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Effect of sintering conditions on the microstructural and mechanical characteristics of porous magnesium materials prepared by powder metallurgy
Materials Science and Engineering C 35 (2014) 21–28
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Materials Science and Engineering C journal homepage: www.elsevier.com/locate/msec
Effect of sintering conditions on the microstructural and mechanical characteristics of porous magnesium materials prepared by powder metallurgy Jaroslav Čapek ⁎, Dalibor Vojtěch Department of Metals and Corrosion Engineering, Institute of Chemical Technology, Prague, Technická 5, 166 28 Prague 6, Czech Republic
a r t i c l e
i n f o
Article history: Received 25 May 2013 Received in revised form 14 October 2013 Accepted 21 October 2013 Available online 1 November 2013 Keywords: Porous magnesium Sintering conditions Powder metallurgy Mechanical properties
a b s t r a c t There has recently been an increased demand for porous magnesium materials in many applications, especially in the medical field. Powder metallurgy appears to be a promising approach for the preparation of such materials. Many works have dealt with the preparation of porous magnesium; however, the effect of sintering conditions on material properties has rarely been investigated. In this work, we investigated porous magnesium samples that were prepared by powder metallurgy using ammonium bicarbonate spacer particles. The effects of the purity of the argon atmosphere and sintering time on the microstructure (SEM, EDX and XRD) and mechanical behaviour (universal loading machine and Vickers hardness tester) of porous magnesium were studied. The porosities of the prepared samples ranged from 24 to 29 vol.% depending on the sintering conditions. The purity of atmosphere played a significant role when the sintering time exceeded 6 h. Under a gettered argon atmosphere, a prolonged sintering time enhanced diffusion connections between magnesium particles and improved the mechanical properties of the samples, whereas under a technical argon atmosphere, oxidation at the particle surfaces caused deterioration in the mechanical properties of the samples. These results suggest that a refined atmosphere is required to improve the mechanical properties of porous magnesium. © 2013 Elsevier B.V. All rights reserved.
1. Introduction Magnesium and magnesium alloys have recently been studied for use in many applications, such as in the automotive and aerospace industries because of their low densities and good mechanical properties [1]. Many successful studies have been performed on biocompatible and biodegradable magnesium-based materials, which are considered suitable materials for orthopaedic applications, such as for nails, screws, splints, etc. [2–6]. For some applications, porous implants, called scaffolds, are required because they possess mechanical properties, such as the modulus of elasticity, that are relatively similar to those found in natural bone tissue [7–14]. The mechanical biocompatibility is important to aid in the remodelling of new tissue [7–14]. Moreover, an interconnected porous structure allows the transport of body fluids to damaged or wounded tissue and supports the incorporation of new tissue in the implant [7,15,16]. Therefore, porous magnesium materials and their preparation methods have been extensively studied in recent years [7–14]. Many methods have been developed for fabricating porous metallic materials [17]. However, because biomaterials should contain interconnected pores [12] and should not be contaminated with harmful impurities [4], only a few of these methods are used for the production of biomaterials. In the available literature, there are five “non-machining” ⁎ Corresponding author. Tel.: +420 220444055. E-mail address: [email protected] (J. Čapek). 0928-4931/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msec.2013.10.014
approaches that have been reported for the fabrication of porous magnesium materials: (1) injection of an inert gas into a melt [4], (2) directional solidification of the metal–gas eutectic (the GASAR process) [9], (3) plaster casting [4], (4) negative salt pattern moulding [12] and (5) powder metallurgical techniques [11,13,14]. However, the first two methods mentioned above do not necessarily produce open-cell structures, and the following two methods may contaminate or corrode the final product during pattern removal [18]. Suitable modifications to powder metallurgy (PM), for example, using space-holder particles, allow the fabrication of materials with interconnected pores. This modification consists of preparing a green compact that contains a powdered mixture of magnesium and a space-holder material, which is later removed by leaching or thermal decomposition. Subsequently, the porous green body is sintered at high temperatures [8,10,13,14]. In principle, any solid matter may be used as the space-holder material; however, in practice, this selection is limited because the spacer material has to be removed without contaminating the product. In the literature, urea and ammonium bicarbonate have been successfully used as spacer materials in the preparation of PM porous magnesium [8,11,13,14,19]. Hao et al. [8] removed the urea by leaching the material in a NaOH-water solution, whereas other authors removed the spaceholder particles by thermal decomposition [10,11,13,14]. The majority of these authors used urea as the spacer material, even though its complete decomposition occurs at temperatures above the melting point of magnesium. Because urea only partially decomposes at lower
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Fig. 1. SEM micrographs (SE detector) of sample cross sections prepared under the following sintering conditions: (a) technical argon for 6 h, (b) technical argon for 24 h, (c) gettered argon for 6 h and (d) gettered argon for 24 h.
temperatures, contamination may occur [20]. For example, Zhuang et al. [13] found traces of carbon and oxygen on pore walls, which they attributed to contamination by urea residues. Ammonium bicarbonate, which decomposes at significantly lower temperatures (36 °C–60 °C), appears to be a more suitable spacer material for the preparation of PM porous magnesium. In addition, leaching the spacer material has the disadvantage of possibly hermetically enclosing some of the spacer particles in the magnesium matrix, which prevents their dissolution. Moreover, water-based dissolving agents are corrosive to magnesium [21]. The mechanical properties of PM porous magnesium are influenced not only by the total porosity but also by the pore size, distribution and shape and the connection between magnesium particles [8,10,13]. These structural characteristics can be adjusted by selecting the optimal sintering time and temperature, compacting pressure and size, shape and volume ratio of the starting material powders [8,11,13,19,22]. Grain coarsening within each magnesium particle, which is expected during the sintering process, may also affect the mechanical properties of the final material [23,24]. In our previous work [19], we showed that porous magnesium can, in principle, be successfully prepared by PM using ammonium bicarbonate as the space-holder material. Although several works on the preparation of porous magnesium have been reported, to the best of our knowledge, no systematic study on the influence of the processing
parameters, namely, the sintering time and atmosphere purity, on the microstructure and mechanical properties of porous magnesium is available. Therefore, our study focused on sintering kinetics and the effect of the sintering atmosphere on the important characteristics of PM porous magnesium.
Table 1 The influence of sintering conditions on sample porosity (in vol.%). Atmosphere/sintering time
0h
3h
6h
12 h
24 h
Technical argon Gettered argon
29 ± 2 29 ± 2
28 ± 1 28 ± 1
28 ± 2 27 ± 2
29 ± 3 25 ± 3
31 ± 4 24 ± 2
Fig. 2. XRD patterns of samples sintered for 6 h.
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compacts were annealed for 4 h at 130 °C in a muffle furnace in air. During this step, the decomposition of ammonium bicarbonate and the evaporation of hexane occurred. Afterwards, sintering was performed at 550 °C in a tube furnace. The green compacts were sintered for 0, 3, 6, 12 and 24 h. The sintering process was performed under two types of flowing atmosphere at a flow rate of 0.1 l/min: (1) argon with technical purity (99.996 vol.%) and (2) argon purified by a 55-mm-thick layer of Mg chips (250–500 μm in size), which was placed around the sintered material and acted as a getter. After sintering, the average material porosity was determined according to Eq. (1) [8,13]: P ¼ 1–ρ=ρMg 100%;
Fig. 3. XRD patterns of samples sintered for 24 h.
2. Experimental A purchased Mg powder (purity of 99.6 wt.%, mesh −100 + 200, Alfa Aesar) and NH4HCO3 powder (p.a. purity, 250–500 μm) were used as starting materials. An Mg/NH4HCO3 volume ratio of 90:10 was used because, according to our previous work [19], this ratio imparts a good combination of strength and corrosion resistance to the resulting material. Materials prepared using this ratio have been shown to possess better mechanical properties that are more comparable to natural bone tissue than porous biomaterials, which have recently been used for medical applications [19]. To avoid segregation, the powders were intensively blended with 30 vol.% hexane for 30 min. Subsequently, pre-weighed mixtures were pressed into cylindrical green compacts (10 mm in diameter and 30 mm in length) at a pressure of 265 MPa using a LabTest 5.250SP1-VM universal loading machine. The green compacts were then subjected to a two-step procedure. First, the
ð1Þ
where P is the porosity, ρ is the density of the porous sample (calculated from the dimensions and weight) and ρMg is the density of pure magnesium (ρMg = 1738 kg/m3). This approach for determining porosity was used because it describes the volumetric 3D porosity more precisely than analysing 2D micrographs. The material mechanical properties were characterised by flexural, compression and Vickers hardness tests that were performed at room temperature. Samples 26 mm and 15 mm in length were used for flexural and compression testing, respectively. Five samples of each series were used for the flexural tests, and three samples of each series were used for the compression tests. The average hardness was calculated from ten values. Standard deviations for all tests were calculated and are shown as error bars in the figure plots. The deformation rates during the flexural and compression tests were 0.5 mm/min and 1 mm/min, respectively. A LabTest 5.250SP1-VM universal loading machine was used for these tests. After flexural testing, the fracture surfaces were observed by a TESCAN VEGA-3 LMU scanning electron microscope (SEM). Metallographic cross sections were also prepared and examined by SEM. Afterwards, the metallographic cross sections were etched (2 g of picric acid, 10 ml of 99% acetic acid, 10 ml of water and 70 ml of ethanol), and the microstructures of the powder particles were observed using an Olympus PME 3 light metallographic microscope. The phase composition and elemental distributions were examined by X-ray diffraction using a PANalytical X'Pert PRO X-ray diffractometer equipped with a Cu anode (XRD) and a TESCAN VEGA-3 LMU SEM equipped with an Oxford Instruments INCA 350 EDX analyser (SEM-EDX). The Vickers hardness was measured for each series with a load of 3 kg (HV 3).
Fig. 4. SEM micrographs (BSE detector) and elemental maps determined by EDX of samples sintered 24 h under (a) technical argon atmosphere and (b) gettered argon atmosphere. Scale bar 50 μm.
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Fig. 5. Microstructure of powder particles during the material fabrication process: (a) the initial powder, (b) compacted powder after annealing at 130 °C for 4 h, (c) compacted powder after annealing at 130 °C for 4 h and sintering at 550 °C for 3 h under technical argon and (d) compacted powder after annealing at 130 °C for 4 h and sintering at 550 °C for 24 h under technical argon.
3. Results and discussion 3.1. Microstructure and porosity The microstructures of samples sintered under two types of atmospheres for different periods of time are shown in Fig. 1. In the magnesium matrix, two types of pores were observed: (1) pores approximately 250–500 μm in size that were formed from the decomposition of spacer particles (“Type I” in Fig. 1) and (2) smaller pores that most likely originated from imperfect compaction during pressing and the expansion of trapped gas during sintering (“Type II” in Fig. 1). No significant difference in microstructure was observed for
Fig. 6. Flexural stress–strain curves of selected samples.
samples sintered for 6 h under the different atmospheres. A prolonged sintering time caused annihilation of “Type II” pores under the gettered atmosphere, whereas under the technical argon atmosphere, the amount of these pores did not significantly change. The sample porosity was measured, and the effects of the sintering conditions on porosity are shown in Table 1. This table indicates that samples sintered under the gettered argon atmosphere became more compact with longer sintering times, i.e., contained less pores. In contrast, the technical argon atmosphere produced materials whose porosities appear almost the same.
Fig. 7. Ultimate flexural strength of prepared samples versus sintering time.
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Fig. 8. Fracture surfaces of samples after flexural testing under the following sintering conditions: (a) technical argon for 6 h - overview, (b) technical argon for 6 h - detail, (c) technical argon for 24 h and (d) gettered argon for 24 h. The white arrow in Fig. 8c denotes oxide particles. In Fig. 8d, the arrow indicates traces of plastic deformation.
This phenomenon suggests that sintering under the gettered argon atmosphere enhanced the diffusion connection between the magnesium particles, which led to annihilation of some “Type II” pores (Fig. 1). In contrast, the use of technical argon caused a slight increase in total porosity. As shown below, porosity is directly related to the presence of oxides in the material.
The XRD patterns of the samples sintered under various conditions are shown in Figs. 2 and 3. The XRD patterns indicate that there is no significant difference in the amount of magnesium oxide in the samples sintered for 6 h, regardless of the atmosphere type (Fig. 2). In this figure, the peaks that are assigned to MgO are similar for both samples sintered for 6 h. However,
Fig. 9. Compressive stress–strain curves of selected samples.
Fig. 10. The compressive yield strength versus sintering time of prepared samples.
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samples that were sintered for 24 h in technical argon (Fig. 3) contained significantly higher amounts of MgO than those sintered in gettered argon. No traces of ammonium bicarbonate or the reaction products of magnesium and ammonium bicarbonate were found by the XRD analysis. To determine the difference between the samples sintered for 24 h, EDX analysis was performed. The results are shown in Fig. 4. The X-ray elemental maps show the presence of Mg particles surrounded by oxide layers. Comparing Fig. 4a and b reveals that more extensive oxidation occurs at the grain boundaries during sintering under the technical argon atmosphere than under the gettered atmosphere. This difference, which is observed after sintering for more than 6 h, is in good agreement with the results acquired from XRD analysis (Fig. 3). Traces of oxygen and water are present in technical argon and cause surface oxidation, which weakens the diffusion connection between the magnesium particles. Thus, annihilation of the smaller “Type II” pores is slowed during longer sintering periods (Table 1). Moreover, oxides on the sample and particle surface may spall off and increase the measured porosity. In contrast, a lower extent of oxidation was observed under the gettered atmosphere, which supports the annihilation of some “Type II” pores because of the diffusive processes of the magnesium particles, as shown in Table 1. The presence of nitrogen and carbon was also investigated as contamination by either ammonium bicarbonate or its decomposition products. No traces of these elements were found; therefore, their distribution is not shown in the XRD elemental maps. Grain coarsening within magnesium particles during sintering was also investigated because it may have an impact on sample properties. Fig. 5 shows detailed views of the microstructure of individual powder particles. The initial powder (Fig. 5a) possesses a relatively fine microstructure (grain size up to 13 μm). After compaction and annealing at 130 °C for 4 h (Fig. 5b), the grain size slightly enlarged; however, the microstructure remained fine (approximate grain size of 15 μm). The grain size significantly increased after sintering at 550 °C for 3 h in technical argon (approximate grain size of 45 μm) (Fig. 5c). Sintering for longer periods did not cause any further significant grain growth. The average grain size after sintering for 24 h at 550 °C was approximately 48 μm (Fig. 5d). The type of atmosphere did not influence grain growth during sintering.
Fig. 11. The ultimate compressive strength versus sintering time of prepared samples.
A study of the fracture surfaces helped to explain this behaviour. The fracture surfaces of samples after flexural testing are illustrated in Fig. 8. For all the samples, the fracture surfaces were similar (Fig. 8a); however, slight differences were observed at higher magnifications. Generally, the fracture surfaces were brittle, but traces of plastic deformation were observed, especially in samples sintered under gettered argon for longer time periods (12 and 24 h) (Fig. 8d). The fracture surfaces were nearly identical and oxide free after sintering for up to 6 h, independent of the atmosphere (Fig. 8b). After longer sintering times, some differences were observed between the fracture surfaces of samples prepared under the different atmospheres. Samples sintered in gettered argon had nearly oxide-free fracture surfaces with features of plastic deformation (Fig. 8d), whereas a large amount of oxide was present on the fracture surfaces of samples sintered in technical argon (Fig. 8c). This observation confirms the XRD and EDX results that were mentioned above (Figs. 2–4).
3.2. Mechanical properties The flexural stress–strain curves of selected samples are shown in Fig. 6. Samples sintered in technical argon had a significantly lower modulus of elasticity than samples sintered under the gettered atmosphere. This result can be explained by the increased amount of oxygen at the interface between powder particles, which results in decreased cohesiveness between the particles. Oxides also impact flexural strength, which will be shown and discussed below. The flexural stress–strain curves of all the materials contained only elastic regions; no regions of macroscopic plastic deformation were observed (Fig. 6). The ultimate flexural strength (UFS) as a function of sintering time is plotted in Fig. 7. The ultimate flexural strength increased with longer sintering times, up to 6 h, independent of the atmosphere. The samples attained approximately the same UFS, 12 MPa, after 6 h of sintering under both atmospheres. For longer sintering times under the gettered argon atmosphere, the UFS increased to approximately 15 MPa. The difference between samples sintered for 12 and 24 h under the gettered argon atmosphere is negligible. In contrast, prolonging the sintering period to more than 6 h under technical argon had a completely opposite effect. The UFS rapidly decreased to 5 MPa after 12 h of sintering and continued to decrease to approximately 3 MPa after 24 h of sintering.
Fig. 12. The HV 3 versus sintering time of prepared samples.
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Table 2 Mechanical properties of porous biomaterials. Material
Porosity (vol.%)
Pore size (μm)
UFS (MPa)
CYS (MPa)
UCS (MPa)
Reference
Natural bone Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Ti Porous hydroxyapatite Porous hydroxyapatite Porous composite (poly-L-lactide + 20–50 wt.% of bioglass) Porous polycaprolactone Porous polycaprolactone Porous polycaprolactone Porous polylactide-co-glycolide Porous composite of polylactide-co-glycolide and 45S5 bioglass
— 29–31 23–38 14–44 36–55 52–70 50 28 35–55 78 50–77 — 77–88 48–77 37–55 55–56 31 43
— 250–500 250–500 250–500 200–400 ~1250 200–500 170 100–400 200–500 200–400 366–444 ~100 — — — 116 89
2–150 3–15 4–17 2–5 14–27 — — — — — 2–7 — 1–4 — —
— 13–53 — — — — — — — — — — — 2–3 — 2–3 — —
2–180 20–70 — — 15–31 4–14 2 24 12–17 35 1–17 30 ~0.4 — 2–3
[13] This study [19] [22] [13] [8] [14] [7] [11] [14] [13] [25] [13] [25] [25] [26] [27] [27]
Fig. 9 shows the compressive stress–strain curves of selected samples. From the figure, the trends of the compressive and flexural behaviour of the material are similar, with the exception of the effect of atmospheric condition on the compressive modulus of elasticity, which is not as strong. The compressive yield strength (CYS) and ultimate compressive strength (UCS) as functions of the sintering time are plotted in Figs. 10 and 11. Under the gettered argon atmosphere, both the yield strength (CYS) and the ultimate strength (UCS) increased with longer sintering times. The maximum CYS and UCS values after sintering for 24 h are 53 and 69 MPa, respectively. These high strength values result from a low amount of oxide, decreased porosity and good contact between the particles in the samples (Figs. 2–4, Table 1). In contrast, for samples sintered under technical argon, the maximum CYS and UCS were attained after 6 h of sintering. Prolonging the sintering period under this atmosphere led to a significant decrease in compressive properties. This finding can be attributed to the extended oxidation that occurs during sintering, which weakens the diffusion bonds between particles. An interesting observation in Figs. 10 and 11 is that the CYS of samples sintered in gettered argon increased almost linearly with sintering time, whereas the UCS increased for up to 6 h of sintering and then increased significantly more slowly. These results indicate that the “Type II” pores, which decrease in quantity during long sintering periods (Table 1), influence the yield strength rather than the ultimate strength. The ultimate compressive strength is probably more influenced by the larger “Type I” pores, whose amount is almost constant for all samples. The reason may be that the large pores concentrate the stress in their vicinity and initiate the crack nucleation and subsequent collapse of the porous structure. On the other hand, the small “Type II” pores concentrate the stress to a smaller extent and plastic deformation takes place in their surroundings. Therefore, their partial vanishing during longer sintering times seems to increase the stress needed for the plastic strain to begin. Fig. 12 reveals the effect of sintering time on sample hardness. Samples sintered in gettered argon possessed a higher hardness than those sintered in technical argon. The effect of sintering time on hardness was small for samples sintered in technical argon. However, prolonging the sintering period in the gettered argon led to an increase in hardness, with a peak hardness occurring in samples sintered for 12 h. Sintering up to 24 h slightly decreased sample hardness. This result may be caused by oxidation beginning at the grain boundaries and decreases the cohesiveness of the magnesium particles. A comparison of our results with previously reported results is difficult because of differences in the porosities and pore sizes of the studied materials. Nevertheless, we conducted a small review that summarises the mechanical properties of natural bone tissue and porous materials
— —
0.5 0.4
that are considered suitable biomaterials. The results of this research are summarised in Table 2. From this table, it is clear that porous magnesium-based metallic materials exhibit mechanical properties similar to those of natural bone tissue. However, non-metallic porous materials (especially polymer-based materials) have weaker mechanical properties compared to natural bone tissue. It is also important to note that the materials prepared in our study have higher UFS and UCS values than porous magnesium materials that have been prepared in the majority of other works and non-metallic porous materials. Moreover, the materials prepared in this work possess mechanical properties that are comparable to natural bone tissue. 4. Conclusions Magnesium samples with porosities of 29–31 vol.% were prepared by powder metallurgy under different sintering conditions. When compared to non-metallic biomaterials, the samples exhibited enhanced mechanical properties that were similar to human bone tissue. After sintering for up to 6 h, no significant effect of atmosphere purity on sample microstructure or mechanical behaviour was observed. Under the gettered argon atmosphere, longer sintering times decreased porosity but enhanced the mechanical properties of the sample. Under the technical argon atmosphere, the opposite trend was observed, which can be attributed to oxidation of the powder surface. These results suggest that when sintering is longer than 6 h, the purity of the argon atmosphere plays an important role in determining the mechanical properties of PM magnesium. Acknowledgements The authors would like to thank the Czech Science Foundation (project no. P108/12/G043) for supporting this research. References [1] F.H. Froes, Mater. Sci. Eng. A 184 (1994) 119–133. [2] N. Li, Y.F. Zheng, J. Mater. Sci. Technol. 29 (2013) 489–502. [3] H. Waizy, J.M. Seitz, J. Reifenrath, A. Weizbauer, F.W. Bach, A. Meyer-Lindenberg, B. Denkena, H. Windhagen, J. Mater. Sci. 48 (2013) 39–50. [4] M.P. Staiger, A.M. Pietak, J. Huadmai, G. Dias, Biomaterials 27 (2006) 1728–1734. [5] S. Virtanen, Mater. Sci. Eng. B 176 (2011) 1600–1608. [6] M. Hradilova, J. Kubasek, P. Lejcek, L. Tanger, Metal 2011: 20th Anniversary International Conference on Metallurgy and Materials, 2011, pp. 920–924, (full text available on http://www.metal2013.com/files/proceedings/metal_11/lists/papers/978. pdf, 31.7.2013). [7] X.N. Gu, W.R. Zhou, Y.F. Zheng, Y. Liu, Y.X. Li, Mater. Lett. 64 (2010) 1871–1874. [8] G. Hao, F. Han, W. Li, J. Porous. Mater. 16 (2009) 251–256.
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Contents lists available at ScienceDirect
Materials Science and Engineering C journal homepage: www.elsevier.com/locate/msec
Effect of sintering conditions on the microstructural and mechanical characteristics of porous magnesium materials prepared by powder metallurgy Jaroslav Čapek ⁎, Dalibor Vojtěch Department of Metals and Corrosion Engineering, Institute of Chemical Technology, Prague, Technická 5, 166 28 Prague 6, Czech Republic
a r t i c l e
i n f o
Article history: Received 25 May 2013 Received in revised form 14 October 2013 Accepted 21 October 2013 Available online 1 November 2013 Keywords: Porous magnesium Sintering conditions Powder metallurgy Mechanical properties
a b s t r a c t There has recently been an increased demand for porous magnesium materials in many applications, especially in the medical field. Powder metallurgy appears to be a promising approach for the preparation of such materials. Many works have dealt with the preparation of porous magnesium; however, the effect of sintering conditions on material properties has rarely been investigated. In this work, we investigated porous magnesium samples that were prepared by powder metallurgy using ammonium bicarbonate spacer particles. The effects of the purity of the argon atmosphere and sintering time on the microstructure (SEM, EDX and XRD) and mechanical behaviour (universal loading machine and Vickers hardness tester) of porous magnesium were studied. The porosities of the prepared samples ranged from 24 to 29 vol.% depending on the sintering conditions. The purity of atmosphere played a significant role when the sintering time exceeded 6 h. Under a gettered argon atmosphere, a prolonged sintering time enhanced diffusion connections between magnesium particles and improved the mechanical properties of the samples, whereas under a technical argon atmosphere, oxidation at the particle surfaces caused deterioration in the mechanical properties of the samples. These results suggest that a refined atmosphere is required to improve the mechanical properties of porous magnesium. © 2013 Elsevier B.V. All rights reserved.
1. Introduction Magnesium and magnesium alloys have recently been studied for use in many applications, such as in the automotive and aerospace industries because of their low densities and good mechanical properties [1]. Many successful studies have been performed on biocompatible and biodegradable magnesium-based materials, which are considered suitable materials for orthopaedic applications, such as for nails, screws, splints, etc. [2–6]. For some applications, porous implants, called scaffolds, are required because they possess mechanical properties, such as the modulus of elasticity, that are relatively similar to those found in natural bone tissue [7–14]. The mechanical biocompatibility is important to aid in the remodelling of new tissue [7–14]. Moreover, an interconnected porous structure allows the transport of body fluids to damaged or wounded tissue and supports the incorporation of new tissue in the implant [7,15,16]. Therefore, porous magnesium materials and their preparation methods have been extensively studied in recent years [7–14]. Many methods have been developed for fabricating porous metallic materials [17]. However, because biomaterials should contain interconnected pores [12] and should not be contaminated with harmful impurities [4], only a few of these methods are used for the production of biomaterials. In the available literature, there are five “non-machining” ⁎ Corresponding author. Tel.: +420 220444055. E-mail address: [email protected] (J. Čapek). 0928-4931/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msec.2013.10.014
approaches that have been reported for the fabrication of porous magnesium materials: (1) injection of an inert gas into a melt [4], (2) directional solidification of the metal–gas eutectic (the GASAR process) [9], (3) plaster casting [4], (4) negative salt pattern moulding [12] and (5) powder metallurgical techniques [11,13,14]. However, the first two methods mentioned above do not necessarily produce open-cell structures, and the following two methods may contaminate or corrode the final product during pattern removal [18]. Suitable modifications to powder metallurgy (PM), for example, using space-holder particles, allow the fabrication of materials with interconnected pores. This modification consists of preparing a green compact that contains a powdered mixture of magnesium and a space-holder material, which is later removed by leaching or thermal decomposition. Subsequently, the porous green body is sintered at high temperatures [8,10,13,14]. In principle, any solid matter may be used as the space-holder material; however, in practice, this selection is limited because the spacer material has to be removed without contaminating the product. In the literature, urea and ammonium bicarbonate have been successfully used as spacer materials in the preparation of PM porous magnesium [8,11,13,14,19]. Hao et al. [8] removed the urea by leaching the material in a NaOH-water solution, whereas other authors removed the spaceholder particles by thermal decomposition [10,11,13,14]. The majority of these authors used urea as the spacer material, even though its complete decomposition occurs at temperatures above the melting point of magnesium. Because urea only partially decomposes at lower
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Fig. 1. SEM micrographs (SE detector) of sample cross sections prepared under the following sintering conditions: (a) technical argon for 6 h, (b) technical argon for 24 h, (c) gettered argon for 6 h and (d) gettered argon for 24 h.
temperatures, contamination may occur [20]. For example, Zhuang et al. [13] found traces of carbon and oxygen on pore walls, which they attributed to contamination by urea residues. Ammonium bicarbonate, which decomposes at significantly lower temperatures (36 °C–60 °C), appears to be a more suitable spacer material for the preparation of PM porous magnesium. In addition, leaching the spacer material has the disadvantage of possibly hermetically enclosing some of the spacer particles in the magnesium matrix, which prevents their dissolution. Moreover, water-based dissolving agents are corrosive to magnesium [21]. The mechanical properties of PM porous magnesium are influenced not only by the total porosity but also by the pore size, distribution and shape and the connection between magnesium particles [8,10,13]. These structural characteristics can be adjusted by selecting the optimal sintering time and temperature, compacting pressure and size, shape and volume ratio of the starting material powders [8,11,13,19,22]. Grain coarsening within each magnesium particle, which is expected during the sintering process, may also affect the mechanical properties of the final material [23,24]. In our previous work [19], we showed that porous magnesium can, in principle, be successfully prepared by PM using ammonium bicarbonate as the space-holder material. Although several works on the preparation of porous magnesium have been reported, to the best of our knowledge, no systematic study on the influence of the processing
parameters, namely, the sintering time and atmosphere purity, on the microstructure and mechanical properties of porous magnesium is available. Therefore, our study focused on sintering kinetics and the effect of the sintering atmosphere on the important characteristics of PM porous magnesium.
Table 1 The influence of sintering conditions on sample porosity (in vol.%). Atmosphere/sintering time
0h
3h
6h
12 h
24 h
Technical argon Gettered argon
29 ± 2 29 ± 2
28 ± 1 28 ± 1
28 ± 2 27 ± 2
29 ± 3 25 ± 3
31 ± 4 24 ± 2
Fig. 2. XRD patterns of samples sintered for 6 h.
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compacts were annealed for 4 h at 130 °C in a muffle furnace in air. During this step, the decomposition of ammonium bicarbonate and the evaporation of hexane occurred. Afterwards, sintering was performed at 550 °C in a tube furnace. The green compacts were sintered for 0, 3, 6, 12 and 24 h. The sintering process was performed under two types of flowing atmosphere at a flow rate of 0.1 l/min: (1) argon with technical purity (99.996 vol.%) and (2) argon purified by a 55-mm-thick layer of Mg chips (250–500 μm in size), which was placed around the sintered material and acted as a getter. After sintering, the average material porosity was determined according to Eq. (1) [8,13]: P ¼ 1–ρ=ρMg 100%;
Fig. 3. XRD patterns of samples sintered for 24 h.
2. Experimental A purchased Mg powder (purity of 99.6 wt.%, mesh −100 + 200, Alfa Aesar) and NH4HCO3 powder (p.a. purity, 250–500 μm) were used as starting materials. An Mg/NH4HCO3 volume ratio of 90:10 was used because, according to our previous work [19], this ratio imparts a good combination of strength and corrosion resistance to the resulting material. Materials prepared using this ratio have been shown to possess better mechanical properties that are more comparable to natural bone tissue than porous biomaterials, which have recently been used for medical applications [19]. To avoid segregation, the powders were intensively blended with 30 vol.% hexane for 30 min. Subsequently, pre-weighed mixtures were pressed into cylindrical green compacts (10 mm in diameter and 30 mm in length) at a pressure of 265 MPa using a LabTest 5.250SP1-VM universal loading machine. The green compacts were then subjected to a two-step procedure. First, the
ð1Þ
where P is the porosity, ρ is the density of the porous sample (calculated from the dimensions and weight) and ρMg is the density of pure magnesium (ρMg = 1738 kg/m3). This approach for determining porosity was used because it describes the volumetric 3D porosity more precisely than analysing 2D micrographs. The material mechanical properties were characterised by flexural, compression and Vickers hardness tests that were performed at room temperature. Samples 26 mm and 15 mm in length were used for flexural and compression testing, respectively. Five samples of each series were used for the flexural tests, and three samples of each series were used for the compression tests. The average hardness was calculated from ten values. Standard deviations for all tests were calculated and are shown as error bars in the figure plots. The deformation rates during the flexural and compression tests were 0.5 mm/min and 1 mm/min, respectively. A LabTest 5.250SP1-VM universal loading machine was used for these tests. After flexural testing, the fracture surfaces were observed by a TESCAN VEGA-3 LMU scanning electron microscope (SEM). Metallographic cross sections were also prepared and examined by SEM. Afterwards, the metallographic cross sections were etched (2 g of picric acid, 10 ml of 99% acetic acid, 10 ml of water and 70 ml of ethanol), and the microstructures of the powder particles were observed using an Olympus PME 3 light metallographic microscope. The phase composition and elemental distributions were examined by X-ray diffraction using a PANalytical X'Pert PRO X-ray diffractometer equipped with a Cu anode (XRD) and a TESCAN VEGA-3 LMU SEM equipped with an Oxford Instruments INCA 350 EDX analyser (SEM-EDX). The Vickers hardness was measured for each series with a load of 3 kg (HV 3).
Fig. 4. SEM micrographs (BSE detector) and elemental maps determined by EDX of samples sintered 24 h under (a) technical argon atmosphere and (b) gettered argon atmosphere. Scale bar 50 μm.
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Fig. 5. Microstructure of powder particles during the material fabrication process: (a) the initial powder, (b) compacted powder after annealing at 130 °C for 4 h, (c) compacted powder after annealing at 130 °C for 4 h and sintering at 550 °C for 3 h under technical argon and (d) compacted powder after annealing at 130 °C for 4 h and sintering at 550 °C for 24 h under technical argon.
3. Results and discussion 3.1. Microstructure and porosity The microstructures of samples sintered under two types of atmospheres for different periods of time are shown in Fig. 1. In the magnesium matrix, two types of pores were observed: (1) pores approximately 250–500 μm in size that were formed from the decomposition of spacer particles (“Type I” in Fig. 1) and (2) smaller pores that most likely originated from imperfect compaction during pressing and the expansion of trapped gas during sintering (“Type II” in Fig. 1). No significant difference in microstructure was observed for
Fig. 6. Flexural stress–strain curves of selected samples.
samples sintered for 6 h under the different atmospheres. A prolonged sintering time caused annihilation of “Type II” pores under the gettered atmosphere, whereas under the technical argon atmosphere, the amount of these pores did not significantly change. The sample porosity was measured, and the effects of the sintering conditions on porosity are shown in Table 1. This table indicates that samples sintered under the gettered argon atmosphere became more compact with longer sintering times, i.e., contained less pores. In contrast, the technical argon atmosphere produced materials whose porosities appear almost the same.
Fig. 7. Ultimate flexural strength of prepared samples versus sintering time.
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Fig. 8. Fracture surfaces of samples after flexural testing under the following sintering conditions: (a) technical argon for 6 h - overview, (b) technical argon for 6 h - detail, (c) technical argon for 24 h and (d) gettered argon for 24 h. The white arrow in Fig. 8c denotes oxide particles. In Fig. 8d, the arrow indicates traces of plastic deformation.
This phenomenon suggests that sintering under the gettered argon atmosphere enhanced the diffusion connection between the magnesium particles, which led to annihilation of some “Type II” pores (Fig. 1). In contrast, the use of technical argon caused a slight increase in total porosity. As shown below, porosity is directly related to the presence of oxides in the material.
The XRD patterns of the samples sintered under various conditions are shown in Figs. 2 and 3. The XRD patterns indicate that there is no significant difference in the amount of magnesium oxide in the samples sintered for 6 h, regardless of the atmosphere type (Fig. 2). In this figure, the peaks that are assigned to MgO are similar for both samples sintered for 6 h. However,
Fig. 9. Compressive stress–strain curves of selected samples.
Fig. 10. The compressive yield strength versus sintering time of prepared samples.
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samples that were sintered for 24 h in technical argon (Fig. 3) contained significantly higher amounts of MgO than those sintered in gettered argon. No traces of ammonium bicarbonate or the reaction products of magnesium and ammonium bicarbonate were found by the XRD analysis. To determine the difference between the samples sintered for 24 h, EDX analysis was performed. The results are shown in Fig. 4. The X-ray elemental maps show the presence of Mg particles surrounded by oxide layers. Comparing Fig. 4a and b reveals that more extensive oxidation occurs at the grain boundaries during sintering under the technical argon atmosphere than under the gettered atmosphere. This difference, which is observed after sintering for more than 6 h, is in good agreement with the results acquired from XRD analysis (Fig. 3). Traces of oxygen and water are present in technical argon and cause surface oxidation, which weakens the diffusion connection between the magnesium particles. Thus, annihilation of the smaller “Type II” pores is slowed during longer sintering periods (Table 1). Moreover, oxides on the sample and particle surface may spall off and increase the measured porosity. In contrast, a lower extent of oxidation was observed under the gettered atmosphere, which supports the annihilation of some “Type II” pores because of the diffusive processes of the magnesium particles, as shown in Table 1. The presence of nitrogen and carbon was also investigated as contamination by either ammonium bicarbonate or its decomposition products. No traces of these elements were found; therefore, their distribution is not shown in the XRD elemental maps. Grain coarsening within magnesium particles during sintering was also investigated because it may have an impact on sample properties. Fig. 5 shows detailed views of the microstructure of individual powder particles. The initial powder (Fig. 5a) possesses a relatively fine microstructure (grain size up to 13 μm). After compaction and annealing at 130 °C for 4 h (Fig. 5b), the grain size slightly enlarged; however, the microstructure remained fine (approximate grain size of 15 μm). The grain size significantly increased after sintering at 550 °C for 3 h in technical argon (approximate grain size of 45 μm) (Fig. 5c). Sintering for longer periods did not cause any further significant grain growth. The average grain size after sintering for 24 h at 550 °C was approximately 48 μm (Fig. 5d). The type of atmosphere did not influence grain growth during sintering.
Fig. 11. The ultimate compressive strength versus sintering time of prepared samples.
A study of the fracture surfaces helped to explain this behaviour. The fracture surfaces of samples after flexural testing are illustrated in Fig. 8. For all the samples, the fracture surfaces were similar (Fig. 8a); however, slight differences were observed at higher magnifications. Generally, the fracture surfaces were brittle, but traces of plastic deformation were observed, especially in samples sintered under gettered argon for longer time periods (12 and 24 h) (Fig. 8d). The fracture surfaces were nearly identical and oxide free after sintering for up to 6 h, independent of the atmosphere (Fig. 8b). After longer sintering times, some differences were observed between the fracture surfaces of samples prepared under the different atmospheres. Samples sintered in gettered argon had nearly oxide-free fracture surfaces with features of plastic deformation (Fig. 8d), whereas a large amount of oxide was present on the fracture surfaces of samples sintered in technical argon (Fig. 8c). This observation confirms the XRD and EDX results that were mentioned above (Figs. 2–4).
3.2. Mechanical properties The flexural stress–strain curves of selected samples are shown in Fig. 6. Samples sintered in technical argon had a significantly lower modulus of elasticity than samples sintered under the gettered atmosphere. This result can be explained by the increased amount of oxygen at the interface between powder particles, which results in decreased cohesiveness between the particles. Oxides also impact flexural strength, which will be shown and discussed below. The flexural stress–strain curves of all the materials contained only elastic regions; no regions of macroscopic plastic deformation were observed (Fig. 6). The ultimate flexural strength (UFS) as a function of sintering time is plotted in Fig. 7. The ultimate flexural strength increased with longer sintering times, up to 6 h, independent of the atmosphere. The samples attained approximately the same UFS, 12 MPa, after 6 h of sintering under both atmospheres. For longer sintering times under the gettered argon atmosphere, the UFS increased to approximately 15 MPa. The difference between samples sintered for 12 and 24 h under the gettered argon atmosphere is negligible. In contrast, prolonging the sintering period to more than 6 h under technical argon had a completely opposite effect. The UFS rapidly decreased to 5 MPa after 12 h of sintering and continued to decrease to approximately 3 MPa after 24 h of sintering.
Fig. 12. The HV 3 versus sintering time of prepared samples.
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Table 2 Mechanical properties of porous biomaterials. Material
Porosity (vol.%)
Pore size (μm)
UFS (MPa)
CYS (MPa)
UCS (MPa)
Reference
Natural bone Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Mg Porous Ti Porous hydroxyapatite Porous hydroxyapatite Porous composite (poly-L-lactide + 20–50 wt.% of bioglass) Porous polycaprolactone Porous polycaprolactone Porous polycaprolactone Porous polylactide-co-glycolide Porous composite of polylactide-co-glycolide and 45S5 bioglass
— 29–31 23–38 14–44 36–55 52–70 50 28 35–55 78 50–77 — 77–88 48–77 37–55 55–56 31 43
— 250–500 250–500 250–500 200–400 ~1250 200–500 170 100–400 200–500 200–400 366–444 ~100 — — — 116 89
2–150 3–15 4–17 2–5 14–27 — — — — — 2–7 — 1–4 — —
— 13–53 — — — — — — — — — — — 2–3 — 2–3 — —
2–180 20–70 — — 15–31 4–14 2 24 12–17 35 1–17 30 ~0.4 — 2–3
[13] This study [19] [22] [13] [8] [14] [7] [11] [14] [13] [25] [13] [25] [25] [26] [27] [27]
Fig. 9 shows the compressive stress–strain curves of selected samples. From the figure, the trends of the compressive and flexural behaviour of the material are similar, with the exception of the effect of atmospheric condition on the compressive modulus of elasticity, which is not as strong. The compressive yield strength (CYS) and ultimate compressive strength (UCS) as functions of the sintering time are plotted in Figs. 10 and 11. Under the gettered argon atmosphere, both the yield strength (CYS) and the ultimate strength (UCS) increased with longer sintering times. The maximum CYS and UCS values after sintering for 24 h are 53 and 69 MPa, respectively. These high strength values result from a low amount of oxide, decreased porosity and good contact between the particles in the samples (Figs. 2–4, Table 1). In contrast, for samples sintered under technical argon, the maximum CYS and UCS were attained after 6 h of sintering. Prolonging the sintering period under this atmosphere led to a significant decrease in compressive properties. This finding can be attributed to the extended oxidation that occurs during sintering, which weakens the diffusion bonds between particles. An interesting observation in Figs. 10 and 11 is that the CYS of samples sintered in gettered argon increased almost linearly with sintering time, whereas the UCS increased for up to 6 h of sintering and then increased significantly more slowly. These results indicate that the “Type II” pores, which decrease in quantity during long sintering periods (Table 1), influence the yield strength rather than the ultimate strength. The ultimate compressive strength is probably more influenced by the larger “Type I” pores, whose amount is almost constant for all samples. The reason may be that the large pores concentrate the stress in their vicinity and initiate the crack nucleation and subsequent collapse of the porous structure. On the other hand, the small “Type II” pores concentrate the stress to a smaller extent and plastic deformation takes place in their surroundings. Therefore, their partial vanishing during longer sintering times seems to increase the stress needed for the plastic strain to begin. Fig. 12 reveals the effect of sintering time on sample hardness. Samples sintered in gettered argon possessed a higher hardness than those sintered in technical argon. The effect of sintering time on hardness was small for samples sintered in technical argon. However, prolonging the sintering period in the gettered argon led to an increase in hardness, with a peak hardness occurring in samples sintered for 12 h. Sintering up to 24 h slightly decreased sample hardness. This result may be caused by oxidation beginning at the grain boundaries and decreases the cohesiveness of the magnesium particles. A comparison of our results with previously reported results is difficult because of differences in the porosities and pore sizes of the studied materials. Nevertheless, we conducted a small review that summarises the mechanical properties of natural bone tissue and porous materials
— —
0.5 0.4
that are considered suitable biomaterials. The results of this research are summarised in Table 2. From this table, it is clear that porous magnesium-based metallic materials exhibit mechanical properties similar to those of natural bone tissue. However, non-metallic porous materials (especially polymer-based materials) have weaker mechanical properties compared to natural bone tissue. It is also important to note that the materials prepared in our study have higher UFS and UCS values than porous magnesium materials that have been prepared in the majority of other works and non-metallic porous materials. Moreover, the materials prepared in this work possess mechanical properties that are comparable to natural bone tissue. 4. Conclusions Magnesium samples with porosities of 29–31 vol.% were prepared by powder metallurgy under different sintering conditions. When compared to non-metallic biomaterials, the samples exhibited enhanced mechanical properties that were similar to human bone tissue. After sintering for up to 6 h, no significant effect of atmosphere purity on sample microstructure or mechanical behaviour was observed. Under the gettered argon atmosphere, longer sintering times decreased porosity but enhanced the mechanical properties of the sample. Under the technical argon atmosphere, the opposite trend was observed, which can be attributed to oxidation of the powder surface. These results suggest that when sintering is longer than 6 h, the purity of the argon atmosphere plays an important role in determining the mechanical properties of PM magnesium. Acknowledgements The authors would like to thank the Czech Science Foundation (project no. P108/12/G043) for supporting this research. References [1] F.H. Froes, Mater. Sci. Eng. A 184 (1994) 119–133. [2] N. Li, Y.F. Zheng, J. Mater. Sci. Technol. 29 (2013) 489–502. [3] H. Waizy, J.M. Seitz, J. Reifenrath, A. Weizbauer, F.W. Bach, A. Meyer-Lindenberg, B. Denkena, H. Windhagen, J. Mater. Sci. 48 (2013) 39–50. [4] M.P. Staiger, A.M. Pietak, J. Huadmai, G. Dias, Biomaterials 27 (2006) 1728–1734. [5] S. Virtanen, Mater. Sci. Eng. B 176 (2011) 1600–1608. [6] M. Hradilova, J. Kubasek, P. Lejcek, L. Tanger, Metal 2011: 20th Anniversary International Conference on Metallurgy and Materials, 2011, pp. 920–924, (full text available on http://www.metal2013.com/files/proceedings/metal_11/lists/papers/978. pdf, 31.7.2013). [7] X.N. Gu, W.R. Zhou, Y.F. Zheng, Y. Liu, Y.X. Li, Mater. Lett. 64 (2010) 1871–1874. [8] G. Hao, F. Han, W. Li, J. Porous. Mater. 16 (2009) 251–256.
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