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EPD combined method
Materials Letters 65 (2011) 926–928
Contents lists available at ScienceDirect
Materials Letters 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 a t l e t
Innovative fabrication of ZrO2–HAp–TiO2 nano/micro-structured composites through MAO/EPD combined method F. Samanipour a, M.R. Bayati a,⁎, F. Golestani-Fard a,b, H.R. Zargar c, A.R. Mirhabibi a, V. Shoaei-Rad a, S. Abbasi a a b c
School of Metallurgy and Materials Engineering, Iran University of Science and Technology, P.O. Box: 16845-161, Tehran, Iran Center of Excellence for Advanced Materials, Iran University of Science and Technology, P.O. Box: 16845-195, Tehran, Iran Department of Metals and Materials Engineering, University of British Columbia, Vancouver, BC V6T 1Z4, Canada
a r t i c l e
i n f o
Article history: Received 21 October 2010 Accepted 15 November 2010 Available online 20 November 2010 Keywords: Composite Coating Porous Hydroxyapatite Zirconia Titania
a b s t r a c t ZrO2–HAp–TiO2 porous layers were fabricated by micro arc oxidation/electrophoretic deposition hybrid processes in the electrolytes consisted of β-glycerophosphate, calcium acetate, sodium phosphate, and yttriastabilized zirconia within short times and only in one step. The layers had a porous structure with a pores size of 50–750 nm. They contained anatase, hydroxyapatite, monoclinic ZrO2, tetragonal ZrO2, and some minor phases with a varying fraction depending on time. Increasing the time, the tetragonal ZrO2 transformed to monoclinic ZrO2. Moreover, hydroxyapatite relative content decreased with time. The nanosized zirconia particles (d = 20–40 nm) were dispersed on the surface and in the layers depth. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Titanium and its alloys are widely used as surgical implants in orthopedics and dentistry due to their excellent mechanical properties and bio-affinity. To improve the osseointegration of the implants, the Ti surface has been modified by bioactive coatings, chemical etching, and sand blasting. Coatings with hydroxyapatite and titania have attracted particular attention over the last decade [1]. Even though hydroxyapatite has been introduced as a substitute material for damaged tissues, its wide and successful application has been limited by its poor mechanical properties especially under loadbearing conditions [2,3]. Therefore, it is essential to enhance its mechanical properties for long-term performance. An approach to enhance the mechanical properties of HAp is reinforcing it with other ceramics [4]. Zirconia has excellent technological properties and is a proper candidate to be combined with HAp [5,6]. Combining hydroxyapatite and zirconia, favorable mechanical properties [7], good corrosion resistance, biocompatibility, and bioactivity are achieved [8,9]. Meanwhile, the bonding strength of HA increases with addition of Y-ZrO2 [10]. Electrophoretic deposition (EPD) is a simple and inexpensive technique making it possible to synthesize a wide range of the ceramic coatings with varying morphologies and chemical compositions [11,12]. Micro arc oxidation (MAO) is a simple
and promising approach for fabrication of different categories of oxide layers especially combined systems [13–18]. In this report, ZrO2–HAp–TiO2 combined layers were derived by MAO/EPD process for the first time and the growth time on physical and chemical features of the layers was investigated. 2. Experimental Commercially pure grade 2 (CP-II) titanium pieces, with dimensions of 3 cm× 1.5 cm× 0.5 mm, were utilized as substrate. Details about the employed experimental setup and substrates surface cleaning procedure can be found in our previous works [13–18]. The electrolytes were prepared by dissolution of the β-glycerophosphate disodium (C3H7Na2O6P, Merck) and calcium acetate (Ca(CH3COO)2, Panareac) salts with concentrations of 1 and 5 g l−1 in 75 vol.% DI-water and 25 vol.% ethanol. After complete dissolution, yttria-stabilized zirconia (8 mol%, Tosoh) powder, with concentrations of 1 g l−1, was added to the suspension by jar-milling for 48 h at room temperature. The pH of the suspensions was adjusted ~11.5 through the addition of sodium phosphate (Na3PO4.12H2O, Panareac). Electrolyte temperature was fixed at 70 ± 3 °C employing a water circulating system. The MAO/EPD treatment was performed under the voltage of 350 V for various times. 3. Results and discussion
⁎ Corresponding author. Tel./fax: +98 21 77240291. E-mail address: [email protected] (M.R. Bayati). URL: http://rrg.iust.ac.ir (M.R. Bayati). 0167-577X/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2010.11.039
SEM surface morphology of the layers grown at different times is depicted in Fig. 1. As is seen, the pores size increases with the treatment time. When a structural pore forms by an electrical spark, it
F. Samanipour et al. / Materials Letters 65 (2011) 926–928
927
Fig. 1. SEM surface morphology of the composite layers fabricated for different times.
Fig. 2. XRD patterns of the composite layers grown for different times. (Ti: Titanium, A: Anatase, H: Hydroxyapatite, mZ: Monoclinic ZrO2, tZ: Tetragonal ZrO2, Z: ZrO, ZT: ZrTiO4, CT: CaTiO3, P: α-TCP).
is more susceptible for next electron avalanches, because these areas have lower breakdown voltages than other areas. Sequence of the electrical sparks at one point makes the pores larger. It is worthy to be noted that these electric sparks are responsible for formation of the structural pores. It is also clear that the concentration of the zirconia particles on the surface increases with the time. Surface of the zirconia particles contains broken bonds and can be simply charged in aqueous solutions by changing the pH. OH− ions can be attached to the particles surface resulting in formation of negatively charged zirconia particles. These particles move toward the anode due to the electrical field between anode and cathode. Fabricating layers for longer times, the zirconia particles can further accumulate on the vicinity of the anode. Fig. 2 illustrates the normalized XRD patterns of the layers grown for different times. Because more zirconia particles move toward the anode and accumulate on the vicinity of the layer under longer growth times, the ZrO2 XRD peaks intensify with the time. Increasing the anode temperature during the MAO treatment results in phase transformation of the metastable stabilized zirconia to its stable monoclinic structure. It has been reported that zirconia-based materials exhibit exceptional toughness due to the martensitic transformation of between tetragonal and monoclinic ZrO2 [19]. Formation of some minor phases is also evident. The temperature inside the pores elevates to ~104 K due to the high energy electron avalanches. This temperature is high enough to melt the substrate and
the layer. The molten system solves the ZrO2 particles and other anions to form these minor phases. It is reminded that negatively charged species are drawn inside the pores because of the strong electric field between the anode and the cathode which is about 106 V m−1 [20,21]. Fig. 3 presents the XPS survey spectra of the sample grown for 6 min where existence of Ti, Zr, P, Ca, O, and C is evident. Fig. 4a
Fig. 3. XPS survey spectrum of the layer grown for 6 min.
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F. Samanipour et al. / Materials Letters 65 (2011) 926–928
Fig. 4. XPS core level binding energies in the layer grown for 6 min.
depicts the Ti(2p3/2) core level binding energy. The peak fitting was performed with 2 components where the dominant peak (Peak B) is assigned to the Ti4+ in Ti–O binding. Since the TiO2 surface loses its oxygen at high temperatures, a small amount of Ti3+ (peak A) corresponding to the Ti2O3 compound was also detected. The Zr(3d) core level binding energy is shown in Fig. 4b. The peaks A and B are consistent with those of metallic Zro(3d5/2) and Zro(3d3/2), respectively. The peaks C and D, at the binding energies of 180.5 and 181.5 eV, represent existence of Zr2+ and formation of ZrO and probably some other non-stoichiometric zirconium oxides (ZrxOy). Finally, formation of ZrO2 compound is confirmed by the peaks E and F located at the binding energies of 182.4 and 183.5 eV [22,23]. The P (2p) core level binding energy is exhibited in Fig. 4c which is deconvoluted to the peaks A and B, with binding energies of 132.7 and 134.2 eV. These peaks, which are corresponding to the 2p3/2 and 2p1/2 core levels, prove the existence of phosphor in HAp structure. Fig. 4d shows the XPS spectra of Ca(2p) core level binding energy. The peaks A and B, with binding energies of 347.7 and 351.2 eV, are attributed to the Ca(2p3/2) and Ca(2p1/2) core levels. They are consistent with those of calcium in the HAp structure. Finally, the O(1s) core level binding energy is depicted in Fig. 4e which can be deconvoluted to 5 distinct peaks representing that there are 5 different O-bindings in the layers. The peak A, located at the binding energy of 529.8 eV, is assigned to the oxygen in the oxides lattice namely ZrO2 and TiO2. The peak B, at the binding energy of 530.7 eV, corresponds to the oxygen in nonstoichiometric zirconium oxides [18]. The peak C at the binding energy of 531.85 eV represents the –OH groups on the surface. Oxide free surfaces are always hydrated in the atmosphere. Therefore, the chemisorbed hydroxyl groups are always seen in the XPS patterns obtained from the oxide films. The peaks C and D, located at the binding energies of 533.1 and 534.4 eV, reveal existence of oxygen as O− and H2O with descending binding energy, respectively. Since the layers were porous and fabricated in aqueous solutions, water may be trapped in the pores. The layers were enshrouded by an oxygen plasma during the MAO treatment which causes of attaching oxygen
radicals to the surface. It is to be noted that all results were interpreted according to the C(1s) core level binding energy at 285.0 eV.
4. Conclusion ZrO2–HAp–TiO2 layers were fabricated by MAO/EPD hybrid technique. The layers had a porous morphology where the pores size increased with the time Anatase, hydroxyapatite, monoclinic ZrO2, and tetragonal ZrO2 were the main phases of the layers. The nanosized zirconia particles (d = 30–60 nm) were loaded on the surface and in the depth of the layers.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
Kim HW, Koh YH, Li LH, Lee S, Kim HE. Biomaterials 2004;25:2533. Hench LL. J Am Ceram Soc 1991;74:1487. Berger S, Müller E, Schnabelrauch M. Mater Lett 2009;63:2714. Ducheyne P, Marcolongo M, Schepers E. In: Hench LL, Wilson J, editors. An Introduction to Bioceramic. Singapore: World Scientific Publishing Co; 1993. p. 281. Jentoft FC, Fischer A, Weinberg G, Wild U, Schlgl R. Stud Surf Sci Catal 2000;130:209. Morks MF. Mater Lett 2010;64:1968. Chang BY, Chang E. J Mater Sci 2002;13:589. Ferraris M, Verne HE, Appendino P, Moisescu C, Krajewski A, Ravaglioli A. Biomaterials 2000;21:765. Hsu HC, Wu SC, Yang CH, Wen-Fu Ho. J Mater Sci Mater Med 2009;20:615. Chang E, Chang WJ, Wang BC, Yang CY. J Mater Sci 1997;8:193. Simchi A, Pishbin F, Boccaccini AR. Mater Lett 2009;63:2253. Riahifar R, Marzbanrad E, Dehkordi BR, Zamani C. Mater Lett 2010;64:559. Bayati MR, Moshfegh AZ, Golestani-Fard F. Electrochim Acta 2010;55:3093. Bayati MR, Golestani-Fard F, Moshfegh AZ. Appl Catal A Gen 2010;382:322. Bayati MR, Moshfegh AZ, Golestani-Fard F. Appl Surf Sci 2010;256:2903. Bayati MR, Moshfegh AZ, Golestani-Fard F. Appl Catal A Gen 2010, doi:10.1016/j. apcata.2010.09.003. Bayati MR, Moshfegh AZ, Golestani-Fard F. Mater Lett 2010;64:2215. Bayati MR, Molaei R, Zargar HR, Kajbafvala A. Mater Lett 2010;64:2498. Rapacz-Kmita A, Slosarczyk A, Paszkiewicz Z. J Eur Ceram Soc 2006;26:1481. Yerokhin AL, Nie X, Leyland A, Matthews A, Dowey SJ. Surf Coat Tech 1999;122:73. Yerokhin AL, Lyubimov VV, Ashitkof RV. Ceram Int 1998;24:1. Park J, Heo JK, Kang YC. Bull Korean Chem Soc 2010;31:397. Wang S, Liu C, Shan F. Acta Metall Sin Engl Lett 2009;22:161.
Contents lists available at ScienceDirect
Materials Letters 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 a t l e t
Innovative fabrication of ZrO2–HAp–TiO2 nano/micro-structured composites through MAO/EPD combined method F. Samanipour a, M.R. Bayati a,⁎, F. Golestani-Fard a,b, H.R. Zargar c, A.R. Mirhabibi a, V. Shoaei-Rad a, S. Abbasi a a b c
School of Metallurgy and Materials Engineering, Iran University of Science and Technology, P.O. Box: 16845-161, Tehran, Iran Center of Excellence for Advanced Materials, Iran University of Science and Technology, P.O. Box: 16845-195, Tehran, Iran Department of Metals and Materials Engineering, University of British Columbia, Vancouver, BC V6T 1Z4, Canada
a r t i c l e
i n f o
Article history: Received 21 October 2010 Accepted 15 November 2010 Available online 20 November 2010 Keywords: Composite Coating Porous Hydroxyapatite Zirconia Titania
a b s t r a c t ZrO2–HAp–TiO2 porous layers were fabricated by micro arc oxidation/electrophoretic deposition hybrid processes in the electrolytes consisted of β-glycerophosphate, calcium acetate, sodium phosphate, and yttriastabilized zirconia within short times and only in one step. The layers had a porous structure with a pores size of 50–750 nm. They contained anatase, hydroxyapatite, monoclinic ZrO2, tetragonal ZrO2, and some minor phases with a varying fraction depending on time. Increasing the time, the tetragonal ZrO2 transformed to monoclinic ZrO2. Moreover, hydroxyapatite relative content decreased with time. The nanosized zirconia particles (d = 20–40 nm) were dispersed on the surface and in the layers depth. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Titanium and its alloys are widely used as surgical implants in orthopedics and dentistry due to their excellent mechanical properties and bio-affinity. To improve the osseointegration of the implants, the Ti surface has been modified by bioactive coatings, chemical etching, and sand blasting. Coatings with hydroxyapatite and titania have attracted particular attention over the last decade [1]. Even though hydroxyapatite has been introduced as a substitute material for damaged tissues, its wide and successful application has been limited by its poor mechanical properties especially under loadbearing conditions [2,3]. Therefore, it is essential to enhance its mechanical properties for long-term performance. An approach to enhance the mechanical properties of HAp is reinforcing it with other ceramics [4]. Zirconia has excellent technological properties and is a proper candidate to be combined with HAp [5,6]. Combining hydroxyapatite and zirconia, favorable mechanical properties [7], good corrosion resistance, biocompatibility, and bioactivity are achieved [8,9]. Meanwhile, the bonding strength of HA increases with addition of Y-ZrO2 [10]. Electrophoretic deposition (EPD) is a simple and inexpensive technique making it possible to synthesize a wide range of the ceramic coatings with varying morphologies and chemical compositions [11,12]. Micro arc oxidation (MAO) is a simple
and promising approach for fabrication of different categories of oxide layers especially combined systems [13–18]. In this report, ZrO2–HAp–TiO2 combined layers were derived by MAO/EPD process for the first time and the growth time on physical and chemical features of the layers was investigated. 2. Experimental Commercially pure grade 2 (CP-II) titanium pieces, with dimensions of 3 cm× 1.5 cm× 0.5 mm, were utilized as substrate. Details about the employed experimental setup and substrates surface cleaning procedure can be found in our previous works [13–18]. The electrolytes were prepared by dissolution of the β-glycerophosphate disodium (C3H7Na2O6P, Merck) and calcium acetate (Ca(CH3COO)2, Panareac) salts with concentrations of 1 and 5 g l−1 in 75 vol.% DI-water and 25 vol.% ethanol. After complete dissolution, yttria-stabilized zirconia (8 mol%, Tosoh) powder, with concentrations of 1 g l−1, was added to the suspension by jar-milling for 48 h at room temperature. The pH of the suspensions was adjusted ~11.5 through the addition of sodium phosphate (Na3PO4.12H2O, Panareac). Electrolyte temperature was fixed at 70 ± 3 °C employing a water circulating system. The MAO/EPD treatment was performed under the voltage of 350 V for various times. 3. Results and discussion
⁎ Corresponding author. Tel./fax: +98 21 77240291. E-mail address: [email protected] (M.R. Bayati). URL: http://rrg.iust.ac.ir (M.R. Bayati). 0167-577X/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2010.11.039
SEM surface morphology of the layers grown at different times is depicted in Fig. 1. As is seen, the pores size increases with the treatment time. When a structural pore forms by an electrical spark, it
F. Samanipour et al. / Materials Letters 65 (2011) 926–928
927
Fig. 1. SEM surface morphology of the composite layers fabricated for different times.
Fig. 2. XRD patterns of the composite layers grown for different times. (Ti: Titanium, A: Anatase, H: Hydroxyapatite, mZ: Monoclinic ZrO2, tZ: Tetragonal ZrO2, Z: ZrO, ZT: ZrTiO4, CT: CaTiO3, P: α-TCP).
is more susceptible for next electron avalanches, because these areas have lower breakdown voltages than other areas. Sequence of the electrical sparks at one point makes the pores larger. It is worthy to be noted that these electric sparks are responsible for formation of the structural pores. It is also clear that the concentration of the zirconia particles on the surface increases with the time. Surface of the zirconia particles contains broken bonds and can be simply charged in aqueous solutions by changing the pH. OH− ions can be attached to the particles surface resulting in formation of negatively charged zirconia particles. These particles move toward the anode due to the electrical field between anode and cathode. Fabricating layers for longer times, the zirconia particles can further accumulate on the vicinity of the anode. Fig. 2 illustrates the normalized XRD patterns of the layers grown for different times. Because more zirconia particles move toward the anode and accumulate on the vicinity of the layer under longer growth times, the ZrO2 XRD peaks intensify with the time. Increasing the anode temperature during the MAO treatment results in phase transformation of the metastable stabilized zirconia to its stable monoclinic structure. It has been reported that zirconia-based materials exhibit exceptional toughness due to the martensitic transformation of between tetragonal and monoclinic ZrO2 [19]. Formation of some minor phases is also evident. The temperature inside the pores elevates to ~104 K due to the high energy electron avalanches. This temperature is high enough to melt the substrate and
the layer. The molten system solves the ZrO2 particles and other anions to form these minor phases. It is reminded that negatively charged species are drawn inside the pores because of the strong electric field between the anode and the cathode which is about 106 V m−1 [20,21]. Fig. 3 presents the XPS survey spectra of the sample grown for 6 min where existence of Ti, Zr, P, Ca, O, and C is evident. Fig. 4a
Fig. 3. XPS survey spectrum of the layer grown for 6 min.
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F. Samanipour et al. / Materials Letters 65 (2011) 926–928
Fig. 4. XPS core level binding energies in the layer grown for 6 min.
depicts the Ti(2p3/2) core level binding energy. The peak fitting was performed with 2 components where the dominant peak (Peak B) is assigned to the Ti4+ in Ti–O binding. Since the TiO2 surface loses its oxygen at high temperatures, a small amount of Ti3+ (peak A) corresponding to the Ti2O3 compound was also detected. The Zr(3d) core level binding energy is shown in Fig. 4b. The peaks A and B are consistent with those of metallic Zro(3d5/2) and Zro(3d3/2), respectively. The peaks C and D, at the binding energies of 180.5 and 181.5 eV, represent existence of Zr2+ and formation of ZrO and probably some other non-stoichiometric zirconium oxides (ZrxOy). Finally, formation of ZrO2 compound is confirmed by the peaks E and F located at the binding energies of 182.4 and 183.5 eV [22,23]. The P (2p) core level binding energy is exhibited in Fig. 4c which is deconvoluted to the peaks A and B, with binding energies of 132.7 and 134.2 eV. These peaks, which are corresponding to the 2p3/2 and 2p1/2 core levels, prove the existence of phosphor in HAp structure. Fig. 4d shows the XPS spectra of Ca(2p) core level binding energy. The peaks A and B, with binding energies of 347.7 and 351.2 eV, are attributed to the Ca(2p3/2) and Ca(2p1/2) core levels. They are consistent with those of calcium in the HAp structure. Finally, the O(1s) core level binding energy is depicted in Fig. 4e which can be deconvoluted to 5 distinct peaks representing that there are 5 different O-bindings in the layers. The peak A, located at the binding energy of 529.8 eV, is assigned to the oxygen in the oxides lattice namely ZrO2 and TiO2. The peak B, at the binding energy of 530.7 eV, corresponds to the oxygen in nonstoichiometric zirconium oxides [18]. The peak C at the binding energy of 531.85 eV represents the –OH groups on the surface. Oxide free surfaces are always hydrated in the atmosphere. Therefore, the chemisorbed hydroxyl groups are always seen in the XPS patterns obtained from the oxide films. The peaks C and D, located at the binding energies of 533.1 and 534.4 eV, reveal existence of oxygen as O− and H2O with descending binding energy, respectively. Since the layers were porous and fabricated in aqueous solutions, water may be trapped in the pores. The layers were enshrouded by an oxygen plasma during the MAO treatment which causes of attaching oxygen
radicals to the surface. It is to be noted that all results were interpreted according to the C(1s) core level binding energy at 285.0 eV.
4. Conclusion ZrO2–HAp–TiO2 layers were fabricated by MAO/EPD hybrid technique. The layers had a porous morphology where the pores size increased with the time Anatase, hydroxyapatite, monoclinic ZrO2, and tetragonal ZrO2 were the main phases of the layers. The nanosized zirconia particles (d = 30–60 nm) were loaded on the surface and in the depth of the layers.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
Kim HW, Koh YH, Li LH, Lee S, Kim HE. Biomaterials 2004;25:2533. Hench LL. J Am Ceram Soc 1991;74:1487. Berger S, Müller E, Schnabelrauch M. Mater Lett 2009;63:2714. Ducheyne P, Marcolongo M, Schepers E. In: Hench LL, Wilson J, editors. An Introduction to Bioceramic. Singapore: World Scientific Publishing Co; 1993. p. 281. Jentoft FC, Fischer A, Weinberg G, Wild U, Schlgl R. Stud Surf Sci Catal 2000;130:209. Morks MF. Mater Lett 2010;64:1968. Chang BY, Chang E. J Mater Sci 2002;13:589. Ferraris M, Verne HE, Appendino P, Moisescu C, Krajewski A, Ravaglioli A. Biomaterials 2000;21:765. Hsu HC, Wu SC, Yang CH, Wen-Fu Ho. J Mater Sci Mater Med 2009;20:615. Chang E, Chang WJ, Wang BC, Yang CY. J Mater Sci 1997;8:193. Simchi A, Pishbin F, Boccaccini AR. Mater Lett 2009;63:2253. Riahifar R, Marzbanrad E, Dehkordi BR, Zamani C. Mater Lett 2010;64:559. Bayati MR, Moshfegh AZ, Golestani-Fard F. Electrochim Acta 2010;55:3093. Bayati MR, Golestani-Fard F, Moshfegh AZ. Appl Catal A Gen 2010;382:322. Bayati MR, Moshfegh AZ, Golestani-Fard F. Appl Surf Sci 2010;256:2903. Bayati MR, Moshfegh AZ, Golestani-Fard F. Appl Catal A Gen 2010, doi:10.1016/j. apcata.2010.09.003. Bayati MR, Moshfegh AZ, Golestani-Fard F. Mater Lett 2010;64:2215. Bayati MR, Molaei R, Zargar HR, Kajbafvala A. Mater Lett 2010;64:2498. Rapacz-Kmita A, Slosarczyk A, Paszkiewicz Z. J Eur Ceram Soc 2006;26:1481. Yerokhin AL, Nie X, Leyland A, Matthews A, Dowey SJ. Surf Coat Tech 1999;122:73. Yerokhin AL, Lyubimov VV, Ashitkof RV. Ceram Int 1998;24:1. Park J, Heo JK, Kang YC. Bull Korean Chem Soc 2010;31:397. Wang S, Liu C, Shan F. Acta Metall Sin Engl Lett 2009;22:161.