angular and irregular particle shape an example being shown by scanning electron microscopy in Figure 1. Alloy powder was prepared by mixing the weighed powders to produce the composition as listed in table 1 for 30 minutes in a roller mixer. Following mixing, approximately 50 gram of alloy powder was filled up into stainless can with inner diameter and length of 28 mm and 30 mm respectively. The cans were subsequently evacuated to approximately 10-2 Pa and backfilled with 0,48 MPa argon and sealed. The canned powder and Argon were densified by hot isostatic press (HIP) at 1000oC for 2 hours or 1100oC for 4 hours at ANSTO (Sydney, Australia) and then furnace cooled. Cubic specimens with approximately of length 10 mm were cut from centre of each HIP-ed billet using electro discharge machining (EDM). The EDM layer damage was removed by light polishing with 600 grit silicon carbide paper and cleaning ultrasonic cleaning using acetone. After densification the billet the specimens were then heated rapidly to isothermal temperature of 1350oC and 1150oC for 10 hours (isothermal foaming process) in graphite element vacuum furnace. Following this foaming process, selected specimens were metallographically prepared by grinding and polishing to 0.05 mm. Some specimens were etched using modified Krolll's solution (2.5% HF, 5 % HNO3 and 92.5 % H2O) and examined by optical microscope (OM) and scanning electron microscope (SEM). Their porosity distribution was evaluated by digital image analysis of several optical microscopy images for each specimens using freely available software (imageJ). Phase constituent within the specimens during those process were identified using X-ray diffraction (XRD) of flat specimens plate under condition of CuK¡, 40 kV and 30 mA.
3. Result and Discussion
Figure 2 shows typical X-ray diffraction patterns for samples before and after densification by HIP-ing. The blended powder patterns indicated peaks related to individual elements, Ti and Zr existed as hexagonal closed packed (hcp) structure. While the Nb and Ta were present as isomorphous body centered cubic (bcc structure). After the HIP-ing the samples revealed peaks attributed to ¡ and ¢ Ti, Nb and Ta, whereas peaks related to Zr, oxide or intermetalic were not identified. Peaks related to ¡ Ti were suppressed along with increasing HIP temperature and will vanish at sintering temperature higher than 1500oC.
The porous Ti-Nb-Ta-Zr alloy was developed and classified as ¡€«¢. The microstructure analysis demonstrate s that the two phase Widmanstattën structure develops with dissolution of Nb particle, that act as ¢-phase nucleator agent  by higher HIPing and foaming temperature. The microstructure of the HIP-ed alloy (Figure 3.a) at 1000oC revealed the formation the microstructure consists of angular titanium particles, niobium and tantalum (the brightest one) which is their initial morphology still visible individually. HIP-ing at high temperature above ¢ transition temperature (1000oC) and subsequent slow rate cooling to room temperature under vacuum resulted in of the beginning of the two phase Widmanstattën structure. Pores can be found due to uncompleted dissolution among the elemental powders and resembled the original inter powder spacing. In HIP-ing at 1100oC, the dissolution of the Nb and Ta particles within the HIP-ed specimens had continued and still distinguishable. The boundaries between the Ti, Nb and Ta particles were noted to have changed from the former angular in the starting powder to diffuse following the HIP procedure (figure 3.b). The dissolution of an ¡€«¢ phase with the Widmanstattën structure that was identified from (mainly) Nb and Ti particles with Nb being the phase nucleator agent. It can be seen that two phase Widmanstattën area become visible on titanium particles on area where the dissolution of ¢-stabilizer elements take places more slowly. The HIP-ed specimens at 1100oC and 1000oC were found to contain < 1 vol% and <3 vol % of porosity respectively with pore sizes of typically 10 m. Previous research has suggested that HIP-ed CP-Ti typically contain initial porosity levels ranging from around 0.14 - 2.4 vol% .
The HIP-ed billets were then exposed into high temperature (1160oC or 1350oC) in a vacuum furnace for 10 hours to undergo isothermal foaming process. The pressuring pores are expected to expand due to creep of the surrounding metal. During initial 10 hours of this foaming process, pore size appears to grow rapidly  with pore growth virtually being stopped following approximately 3 days of foaming .
Pore morphologies of the foamed specimen at 1160oC and 1350oC with backfilled argon pressure of 0.48 MPa considerably distinguishable. The first shows pore morphology more uniform with porosity level of approximation 30% due to their initial pores size and distribution which is typically uniform through the specimen (figure 4.a) ). On the other hand, pore distribution for specimen isothermal foamed at 1350oC showed not uniform pore size distribution (figure 4.b). Most pores resulted in primary equiaxed and more rounded with porosity level at center of the specimens of approximately 14% with pore size typically 20-50 m. The pores remain discrete with few pores coalescence (figure 4.c). In contrast, on the near surfaces of the specimen the pores expanded noticeably larger with the porosity level achieving approximately of 45 vol% and pore size of typically 100-200 m with (Figure 4.d). This was occurred due to fewer constraints being present at the surface region for the expansion process. Some interconnected pores has been found indicating that many pores had coalesced due to gas expansion sometimes collapsing multi interpore walls. This expansion will be inhibited if the driving force for further expansion reduced because of an increase in pore volume or escaping the pressurized gas to the atmosphere when the pore connect to the specimen surface.
Pore morphology involving pore size and interconnectivity are known to be important factor for the protein absorptions which initiate bone ingrowth into porous implant. Previous research noted that bone ingrowth into porous coating with pore size down to 50 m , while for generating mineralized bone, the interconnections of the porosity should be larger than 100 m . Hence the porous titanium alloy prepared in this work may have a pore size diameter appropriate for biomedical implant applications.
The homogenization of the HIP-ed alloy isothermal foamed at 1150oC is still incomplete indicated that the dissolution of Nb and Ta had continued. The Ta particles apparently had dissolved at foaming temperature of 1350oC resulting in more homogenous microstructure consisting of ¡ plates and hcp-martensite, ¡', dispersed in ¢ phase (figure 5). This existence of the martensite phase suggested that the alloy was likely to be an ¡€«¢ structure.
The current study has successfully demonstrated the potential fabrication of porous titanium alloys using a powder metallurgy with non-toxic elemental powder i.e. Ti, Nb, Ta, Zr as the starting material through pressurized gas entrapment technique. The alloys present porosity and pore size ranging from approximately 20-40 vol% and 20-200 m, respectively with some interconnections pores formed. The porous titanium alloy under investigation is expected to show great potential as a biocompatible implant material due to its porous structure allowing the ingrowth of new bone tissue.
This work is financially support by TPSDP, SPMU: Muhammadiyah University of Yogyakarta. The authors also wish to thank the Australian Nuclear Science and Technology Organisation (ANSTO) for carrying out the HIPing and foaming procedures.