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Effect of ball-milling time on the structural characteristics of biomedical porous Ti–Sn–Nb alloy
Materials Science and Engineering C 31 (2011) 921–928
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Materials Science and Engineering C j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / m s e c
Effect of ball-milling time on the structural characteristics of biomedical porous Ti–Sn–Nb alloy Alireza Nouri a,b,⁎, Peter D. Hodgson b, Cuie Wen c a b c
CQM-Centro de Química da Madeira, MMRG, Universidade da Madeira, Campus Universitário da Penteada, 9000-390 Funchal, Portugal Institute for Technology Research and Innovation, Deakin University, Geelong, Victoria 3217, Australia IRIS, Faculty of Engineering and Industrial Sciences, Swinburne University of Technology, 543-454 Burwood Road, Hawthorn, Victoria 3122 Australia
a r t i c l e
i n f o
Article history: Received 16 February 2010 Received in revised form 2 February 2011 Accepted 17 February 2011 Available online 24 February 2011 Keywords: Mechanical alloying Titanium alloys Porous materials Powder metallurgy Biomaterials
a b s t r a c t The structural characteristics of biomedical porous materials are crucial for bone tissue to grow into a porous structure and can also influence the fixation and remodeling between the implant and the human tissues. The current study has been investigating the effect of the ball-milling variable of time on the structural characteristics and pore morphology of a biomedical porous Ti–16Sn–4Nb (wt.%) alloy. The alloy was synthesized using high-energy ball milling for different periods of time, and the porous Ti–16Sn–4Nb alloy was fabricated by using a space holder sintering process. The resultant powder particles, bulk, and porous samples were characterized using a scanning electron microscope (SEM), laser particle-size analyzer, chemical analysis, X-ray diffraction analysis (XRD), and the Vickers hardness test. The results indicated that the inner pore surface, pore wall architecture, degree of porosity, pore size and the inter-pore connectivity of the sintered porous alloy are all considerably affected by ball-milling time. Crown Copyright © 2011 Published by Elsevier B.V. All rights reserved.
1. Introduction Bone is one of the most intensively studied tissues, due to the defects frequently caused by trauma or disease. The development of artificial organs and implants to replace injured and diseased hard tissues such as bones, teeth and joints is highly desirable in orthopedic surgery. When designing and fabricating an orthopedic biomaterial, the main goal is to restore the function and mobility of the native tissue requiring replacement. An ideal orthopedic implant must not only meet the biological requirements, but must also have adequate mechanical strength. Specifically, the implant aims to mimic the microand macro-porous architecture of the inorganic matrix of natural bone [1–3]. Since bone is a porous tissue material, there is a physiological rationale for the use of porous materials in its replacement. Bone surrounding the implant with high stiffness would be stress-shielded, and might therefore result in detrimental bone resorption [4,5]. Introducing pores into a material structure reduces the stiffness values close to that of natural bone and can provide a good load transfer and stimulate the formation of new bones [6]. The porosity also allows the bone tissue to grow into the porous structure, providing an adequate biological fixation and an efficient nutritional supply and vascularization. ⁎ Corresponding author at: CQM-Centro de Química da Madeira, MMRG, Universidade da Madeira, Campus Universitário da Penteada, 9000-390 Funchal, Portugal. Tel.: + 351 966 131 895; fax: + 351 291 705 149. E-mail address: [email protected] (A. Nouri).
Compared to other biomaterials such as ceramics and polymers, metallic biomaterials offer a wider range of the mechanical properties required for most load bearing applications in fracture fixation and bone replacement (total joint arthroplasty) [7,8]. Amongst them, Ti and most of its alloys are extensively used in dental and orthopedic applications due to their appropriate biological and mechanical properties [9–11]. The biomedical applications of Ti alloys with alloying elements of Sn and Nb have previously been studied by Okazaki et al. [12,13], Kawahara [14], Ito et al. [15], and Niinomi [16]. According to these studies, Sn and Nb are non-toxic elements and have not been associated with any adverse tissue reaction. In addition, Sn is a solid solution strengthener for Ti alloys [17,18], whereas Nb, as a β-stabilizing element, contributes to a decrease in the bulk elastic modulus of the alloy which is required for orthopedic applications [12]. Over the past few years, mechanical alloying (MA) has shown great potential in synthesizing a wide variety of nanocrystalline, supersaturated solid solutions and amorphous phase with unique characteristics [19–21]. The MA is a solid-state powder metallurgical process and is capable of processing Ti alloys with homogeneous microstructures and improved mechanical properties than the conventional powder metallurgy or casting techniques [22–24]. In the present study, a biomedical porous Ti–16Sn–4Nb (wt.%) alloy was fabricated from elemental powders via the MA process for different ball-milling times. One major aspect in the development of suitable porous metallic biomaterials is to obtain the appropriate structural characteristics in order to improve their biological performance. These structural characteristics such as pore wall architecture, degree of porosity, pore size
0928-4931/$ – see front matter. Crown Copyright © 2011 Published by Elsevier B.V. All rights reserved. doi:10.1016/j.msec.2011.02.011
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Table 1 Characteristics of the elemental metal powders in the current study. Elements
Purity (%)
Size (μm)
Morphology
Ti Sn Nb
99.9 99.9 99.8
≤ 45 (− 325 mesh) ≤ 45 ≤ 45
Irregular Spherical Irregular
and inter-pore connectivity, have all been shown to influence bone ingrowth into porous implants and the mechanical properties of the implants [25–28]. The present work aimed to investigate the effect of the ball-milling process underlying structure–property relationships in the biomedical porous Ti–16Sn–4Nb alloy. 2. Materials and methods Elemental metal powders of Ti, Sn and Nb were mixed together according to the desired composition of Ti–16Sn–4Nb alloy. The characteristics of the elemental metal powders are shown in Table 1. The MA process was conducted in a planetary ball mill (Retsch, PM400) using steel containers and balls at room temperature. The weight ratio of ball to powder was maintained at 10:1 and the ballmilling was carried out at a rotation rate of 200 rpm. The powders were handled in a glove box chamber under argon gas to prevent any atmospheric contamination. The containers were also sealed and purged with high purity argon prior to the ball-milling process. Furthermore, the ball-milling was carried out for various times, namely 1 h, 10 h and 20 h without using any process control agent (PCA). For fabrication of the porous Ti–16Sn–4Nb alloy, the ball-milled powders were thoroughly mixed with 50 wt.% ammonium hydrogen carbonate, (NH4)HCO3, as a space-holding material with a particle size of 300–500 μm. A small amount of ethanol was added during the mixing of the powders and ammonium hydrogen carbonate to prevent segregation. However, for the fabrication of bulk samples no space-holding material was utilized. The powders were subsequently compacted by a uniaxial cold press under the pressure of 500 MPa. The first step used in fabricating the porous samples was the heat treatment of the green compact at 170 °C for 3 h. The samples were then subjected to 1200 °C sintering for 5 h in a high vacuum furnace
with a heating/cooling rate of 10 °C/min. However, for fabrication of the bulk samples, the green compacts were directly heated to 1200 °C with the aforementioned sintering time and heating/cooling rate. It should be noted that one series of bulk and porous samples was also synthesized from the blended powder mixture without using the ballmilling process (Fig. 1a). A quantitative chemical analysis of the bulk samples after sintering was carried out using the inductively coupled plasma-atomic emission spectrometry (ICP-AES) and Leco combustion methods. The morphology and microstructure of the powder particles as well as the structure of the porous samples were characterized by means of a scanning electron microscope (Leica S440) combined with secondary electron/backscattered electron imaging (SEI/BEI). Particle-size distribution was measured using Malvern Instruments Mastersizer 2000 with a Hydro 2000S side feeder. The porosity of each porous sample was calculated by measuring their weight and dimensions. Accordingly, 6 samples from each group were weighed by a precision digital balance (Sartorius BP 221S) and their dimensions were measured using a digital caliper. The height and diameter of cylindrical samples were measured at three different positions in order to minimize the measurement error. Phase formation was characterized using XRD in the powder and bulk samples by Cu Kα radiation (35 kV, 28 mA) at a scanning rate of 1°/min using a Philips PW 1820 diffractometer. The Vickers hardness of the sintered bulk samples was measured using a 9.8 N (1 kgf) load for 15 s. Average hardness values were obtained from six indents on each of the bulk samples. 3. Results and discussion 3.1. Morphological and microstructural analyses of Ti–16Sn–4Nb alloy particles The morphology of the blended Ti–16Sn–4Nb powders is shown in Fig. 1a. The powder particles did not undergo any noticeable change after 1 h of ball-milling and the majority of the particles retained their original shape at this stage, as shown in Fig. 1b. The particles after the short ball-milling time of 1 h had not collided sufficiently with the balls, although they became slightly work hardened and agglomerated. The small difference in work hardening rate between
Fig. 1. SEM micrographs of (a) blended powders and the Ti–16Sn–4Nb powders ball-milled for: (b) 1 h; (c) 10 h; and (d) 20 h.
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the samples made from the powders ball-milled for 1 h and those of made from the blended elemental powders can be explained by the lower degree of compaction (larger degree of porosity) and slightly higher hardness of the former. The effect of work hardening on powder compaction will be discussed further in Section 3.5. After ballmilling for 10 h the majority of the particles became flattened and showed a relatively smaller size compared to that of the particles ballmilled for 1 h (Fig. 1c). Nevertheless, ball-milling up to 10 h did not result in any significant particle refinement. The presence of flattened and thin flaky particles after 10 h ball-milling is indicative of the micro-forging and cold welding of the particles at this stage. As seen in Fig. 1d, the morphology and size of powders drastically changed after 20 h ball-milling, and the powders exhibited a rather equiaxed shape and a considerable smaller size than the powders ball-milled for 1 and 10 h. It is notable that the absence of PCA in the ball-milling process resulted in an intense cold welding of powders to the milling tools for the long ball-milling time of 20 h and, thus, a considerably small amount of powder was recovered after 20 h ball-milling. As a result of the work hardening effect and high density of defects such as dislocations and vacancies for long ball-milling times, the fracture and welding of powders reach an equilibrium state. This state leads to the formation of rather equiaxed and small particles. However, due to the higher bonding strength of the finer particle size, the ability of particles to stand further plastic deformation is decreased and a higher force is required to fracture the small particles during the ballmilling process [20]. It is worth mentioning that the reduction of particle size at a given milling time can also be influenced by milling technique. For instance, high-energy ring mills are more efficient to reduce the particle size than ball-milling process. In the ring mill technique, particles are under the simultaneous effect of compression and shear forces [29–31]. Fig. 2 shows the particle-size distribution of Ti–16Sn–4Nb alloy after ball-milling for 1, 10, and 20 h. Besides a slight decrease in particle size, there is no considerable difference in powder size distribution up to 10 h ball-milling. At this stage, the colliding balls are mostly used to flatten the particles through shearing and the compressive force. The particle-size distribution of the powder ball-milled for 20 h was wider and showed a considerable shift to the left-hand end of the particle-size scale, indicating a decrease in particle size after ball-milling for 20 h. The broadening of particle-size distribution at longer ball-milling times is a typical behavior of the high-energy ball milling process [19,32,33]. There is also an overall tendency to weld small particles into larger pieces, as evidenced by a small hump at the right-hand end of the particle-size scale. The agglomeration of 20 h ball-milled powders is mostly due to the increasing resistance to fracture and the strong tendency of cohesion between particles with
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decreasing particle size. In the present study, ball-milling for 1 and 10 h resulted in a more symmetrical particle-size distribution. Fig. 3 shows the SEM backscattered electron images of the crosssection of the Ti–16Sn–4Nb alloy particles after ball-milling for 1, 10, and 20 h. Upon ball-milling for 1 h, the majority of particles remained unalloyed. The particles are distinguished by dark contrast areas for Ti and bright contrast areas for the Sn or Nb (Fig. 3a). At this stage, the ball-milling time is too short to achieve alloy formation. Principally, for a short ball-milling time, the particles are not subjected to an adequate compressive force or kinetic energy input, and therefore cannot be mechanically alloyed [20,34]. With continued ball-milling to 10 h, only small fragments of Sn and Nb were trapped into Ti powders and the particles were only mechanically bonded together, as shown in Fig. 3b. Increasing the ball-milling time to 20 h resulted in significant changes in the microstructure of the particles. As seen in Fig. 3c, the solute elements of Sn and Nb were embedded in the interfacial boundaries and were incorporated into the Ti matrix. As such, each particle displayed a lamellar-like structure consisting of a fine and relatively homogeneous distribution of the solute elements within the Ti matrix. In general, a more homogeneous distribution is expected when the ball-milling time is increased. Homogeneity of the particles is the result of a balance between fracturing and cold welding during the MA process [20]. The heavy plastic deformation of particles at longer ball-milling times gives rise to a large number of defects, providing a higher diffusivity of the constituent elements [34,35]. In other words,
8
Volume (%)
MA-1 h
MA-10 h
6
MA-20 h 4
2
0 0.01
0.1
1
10
100
1000
10000
Particle size ( m) Fig. 2. Particle size distribution of Ti–16Sn–4Nb alloy at different ball-milling times.
Fig. 3. The SEM backscattered electron images of the cross-section of the Ti–16Sn–4Nb particles for different ball-milling times: (a) 1 h; (b) 10 h; and (c) 20 h.
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continuous fracturing and welding of particles at long ball-milling times reduce the diffusion distance of alloying elements to micrometer range, leading to the lamellar microstructure, as seen in Fig. 3c. In contrast to a thermally induced diffusion process, the diffusion during the MA process is mostly controlled by mechanical energies [35]. The amount of powder recovered from the container at different ball-milling times is plotted in Fig. 4. It is clear that an increase in the ball-milling time led to a decrease in the powder yield. Almost the whole amount of powder was recovered after ball-milling for 1 h. By increasing the ball-milling time to 10 h, about 85% of the powder was recovered from the containers. This value is considered to be a high powder yield since no PCA was used in the present study. When the ball-milling time increased to 20 h, only a low powder yield of about 30% was obtained. At this point, a large amount of powder remained in the grinding media as evidenced by a relatively thick layer of powder over the balls and the inner wall of the container. Principally, due to the frequent mechanical impact of the balls on the powder particles, part of the milled powders are cold-welded over the grinding media, forming a coating layer [36,37]. However, the thickness of the coating layer is mainly dependent on the ball-milling time, the ductility of the initial powders, the amount of PCA and the intensity of milling (e.g. ball to powder weight ratio and milling speed). 3.2. XRD analysis on powder and bulk Ti–16Sn–4Nb alloy samples The XRD patterns of the Ti–16Sn–4Nb alloy powder mixtures and bulk samples sintered from the powders ball-milled for different times are shown in Fig. 5 and Fig. 6, respectively. The initial powder mixture consists of the elemental metals of Ti, Sn, and Nb (Fig. 5a). After 1 and 10 h ball-milling, the intensity of peaks corresponding to Sn and Nb did not show any significant change (Fig. 5b and c). However, on continuous ball-milling for 20 h, the sharp peaks from the crystalline elements of Sn and Nb disappeared, as seen in Fig. 5d. Peak broadening and the reduction in intensity of the diffraction peaks of ball-milled powders are associated with the refinement in crystallite size, lattice internal strain, and the instrumental effects [19]. With 10 h ball-milling or less, little solid solution took place between the powder particles, as evidenced by the presence of the diffraction peaks of Sn and Nb. The presence of Sn and Nb diffraction peaks at this stage is in agreement with the inhomogeneity of powder microstructure, as shown in Fig. 3b. Increasing the ball-milling time to 20 h led to the disappearance of the Sn and Nb diffraction due to the grain size reduction and/or accumulation of mechanical strains (Fig. 5d). The formation of an hcp α-Ti phase after 20 h ball-milling can be attributed to the dissolution of the alloying elements in the Ti lattice. The aforementioned result is consistent with the formation of the lamellar-like microstructure
Ti-16Sn-4Nb (wt.%) −Ti Sn Nb
Intensity (a.u.)
924
d 30
70 60 50 40
40
50 2 (deg.)
60
70
Fig. 5. XRD patterns of (a) elemental powders and the Ti–16Sn–4Nb powders ballmilled for: (b) 1 h; (c) 10 h; and (d) 20 h.
seen in Fig. 3c. However, the disappearance of the Sn and Nb peaks can be mainly associated with the formation of a titanium solid solution. With sintering of the ball-milled powders at a high temperature, no diffraction peaks of Sn and Nb were detected and only the Ti diffraction peaks were identified (Fig. 6). The intensity of Ti diffraction peaks in the sintered bulk samples decreased with the increase of ball-milling time. The diffraction patterns of the sintered bulk samples made from the powders ball-milled for 1 and 10 h also indicated a small fraction of β-Ti phase, as seen in Fig. 6b and c. This peak represents the small amount of Nb used as a β stabilizer in the synthesis of the present Ti alloy. However, after ball-milling for 20 h, the β phase became almost indistinguishable and merged into a single α-Ti phase, suggesting the complete solid solution of Nb in α-Ti (Fig. 6d). 3.3. Chemical analysis of bulk Ti–16Sn–4Nb alloy samples The content of contamination in the bulk Ti–16Sn–4Nb alloy samples made from the powders ball-milled for various ball-milling
Ti-16Sn-4Nb (wt.%) Ti
Intensity (a.u.)
Weight Recovery (%)
80
b c
100 90
a
Ti Sn Nb
a b c
30
d
20 10 0 0
5
10
15
20
25
Milling Time (h) Fig. 4. The amount of powder recovered from the ball-milling container as a function of ball-milling time.
30
40
50
60
70
(deg.) Fig. 6. XRD patterns of (a) elemental powders and the bulk Ti–16Sn–4Nb made from the powders ball-milled for: (b) 1 h; (c) 10 h; and (d) 20 h.
A. Nouri et al. / Materials Science and Engineering C 31 (2011) 921–928
3.4. Hardness of the bulk Ti–16Sn–4Nb alloy sample The hardness values were measured on the bulk Ti–16Sn–4Nb samples made from a blended powder mixture as well as those of samples made from the powders ball-milled for different times. The variation in hardness versus different ball-milling times is plotted in Fig. 7. It can be seen that the hardness increased with increasing ballmilling time. The lowest hardness belonged to the sample made from the blended powder mixture showing the value of about 300 HV1.0. However, the hardness did not show any considerable change after ball-milling for 1 h. For the samples made from the powders ballmilled for 10 h, the Vickers hardness increased to approximately 370 HV1.0. Subsequently, a 40% increase in hardness was achieved when ball-milling time was increased from 10 h to 20 h, reaching a maximum of approximately 520 HV1.0. A progressive increase in hardness of the bulk Ti–16Sn–4Nb alloy made from the 20 h ball-milled powders could be put down to: (i) the work hardening of the powders; (ii) crystallite size refinement; (iii) dispersoid strengthening due to the increased contamination at longer ball-milling times; and (iv) solid solution strengthening. All the aforementioned factors impede dislocation movement and lead to a strengthening of the material. However, compared to
500
Hardness (HV1.0)
times is given in Table 2. The results indicate that the increase of ballmilling time leads to a noticeable contamination of oxygen in the bulk samples. The contaminations from oxygen have frequently been reported during the milling of Ti alloys [38–41], which can be traced to the milling atmosphere, chemical purity of the starting powders, and the decomposition of the PCA used during milling. However, in the present study the resultant contamination from oxygen is mainly attributed to the leakage of air into an inadequate sealed container which interacts with the large and fresh surface areas of powder particles during high-energy ball milling. It is also notable that the longer ball-milling times of 10 and 20 h also resulted in increased contamination of Fe and Cr. The Fe and Cr contaminations are from the wear of the stainless steel milling balls and containers during the ballmilling process. These metallic contaminations are more pronounced for the samples made from the powders ball-milled for 20 h. Higher levels of C, Fe and Cr contamination can be expected if the powders are ball-milled with a PCA addition, as previously reported by the same authors [38,42]. It is possible that the coating film formed on the milling tools during the longer ball-milling times prevents excessive wear of the milling media. This is indicated by the smaller difference in the amount of Fe and Cr contamination between 20 and 10 h ball-milled powders as compared to the difference in metallic contamination between 10 and 1 h ball-milled powders. However, the thickness of the film should be kept to a minimum to increase the powder yield and avoid forming a heterogeneous product [19,43]. The increase in content of contamination with ball-milling time is a common problem with the MA technique. In fact, an argon atmosphere is incapable of reducing the oxygen partial pressure sufficiently and it is practically impossible to prevent the surrounding atmosphere from leaking into the milling container during the MA process [44,45].
925
400 300 200 100 0 0
5
10
15
20
Milling Time (h) Fig. 7. The Vickers hardness of the bulk Ti–16Sn–4Nb alloy made from the powders ballmilled for different times.
the hardness values obtained from the bulk samples with different amounts of PCA [38], the dispersoid strengthening may not be an influential factor in the high hardness of the resultant samples at longer ball-milling times. Finally, most pure metals can be strengthened by the alloying elements. This is due to the fact that the presence of alloying atoms along the core of dislocation reduces the strain energy of the crystal, suggesting more stress to move the dislocation away from each other [17]. In particular, it has been reported that Sn has a strengthening effect in Ti lattice [12,46].
3.5. Structural characteristics of the porous Ti–16Sn–4Nb alloy samples 3.5.1. Degree of porosity The degree of porosity in the porous Ti–16Sn–4Nb samples as a function of ball-milling time is shown in Fig. 8. Although the weight ratios of the metal powders to the amount of space-holding material in this experiment were calculated to produce overall porosities of around 70%, the resultant samples made from the powders ball-milled for 20 h revealed a noticeable deviation from the calculated value. As seen in Fig. 8, the porosity of samples increases with an increase in ball-milling time, reaching a maximum value of around 74% after ballmilling for 20 h. This deviation can also be seen by the increase in error bars for longer ball-milling times. The error bar was the mean square deviation obtained from the measurements of six different porous samples for a given milling time. Generally, the porous powder metallurgical products synthesized from the particles with a broad particle-size distribution have a lower reproducibility and uniformity of the final porosity [47]. The minimum porosity was obtained for the porous samples synthesized from the blended powder mixture, showed a porosity value of 68%.
Table 2 The content of contamination in bulk Ti–16Sn–4Nb alloy made from the powders ballmilled for various periods and sintered at 1200 °C for 5 h. Ball-milling time
Contamination content (wt.%) O
N
C
Fe
Cr
1h 10 h 20 h
1.25 2.15 2.65
0.01 0.01 0.03
0.1 0.1 0.22
0.2 0.41 0.47
0.11 0.28 0.35
Fig. 8. The degree of porosity for the porous Ti–16Sn–4Nb samples versus different ballmilling times.
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3.5.2. Morphology of the porous alloy samples Fig. 9 shows the SEM micrographs of the porous Ti–16Sn–4Nb samples made from the powders ball-milled for different times at two different magnifications (250× and 1000×). It should be noted that the morphology of pores and pore walls in a porous material can be influenced by the morphology of its constituent powders [48,49]. The pore walls and inner pore surfaces in the porous samples made from the blended powder mixture or powder ball-milled for 1 h were smoother and only a few micro-pores or interstices could be seen, as shown in Fig. 9(a,b) and Fig. 9(c,d), respectively. In contrast, the rather round and equiaxed particles obtained from the powders ball-milled for 20 h, resulted in corrugated and rough pore walls and
inner pore surfaces (Fig. 9(g,h)). An increase in ball-milling time also resulted in a larger number of dents and concavities on the surface of the porous material. Moreover, in the porous sample made from the powders ball-milled for 20 h, a smaller inter-particle neck size and a larger number of micro-pores could be seen while the individual particles could usually be distinguished. However, the effect of powder morphology on the size of the pores was negligible, since the pore size is strongly controlled by the size of the space holder particles and the sintering temperature. Fig. 10 shows the backscattered electron images of the cell wall microstructure in the porous Ti–16Sn–4Nb samples for different ballmilling times. As shown in Fig. 10a, the porous sample made from the
Fig. 9. SEM micrographs of the porous Ti–16Sn–4Nb samples made from (a,b) a blended powder mixture and powders ball-milled for: (c,d) 1 h; (e,f) 10 h; and (g,h) 20 h.
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Fig. 10. The SEM backscattered electron images of the cell wall microstructure in the porous Ti–16Sn–4Nb samples made from the powders ball-milled for: (a) 1 h; (b) 10 h; and (c) 20 h.
powders ball-milled for 1 h exhibited a heterogeneous microstructure, as evidenced by the high volume fraction of the β Nb-rich phase (bright areas) in α-Ti matrix (dark areas). The composition of the aforementioned phases has already been reported in our previous papers using EDX analysis [42,50]. With increasing ball-milling time to 10 h, the volume fraction of the β phase slightly decreased and a more homogeneous microstructure was achieved (Fig. 10b). However, after 20 h ball-milling, the bright Nb-rich colonies were mostly homogenized in the dark Ti matrix showing a considerable decrease in size and volume fraction of the Nb-rich phase, as shown in Fig. 10c. The biomedical characteristics and mechanical properties of a porous metallic implant are highly affected by its total porosity, size and the morphology of the pores [51–53]. Apart from the type of processing method, the total porosity, density, and pore morphology in a powder metallurgical porous material are significantly controlled by its particle packing characteristics. However, the selection of the processing method is also based on the powder characteristics and the type of porosity required by the application. In general, the density of a green powder compact consisting of coarse powder (i.e. blended and 1 h ball-milled powders in the present study) is saturated under a low compaction pressure. This phenomenon is caused by weak interparticle attractive forces and low friction between the powder particles, due to the small number of contacts between the powder particles [54,55]. It can be assumed that the fine powders obtained
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with a ball-milling time of 20 h have stronger inter-particle attractive forces. These forces are a result of the electrostatic field, Van der Waals and moisture/surface adsorption forces that become much more significant with a decrease in particle size [54,56]. In addition, the powders ball-milled for long times cannot move and deform effectively during compaction. This may be primarily due to the effect of a decrease in particle size, solid solution strengthening and work hardening of the particles during the ball-milling process [57,58]. Amongst these factors, the work hardening effect appears to be the main determinant in governing the plastic behavior of powder particles during compaction. In the ball-milling process, the work hardening is caused by the interaction of dislocations with each other through the continuous mechanical impact of the balls on the powder particles. Parvin et al. [59], in a study on mechanically alloyed SiCreinforced Al composites, concluded that the effect of work hardening due to milling time is more pronounced than the morphological changes of the powder particles. Sadek and Salem [60], also suggested that the work hardening of the nanopowders after the ball-milling process resulted in the resistance of the particles to plastic deformation and hence retarded the compaction. Besides the shape, size and distribution of the particles, the particle packing characteristics of a green compact depend directly on the properties of the materials such as yield strength, impurities, and hardness [49,54]. These factors seem to be dominant for the total porosity of the samples after longer ball-milling times. Hewitt and Kibble [61] recently investigated the effects of ball-milling time of 0.5, 1, 3, and 5 h on the synthesis and consolidation of WC-10%Co powder at room temperature. They have shown that the relative density of compacted powders decreases linearly with milling time due to the combined effects of particle size, morphology and hardness. At 400 MPa compaction pressure, 64% relative density was achieved for the powders milled for 30 min compared to 61% after milling for 300 min. The small powder particles give rise to a lower packing density. In other words, a higher degree of porosity or lower relative density in compacts is achieved when the surface area of powders is increased. The reason for this is that small particles result in a larger number of contacts per particle and particle bridging, thereby leading to an increase in inter-particle friction and resistance to compaction [49,54]. Furthermore, the small particles provide less available sites for mechanical interlocking. In the present study, the absence of PCA during the MA process also intensified the interparticle friction between the resultant ball-milled powders since the rearrangement of the particles, which is the first densification step of cold die compaction, is extensively controlled by the lubricating effect of PCA. In another study by the same authors [38], a decidedly upward trend in the relative density of the sintered bulk samples was observed with an increase in PCA content. In the present study, the porous samples made from the powders ball-milled for 20 h exhibited more interconnected micro- and macropores. A large number of interconnected porosities enhance the circulation and exchange of body fluids, increase the rate and distribution of osteogenesis around and throughout the implant, promote vascularization and nutritional supply, and improve implant osseointegration [62]. Since the vascular network is strongly influenced by the continuity of pores, the degree of interconnectivity between pores is more crucial for new bone formation than pore size and porosity [63–65]. The interconnectivity of the pores can be theoretically increased by an increase in pore size and/or porosity, which in turn compromises the mechanical strength of the implant. Hence, an optimum balance between porosity and strength must be achieved, to ensure that the implant can at least withstand the forces applied during the surgical procedure and initially at the implantation site. In the current study, the porous samples synthesized with different ball-milled powders exhibited the macro-pore-structure of 100–400 μm which is desired for optimal bone cell ingrowth [66,67].
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The ball-milling process can be a promising method in synthesizing a desired porous structure for optimal bone cell ingrowth with increased overall porosity, as long as the mechanical strength is accommodated. The diverse porous structures obtained with different ball-milling times can provide new substrates for evaluating the mechanical strength and influence of pore architecture on in vitro cell proliferation.
[13] [14] [15] [16] [17]
4. Conclusion
[22] [23] [24] [25] [26]
The effect of different ball-milling times on the structural characteristics of powders, bulk and porous Ti–16Sn–4Nb alloy was investigated for various periods of time from 1 to 20 h. Ball-milling for 20 h led to the formation of a relatively equiaxed shape and smaller size powder particles with more asymmetrical particle-size distribution than the powders ball-milled for 1 and 10 h. The diffraction patterns of the sintered bulk samples made from the powders ballmilled for 1 and 10 h indicated a small fraction of β-Ti phase. However, further ball-milling to 20 h resulted in merging the β phase into a single α-Ti phase and, presumably, the formation of a partial amorphous phase. An increase in ball-milling time also revealed some increase in the amount of contamination introduced by the milling atmosphere and grinding media. The hardness of the sintered bulk samples showed a decidedly upward trend with an increase in ballmilling time reaching to the maximum value of ~ 520 HV1.0 for the sample made from 20 h ball-milled powder. The pore walls and inner pore surfaces in the porous samples made from the blended powder mixture and powder ball-milled for 1 and 10 h were smoother with a small number of micro-pores. However, the porous samples made from the 20 h ball-milled powders exhibited corrugated pore walls and rough inner pore surfaces with a larger number of micro-pores mostly due to the combined effects of particle size, morphology, and work hardening. The resultant porous samples made from the 20 h ball-milled powders showed a higher degree of porosity with more interconnected micro- and macro-pores. Acknowledgments The authors acknowledge the financial support from the VCAMM (Victorian Centre for Advanced Materials Manufacturing) and ARC (Australian Research Council) through the ARC Discovery Project DP0770021. Peter Hodgson is also supported by the ARC through a Federation Fellowship. References [1] [2] [3] [4] [5] [6] [7]
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Contents lists available at ScienceDirect
Materials Science and Engineering C j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / m s e c
Effect of ball-milling time on the structural characteristics of biomedical porous Ti–Sn–Nb alloy Alireza Nouri a,b,⁎, Peter D. Hodgson b, Cuie Wen c a b c
CQM-Centro de Química da Madeira, MMRG, Universidade da Madeira, Campus Universitário da Penteada, 9000-390 Funchal, Portugal Institute for Technology Research and Innovation, Deakin University, Geelong, Victoria 3217, Australia IRIS, Faculty of Engineering and Industrial Sciences, Swinburne University of Technology, 543-454 Burwood Road, Hawthorn, Victoria 3122 Australia
a r t i c l e
i n f o
Article history: Received 16 February 2010 Received in revised form 2 February 2011 Accepted 17 February 2011 Available online 24 February 2011 Keywords: Mechanical alloying Titanium alloys Porous materials Powder metallurgy Biomaterials
a b s t r a c t The structural characteristics of biomedical porous materials are crucial for bone tissue to grow into a porous structure and can also influence the fixation and remodeling between the implant and the human tissues. The current study has been investigating the effect of the ball-milling variable of time on the structural characteristics and pore morphology of a biomedical porous Ti–16Sn–4Nb (wt.%) alloy. The alloy was synthesized using high-energy ball milling for different periods of time, and the porous Ti–16Sn–4Nb alloy was fabricated by using a space holder sintering process. The resultant powder particles, bulk, and porous samples were characterized using a scanning electron microscope (SEM), laser particle-size analyzer, chemical analysis, X-ray diffraction analysis (XRD), and the Vickers hardness test. The results indicated that the inner pore surface, pore wall architecture, degree of porosity, pore size and the inter-pore connectivity of the sintered porous alloy are all considerably affected by ball-milling time. Crown Copyright © 2011 Published by Elsevier B.V. All rights reserved.
1. Introduction Bone is one of the most intensively studied tissues, due to the defects frequently caused by trauma or disease. The development of artificial organs and implants to replace injured and diseased hard tissues such as bones, teeth and joints is highly desirable in orthopedic surgery. When designing and fabricating an orthopedic biomaterial, the main goal is to restore the function and mobility of the native tissue requiring replacement. An ideal orthopedic implant must not only meet the biological requirements, but must also have adequate mechanical strength. Specifically, the implant aims to mimic the microand macro-porous architecture of the inorganic matrix of natural bone [1–3]. Since bone is a porous tissue material, there is a physiological rationale for the use of porous materials in its replacement. Bone surrounding the implant with high stiffness would be stress-shielded, and might therefore result in detrimental bone resorption [4,5]. Introducing pores into a material structure reduces the stiffness values close to that of natural bone and can provide a good load transfer and stimulate the formation of new bones [6]. The porosity also allows the bone tissue to grow into the porous structure, providing an adequate biological fixation and an efficient nutritional supply and vascularization. ⁎ Corresponding author at: CQM-Centro de Química da Madeira, MMRG, Universidade da Madeira, Campus Universitário da Penteada, 9000-390 Funchal, Portugal. Tel.: + 351 966 131 895; fax: + 351 291 705 149. E-mail address: [email protected] (A. Nouri).
Compared to other biomaterials such as ceramics and polymers, metallic biomaterials offer a wider range of the mechanical properties required for most load bearing applications in fracture fixation and bone replacement (total joint arthroplasty) [7,8]. Amongst them, Ti and most of its alloys are extensively used in dental and orthopedic applications due to their appropriate biological and mechanical properties [9–11]. The biomedical applications of Ti alloys with alloying elements of Sn and Nb have previously been studied by Okazaki et al. [12,13], Kawahara [14], Ito et al. [15], and Niinomi [16]. According to these studies, Sn and Nb are non-toxic elements and have not been associated with any adverse tissue reaction. In addition, Sn is a solid solution strengthener for Ti alloys [17,18], whereas Nb, as a β-stabilizing element, contributes to a decrease in the bulk elastic modulus of the alloy which is required for orthopedic applications [12]. Over the past few years, mechanical alloying (MA) has shown great potential in synthesizing a wide variety of nanocrystalline, supersaturated solid solutions and amorphous phase with unique characteristics [19–21]. The MA is a solid-state powder metallurgical process and is capable of processing Ti alloys with homogeneous microstructures and improved mechanical properties than the conventional powder metallurgy or casting techniques [22–24]. In the present study, a biomedical porous Ti–16Sn–4Nb (wt.%) alloy was fabricated from elemental powders via the MA process for different ball-milling times. One major aspect in the development of suitable porous metallic biomaterials is to obtain the appropriate structural characteristics in order to improve their biological performance. These structural characteristics such as pore wall architecture, degree of porosity, pore size
0928-4931/$ – see front matter. Crown Copyright © 2011 Published by Elsevier B.V. All rights reserved. doi:10.1016/j.msec.2011.02.011
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Table 1 Characteristics of the elemental metal powders in the current study. Elements
Purity (%)
Size (μm)
Morphology
Ti Sn Nb
99.9 99.9 99.8
≤ 45 (− 325 mesh) ≤ 45 ≤ 45
Irregular Spherical Irregular
and inter-pore connectivity, have all been shown to influence bone ingrowth into porous implants and the mechanical properties of the implants [25–28]. The present work aimed to investigate the effect of the ball-milling process underlying structure–property relationships in the biomedical porous Ti–16Sn–4Nb alloy. 2. Materials and methods Elemental metal powders of Ti, Sn and Nb were mixed together according to the desired composition of Ti–16Sn–4Nb alloy. The characteristics of the elemental metal powders are shown in Table 1. The MA process was conducted in a planetary ball mill (Retsch, PM400) using steel containers and balls at room temperature. The weight ratio of ball to powder was maintained at 10:1 and the ballmilling was carried out at a rotation rate of 200 rpm. The powders were handled in a glove box chamber under argon gas to prevent any atmospheric contamination. The containers were also sealed and purged with high purity argon prior to the ball-milling process. Furthermore, the ball-milling was carried out for various times, namely 1 h, 10 h and 20 h without using any process control agent (PCA). For fabrication of the porous Ti–16Sn–4Nb alloy, the ball-milled powders were thoroughly mixed with 50 wt.% ammonium hydrogen carbonate, (NH4)HCO3, as a space-holding material with a particle size of 300–500 μm. A small amount of ethanol was added during the mixing of the powders and ammonium hydrogen carbonate to prevent segregation. However, for the fabrication of bulk samples no space-holding material was utilized. The powders were subsequently compacted by a uniaxial cold press under the pressure of 500 MPa. The first step used in fabricating the porous samples was the heat treatment of the green compact at 170 °C for 3 h. The samples were then subjected to 1200 °C sintering for 5 h in a high vacuum furnace
with a heating/cooling rate of 10 °C/min. However, for fabrication of the bulk samples, the green compacts were directly heated to 1200 °C with the aforementioned sintering time and heating/cooling rate. It should be noted that one series of bulk and porous samples was also synthesized from the blended powder mixture without using the ballmilling process (Fig. 1a). A quantitative chemical analysis of the bulk samples after sintering was carried out using the inductively coupled plasma-atomic emission spectrometry (ICP-AES) and Leco combustion methods. The morphology and microstructure of the powder particles as well as the structure of the porous samples were characterized by means of a scanning electron microscope (Leica S440) combined with secondary electron/backscattered electron imaging (SEI/BEI). Particle-size distribution was measured using Malvern Instruments Mastersizer 2000 with a Hydro 2000S side feeder. The porosity of each porous sample was calculated by measuring their weight and dimensions. Accordingly, 6 samples from each group were weighed by a precision digital balance (Sartorius BP 221S) and their dimensions were measured using a digital caliper. The height and diameter of cylindrical samples were measured at three different positions in order to minimize the measurement error. Phase formation was characterized using XRD in the powder and bulk samples by Cu Kα radiation (35 kV, 28 mA) at a scanning rate of 1°/min using a Philips PW 1820 diffractometer. The Vickers hardness of the sintered bulk samples was measured using a 9.8 N (1 kgf) load for 15 s. Average hardness values were obtained from six indents on each of the bulk samples. 3. Results and discussion 3.1. Morphological and microstructural analyses of Ti–16Sn–4Nb alloy particles The morphology of the blended Ti–16Sn–4Nb powders is shown in Fig. 1a. The powder particles did not undergo any noticeable change after 1 h of ball-milling and the majority of the particles retained their original shape at this stage, as shown in Fig. 1b. The particles after the short ball-milling time of 1 h had not collided sufficiently with the balls, although they became slightly work hardened and agglomerated. The small difference in work hardening rate between
Fig. 1. SEM micrographs of (a) blended powders and the Ti–16Sn–4Nb powders ball-milled for: (b) 1 h; (c) 10 h; and (d) 20 h.
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the samples made from the powders ball-milled for 1 h and those of made from the blended elemental powders can be explained by the lower degree of compaction (larger degree of porosity) and slightly higher hardness of the former. The effect of work hardening on powder compaction will be discussed further in Section 3.5. After ballmilling for 10 h the majority of the particles became flattened and showed a relatively smaller size compared to that of the particles ballmilled for 1 h (Fig. 1c). Nevertheless, ball-milling up to 10 h did not result in any significant particle refinement. The presence of flattened and thin flaky particles after 10 h ball-milling is indicative of the micro-forging and cold welding of the particles at this stage. As seen in Fig. 1d, the morphology and size of powders drastically changed after 20 h ball-milling, and the powders exhibited a rather equiaxed shape and a considerable smaller size than the powders ball-milled for 1 and 10 h. It is notable that the absence of PCA in the ball-milling process resulted in an intense cold welding of powders to the milling tools for the long ball-milling time of 20 h and, thus, a considerably small amount of powder was recovered after 20 h ball-milling. As a result of the work hardening effect and high density of defects such as dislocations and vacancies for long ball-milling times, the fracture and welding of powders reach an equilibrium state. This state leads to the formation of rather equiaxed and small particles. However, due to the higher bonding strength of the finer particle size, the ability of particles to stand further plastic deformation is decreased and a higher force is required to fracture the small particles during the ballmilling process [20]. It is worth mentioning that the reduction of particle size at a given milling time can also be influenced by milling technique. For instance, high-energy ring mills are more efficient to reduce the particle size than ball-milling process. In the ring mill technique, particles are under the simultaneous effect of compression and shear forces [29–31]. Fig. 2 shows the particle-size distribution of Ti–16Sn–4Nb alloy after ball-milling for 1, 10, and 20 h. Besides a slight decrease in particle size, there is no considerable difference in powder size distribution up to 10 h ball-milling. At this stage, the colliding balls are mostly used to flatten the particles through shearing and the compressive force. The particle-size distribution of the powder ball-milled for 20 h was wider and showed a considerable shift to the left-hand end of the particle-size scale, indicating a decrease in particle size after ball-milling for 20 h. The broadening of particle-size distribution at longer ball-milling times is a typical behavior of the high-energy ball milling process [19,32,33]. There is also an overall tendency to weld small particles into larger pieces, as evidenced by a small hump at the right-hand end of the particle-size scale. The agglomeration of 20 h ball-milled powders is mostly due to the increasing resistance to fracture and the strong tendency of cohesion between particles with
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decreasing particle size. In the present study, ball-milling for 1 and 10 h resulted in a more symmetrical particle-size distribution. Fig. 3 shows the SEM backscattered electron images of the crosssection of the Ti–16Sn–4Nb alloy particles after ball-milling for 1, 10, and 20 h. Upon ball-milling for 1 h, the majority of particles remained unalloyed. The particles are distinguished by dark contrast areas for Ti and bright contrast areas for the Sn or Nb (Fig. 3a). At this stage, the ball-milling time is too short to achieve alloy formation. Principally, for a short ball-milling time, the particles are not subjected to an adequate compressive force or kinetic energy input, and therefore cannot be mechanically alloyed [20,34]. With continued ball-milling to 10 h, only small fragments of Sn and Nb were trapped into Ti powders and the particles were only mechanically bonded together, as shown in Fig. 3b. Increasing the ball-milling time to 20 h resulted in significant changes in the microstructure of the particles. As seen in Fig. 3c, the solute elements of Sn and Nb were embedded in the interfacial boundaries and were incorporated into the Ti matrix. As such, each particle displayed a lamellar-like structure consisting of a fine and relatively homogeneous distribution of the solute elements within the Ti matrix. In general, a more homogeneous distribution is expected when the ball-milling time is increased. Homogeneity of the particles is the result of a balance between fracturing and cold welding during the MA process [20]. The heavy plastic deformation of particles at longer ball-milling times gives rise to a large number of defects, providing a higher diffusivity of the constituent elements [34,35]. In other words,
8
Volume (%)
MA-1 h
MA-10 h
6
MA-20 h 4
2
0 0.01
0.1
1
10
100
1000
10000
Particle size ( m) Fig. 2. Particle size distribution of Ti–16Sn–4Nb alloy at different ball-milling times.
Fig. 3. The SEM backscattered electron images of the cross-section of the Ti–16Sn–4Nb particles for different ball-milling times: (a) 1 h; (b) 10 h; and (c) 20 h.
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continuous fracturing and welding of particles at long ball-milling times reduce the diffusion distance of alloying elements to micrometer range, leading to the lamellar microstructure, as seen in Fig. 3c. In contrast to a thermally induced diffusion process, the diffusion during the MA process is mostly controlled by mechanical energies [35]. The amount of powder recovered from the container at different ball-milling times is plotted in Fig. 4. It is clear that an increase in the ball-milling time led to a decrease in the powder yield. Almost the whole amount of powder was recovered after ball-milling for 1 h. By increasing the ball-milling time to 10 h, about 85% of the powder was recovered from the containers. This value is considered to be a high powder yield since no PCA was used in the present study. When the ball-milling time increased to 20 h, only a low powder yield of about 30% was obtained. At this point, a large amount of powder remained in the grinding media as evidenced by a relatively thick layer of powder over the balls and the inner wall of the container. Principally, due to the frequent mechanical impact of the balls on the powder particles, part of the milled powders are cold-welded over the grinding media, forming a coating layer [36,37]. However, the thickness of the coating layer is mainly dependent on the ball-milling time, the ductility of the initial powders, the amount of PCA and the intensity of milling (e.g. ball to powder weight ratio and milling speed). 3.2. XRD analysis on powder and bulk Ti–16Sn–4Nb alloy samples The XRD patterns of the Ti–16Sn–4Nb alloy powder mixtures and bulk samples sintered from the powders ball-milled for different times are shown in Fig. 5 and Fig. 6, respectively. The initial powder mixture consists of the elemental metals of Ti, Sn, and Nb (Fig. 5a). After 1 and 10 h ball-milling, the intensity of peaks corresponding to Sn and Nb did not show any significant change (Fig. 5b and c). However, on continuous ball-milling for 20 h, the sharp peaks from the crystalline elements of Sn and Nb disappeared, as seen in Fig. 5d. Peak broadening and the reduction in intensity of the diffraction peaks of ball-milled powders are associated with the refinement in crystallite size, lattice internal strain, and the instrumental effects [19]. With 10 h ball-milling or less, little solid solution took place between the powder particles, as evidenced by the presence of the diffraction peaks of Sn and Nb. The presence of Sn and Nb diffraction peaks at this stage is in agreement with the inhomogeneity of powder microstructure, as shown in Fig. 3b. Increasing the ball-milling time to 20 h led to the disappearance of the Sn and Nb diffraction due to the grain size reduction and/or accumulation of mechanical strains (Fig. 5d). The formation of an hcp α-Ti phase after 20 h ball-milling can be attributed to the dissolution of the alloying elements in the Ti lattice. The aforementioned result is consistent with the formation of the lamellar-like microstructure
Ti-16Sn-4Nb (wt.%) −Ti Sn Nb
Intensity (a.u.)
924
d 30
70 60 50 40
40
50 2 (deg.)
60
70
Fig. 5. XRD patterns of (a) elemental powders and the Ti–16Sn–4Nb powders ballmilled for: (b) 1 h; (c) 10 h; and (d) 20 h.
seen in Fig. 3c. However, the disappearance of the Sn and Nb peaks can be mainly associated with the formation of a titanium solid solution. With sintering of the ball-milled powders at a high temperature, no diffraction peaks of Sn and Nb were detected and only the Ti diffraction peaks were identified (Fig. 6). The intensity of Ti diffraction peaks in the sintered bulk samples decreased with the increase of ball-milling time. The diffraction patterns of the sintered bulk samples made from the powders ball-milled for 1 and 10 h also indicated a small fraction of β-Ti phase, as seen in Fig. 6b and c. This peak represents the small amount of Nb used as a β stabilizer in the synthesis of the present Ti alloy. However, after ball-milling for 20 h, the β phase became almost indistinguishable and merged into a single α-Ti phase, suggesting the complete solid solution of Nb in α-Ti (Fig. 6d). 3.3. Chemical analysis of bulk Ti–16Sn–4Nb alloy samples The content of contamination in the bulk Ti–16Sn–4Nb alloy samples made from the powders ball-milled for various ball-milling
Ti-16Sn-4Nb (wt.%) Ti
Intensity (a.u.)
Weight Recovery (%)
80
b c
100 90
a
Ti Sn Nb
a b c
30
d
20 10 0 0
5
10
15
20
25
Milling Time (h) Fig. 4. The amount of powder recovered from the ball-milling container as a function of ball-milling time.
30
40
50
60
70
(deg.) Fig. 6. XRD patterns of (a) elemental powders and the bulk Ti–16Sn–4Nb made from the powders ball-milled for: (b) 1 h; (c) 10 h; and (d) 20 h.
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3.4. Hardness of the bulk Ti–16Sn–4Nb alloy sample The hardness values were measured on the bulk Ti–16Sn–4Nb samples made from a blended powder mixture as well as those of samples made from the powders ball-milled for different times. The variation in hardness versus different ball-milling times is plotted in Fig. 7. It can be seen that the hardness increased with increasing ballmilling time. The lowest hardness belonged to the sample made from the blended powder mixture showing the value of about 300 HV1.0. However, the hardness did not show any considerable change after ball-milling for 1 h. For the samples made from the powders ballmilled for 10 h, the Vickers hardness increased to approximately 370 HV1.0. Subsequently, a 40% increase in hardness was achieved when ball-milling time was increased from 10 h to 20 h, reaching a maximum of approximately 520 HV1.0. A progressive increase in hardness of the bulk Ti–16Sn–4Nb alloy made from the 20 h ball-milled powders could be put down to: (i) the work hardening of the powders; (ii) crystallite size refinement; (iii) dispersoid strengthening due to the increased contamination at longer ball-milling times; and (iv) solid solution strengthening. All the aforementioned factors impede dislocation movement and lead to a strengthening of the material. However, compared to
500
Hardness (HV1.0)
times is given in Table 2. The results indicate that the increase of ballmilling time leads to a noticeable contamination of oxygen in the bulk samples. The contaminations from oxygen have frequently been reported during the milling of Ti alloys [38–41], which can be traced to the milling atmosphere, chemical purity of the starting powders, and the decomposition of the PCA used during milling. However, in the present study the resultant contamination from oxygen is mainly attributed to the leakage of air into an inadequate sealed container which interacts with the large and fresh surface areas of powder particles during high-energy ball milling. It is also notable that the longer ball-milling times of 10 and 20 h also resulted in increased contamination of Fe and Cr. The Fe and Cr contaminations are from the wear of the stainless steel milling balls and containers during the ballmilling process. These metallic contaminations are more pronounced for the samples made from the powders ball-milled for 20 h. Higher levels of C, Fe and Cr contamination can be expected if the powders are ball-milled with a PCA addition, as previously reported by the same authors [38,42]. It is possible that the coating film formed on the milling tools during the longer ball-milling times prevents excessive wear of the milling media. This is indicated by the smaller difference in the amount of Fe and Cr contamination between 20 and 10 h ball-milled powders as compared to the difference in metallic contamination between 10 and 1 h ball-milled powders. However, the thickness of the film should be kept to a minimum to increase the powder yield and avoid forming a heterogeneous product [19,43]. The increase in content of contamination with ball-milling time is a common problem with the MA technique. In fact, an argon atmosphere is incapable of reducing the oxygen partial pressure sufficiently and it is practically impossible to prevent the surrounding atmosphere from leaking into the milling container during the MA process [44,45].
925
400 300 200 100 0 0
5
10
15
20
Milling Time (h) Fig. 7. The Vickers hardness of the bulk Ti–16Sn–4Nb alloy made from the powders ballmilled for different times.
the hardness values obtained from the bulk samples with different amounts of PCA [38], the dispersoid strengthening may not be an influential factor in the high hardness of the resultant samples at longer ball-milling times. Finally, most pure metals can be strengthened by the alloying elements. This is due to the fact that the presence of alloying atoms along the core of dislocation reduces the strain energy of the crystal, suggesting more stress to move the dislocation away from each other [17]. In particular, it has been reported that Sn has a strengthening effect in Ti lattice [12,46].
3.5. Structural characteristics of the porous Ti–16Sn–4Nb alloy samples 3.5.1. Degree of porosity The degree of porosity in the porous Ti–16Sn–4Nb samples as a function of ball-milling time is shown in Fig. 8. Although the weight ratios of the metal powders to the amount of space-holding material in this experiment were calculated to produce overall porosities of around 70%, the resultant samples made from the powders ball-milled for 20 h revealed a noticeable deviation from the calculated value. As seen in Fig. 8, the porosity of samples increases with an increase in ball-milling time, reaching a maximum value of around 74% after ballmilling for 20 h. This deviation can also be seen by the increase in error bars for longer ball-milling times. The error bar was the mean square deviation obtained from the measurements of six different porous samples for a given milling time. Generally, the porous powder metallurgical products synthesized from the particles with a broad particle-size distribution have a lower reproducibility and uniformity of the final porosity [47]. The minimum porosity was obtained for the porous samples synthesized from the blended powder mixture, showed a porosity value of 68%.
Table 2 The content of contamination in bulk Ti–16Sn–4Nb alloy made from the powders ballmilled for various periods and sintered at 1200 °C for 5 h. Ball-milling time
Contamination content (wt.%) O
N
C
Fe
Cr
1h 10 h 20 h
1.25 2.15 2.65
0.01 0.01 0.03
0.1 0.1 0.22
0.2 0.41 0.47
0.11 0.28 0.35
Fig. 8. The degree of porosity for the porous Ti–16Sn–4Nb samples versus different ballmilling times.
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3.5.2. Morphology of the porous alloy samples Fig. 9 shows the SEM micrographs of the porous Ti–16Sn–4Nb samples made from the powders ball-milled for different times at two different magnifications (250× and 1000×). It should be noted that the morphology of pores and pore walls in a porous material can be influenced by the morphology of its constituent powders [48,49]. The pore walls and inner pore surfaces in the porous samples made from the blended powder mixture or powder ball-milled for 1 h were smoother and only a few micro-pores or interstices could be seen, as shown in Fig. 9(a,b) and Fig. 9(c,d), respectively. In contrast, the rather round and equiaxed particles obtained from the powders ball-milled for 20 h, resulted in corrugated and rough pore walls and
inner pore surfaces (Fig. 9(g,h)). An increase in ball-milling time also resulted in a larger number of dents and concavities on the surface of the porous material. Moreover, in the porous sample made from the powders ball-milled for 20 h, a smaller inter-particle neck size and a larger number of micro-pores could be seen while the individual particles could usually be distinguished. However, the effect of powder morphology on the size of the pores was negligible, since the pore size is strongly controlled by the size of the space holder particles and the sintering temperature. Fig. 10 shows the backscattered electron images of the cell wall microstructure in the porous Ti–16Sn–4Nb samples for different ballmilling times. As shown in Fig. 10a, the porous sample made from the
Fig. 9. SEM micrographs of the porous Ti–16Sn–4Nb samples made from (a,b) a blended powder mixture and powders ball-milled for: (c,d) 1 h; (e,f) 10 h; and (g,h) 20 h.
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Fig. 10. The SEM backscattered electron images of the cell wall microstructure in the porous Ti–16Sn–4Nb samples made from the powders ball-milled for: (a) 1 h; (b) 10 h; and (c) 20 h.
powders ball-milled for 1 h exhibited a heterogeneous microstructure, as evidenced by the high volume fraction of the β Nb-rich phase (bright areas) in α-Ti matrix (dark areas). The composition of the aforementioned phases has already been reported in our previous papers using EDX analysis [42,50]. With increasing ball-milling time to 10 h, the volume fraction of the β phase slightly decreased and a more homogeneous microstructure was achieved (Fig. 10b). However, after 20 h ball-milling, the bright Nb-rich colonies were mostly homogenized in the dark Ti matrix showing a considerable decrease in size and volume fraction of the Nb-rich phase, as shown in Fig. 10c. The biomedical characteristics and mechanical properties of a porous metallic implant are highly affected by its total porosity, size and the morphology of the pores [51–53]. Apart from the type of processing method, the total porosity, density, and pore morphology in a powder metallurgical porous material are significantly controlled by its particle packing characteristics. However, the selection of the processing method is also based on the powder characteristics and the type of porosity required by the application. In general, the density of a green powder compact consisting of coarse powder (i.e. blended and 1 h ball-milled powders in the present study) is saturated under a low compaction pressure. This phenomenon is caused by weak interparticle attractive forces and low friction between the powder particles, due to the small number of contacts between the powder particles [54,55]. It can be assumed that the fine powders obtained
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with a ball-milling time of 20 h have stronger inter-particle attractive forces. These forces are a result of the electrostatic field, Van der Waals and moisture/surface adsorption forces that become much more significant with a decrease in particle size [54,56]. In addition, the powders ball-milled for long times cannot move and deform effectively during compaction. This may be primarily due to the effect of a decrease in particle size, solid solution strengthening and work hardening of the particles during the ball-milling process [57,58]. Amongst these factors, the work hardening effect appears to be the main determinant in governing the plastic behavior of powder particles during compaction. In the ball-milling process, the work hardening is caused by the interaction of dislocations with each other through the continuous mechanical impact of the balls on the powder particles. Parvin et al. [59], in a study on mechanically alloyed SiCreinforced Al composites, concluded that the effect of work hardening due to milling time is more pronounced than the morphological changes of the powder particles. Sadek and Salem [60], also suggested that the work hardening of the nanopowders after the ball-milling process resulted in the resistance of the particles to plastic deformation and hence retarded the compaction. Besides the shape, size and distribution of the particles, the particle packing characteristics of a green compact depend directly on the properties of the materials such as yield strength, impurities, and hardness [49,54]. These factors seem to be dominant for the total porosity of the samples after longer ball-milling times. Hewitt and Kibble [61] recently investigated the effects of ball-milling time of 0.5, 1, 3, and 5 h on the synthesis and consolidation of WC-10%Co powder at room temperature. They have shown that the relative density of compacted powders decreases linearly with milling time due to the combined effects of particle size, morphology and hardness. At 400 MPa compaction pressure, 64% relative density was achieved for the powders milled for 30 min compared to 61% after milling for 300 min. The small powder particles give rise to a lower packing density. In other words, a higher degree of porosity or lower relative density in compacts is achieved when the surface area of powders is increased. The reason for this is that small particles result in a larger number of contacts per particle and particle bridging, thereby leading to an increase in inter-particle friction and resistance to compaction [49,54]. Furthermore, the small particles provide less available sites for mechanical interlocking. In the present study, the absence of PCA during the MA process also intensified the interparticle friction between the resultant ball-milled powders since the rearrangement of the particles, which is the first densification step of cold die compaction, is extensively controlled by the lubricating effect of PCA. In another study by the same authors [38], a decidedly upward trend in the relative density of the sintered bulk samples was observed with an increase in PCA content. In the present study, the porous samples made from the powders ball-milled for 20 h exhibited more interconnected micro- and macropores. A large number of interconnected porosities enhance the circulation and exchange of body fluids, increase the rate and distribution of osteogenesis around and throughout the implant, promote vascularization and nutritional supply, and improve implant osseointegration [62]. Since the vascular network is strongly influenced by the continuity of pores, the degree of interconnectivity between pores is more crucial for new bone formation than pore size and porosity [63–65]. The interconnectivity of the pores can be theoretically increased by an increase in pore size and/or porosity, which in turn compromises the mechanical strength of the implant. Hence, an optimum balance between porosity and strength must be achieved, to ensure that the implant can at least withstand the forces applied during the surgical procedure and initially at the implantation site. In the current study, the porous samples synthesized with different ball-milled powders exhibited the macro-pore-structure of 100–400 μm which is desired for optimal bone cell ingrowth [66,67].
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The ball-milling process can be a promising method in synthesizing a desired porous structure for optimal bone cell ingrowth with increased overall porosity, as long as the mechanical strength is accommodated. The diverse porous structures obtained with different ball-milling times can provide new substrates for evaluating the mechanical strength and influence of pore architecture on in vitro cell proliferation.
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4. Conclusion
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The effect of different ball-milling times on the structural characteristics of powders, bulk and porous Ti–16Sn–4Nb alloy was investigated for various periods of time from 1 to 20 h. Ball-milling for 20 h led to the formation of a relatively equiaxed shape and smaller size powder particles with more asymmetrical particle-size distribution than the powders ball-milled for 1 and 10 h. The diffraction patterns of the sintered bulk samples made from the powders ballmilled for 1 and 10 h indicated a small fraction of β-Ti phase. However, further ball-milling to 20 h resulted in merging the β phase into a single α-Ti phase and, presumably, the formation of a partial amorphous phase. An increase in ball-milling time also revealed some increase in the amount of contamination introduced by the milling atmosphere and grinding media. The hardness of the sintered bulk samples showed a decidedly upward trend with an increase in ballmilling time reaching to the maximum value of ~ 520 HV1.0 for the sample made from 20 h ball-milled powder. The pore walls and inner pore surfaces in the porous samples made from the blended powder mixture and powder ball-milled for 1 and 10 h were smoother with a small number of micro-pores. However, the porous samples made from the 20 h ball-milled powders exhibited corrugated pore walls and rough inner pore surfaces with a larger number of micro-pores mostly due to the combined effects of particle size, morphology, and work hardening. The resultant porous samples made from the 20 h ball-milled powders showed a higher degree of porosity with more interconnected micro- and macro-pores. Acknowledgments The authors acknowledge the financial support from the VCAMM (Victorian Centre for Advanced Materials Manufacturing) and ARC (Australian Research Council) through the ARC Discovery Project DP0770021. Peter Hodgson is also supported by the ARC through a Federation Fellowship. References [1] [2] [3] [4] [5] [6] [7]
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