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Fabrication of a porous CuAlMn shape memory alloy by the sintering–dissolution process
Materials Letters 65 (2011) 2735–2738
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
Fabrication of a porous CuAlMn shape memory alloy by the sintering–dissolution process Qingzhou Wang ⁎, Chunxiang Cui, Qian Wang, Najun Yan School of Materials Science and Engineering, Hebei University of Technology, Road No. 1, Dingzigu, Hongqiao District, Tianjin 300130, China
a r t i c l e
i n f o
Article history: Received 14 March 2011 Accepted 21 May 2011 Available online 30 May 2011 Keywords: Porous materials Shape memory materials Powder technology
a b s t r a c t A porous CuAlMn shape memory alloy with adjustable pore characteristics and mechanical properties has been successfully fabricated via the sintering–dissolution process for the first time. Micro amount of fluoride was added to accelerate the sintering of CuAlMn powders covered with the oxide film. The resultant material exhibits uniformly distributed and interconnected macroscopic pores. In order to eliminate possible adverse effects of the microstructures resulted from slow cooling on the properties of the resultant materials, appropriate solution and quench treatments were carried out. The microstructures and quasi-static compressive deformation behaviors of the resultant materials were also examined before and after treatments. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Initiating from early studies in Russian and further developments in China, porous shape memory alloys (SMAs) have been extensively investigated [1]. They have lower density and the possibility of further enhancing energy absorption and damping capabilities. Up to now investigations of porous SMAs are mainly focused on the Ni–Ti alloys [2–5]. But the high cost is its main limitation for engineering applications. As a result, Cu-based SMAs are attracting more and more attention in recent years. In our previous work a porous Cu-based SMA was fabricated via an air pressure infiltration process [6]. However, in the fabrication process there are still strict restrictions to the superheating temperature of Cu-based SMA melt, the choice of dissoluble space-holders as well as the velocity of infiltration process due to the relatively high melting point of Cu-based SMAs (over 1000 °C). Fast solidification of melting metals enhances the level of difficulty in techniques. It is therefore necessary to develop new techniques to offer a wider range of structures and properties while continuing to improve original infiltration technique. Powder metallurgy technique possesses unique advantages in fabrication of porous metals [7,8]. Among the powder metallurgical techniques sintering–dissolution process (SDP) is a promising route for manufacturing open-celled porous metals [8–10]. The most outstanding technical advantage of SDP is much wider range of structure parameters, thus providing more opportunities for resultant porous metals to meet the desired behaviors. Thus, in the present study the SDP is employed to fabricate a porous CuAlMn SMA using NaCl particulates as space-holders. It is interesting to note that ⁎ Corresponding author. Fax: + 86 22 60204125. E-mail address: [email protected] (Q. Wang). 0167-577X/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2011.05.092
using NaCl as space-holders possesses numerous advantages, such as low cost, fast dissolution in water, reduced corrosive attack of metal during dissolution, free of toxicity, in addition to significant widening of structure parameters such as porosity, pore shape, pore size and homogeneity of pore distribution, etc. [11]. 2. Experimental An SDP process was employed to fabricate the porous SMA specimens from the Cu–11.7Al–2.49Mn (wt.%) SMA powders with a purity of more than 99% (shown in Fig. 1a) prepared by an atomization method and domestic NaCl particulates (shown in Fig. 1d). From the figure it can be seen that the CuAlMn powders with an average diameter of about 15 μm are round or elliptic, while NaCl particles are around 0.4 mm in size and cuboidal in shape. Fig. 1b shows the DSC analysis result of the starting CuAlMn powders. The characteristic temperatures As, Af, Ms and Mf of the Cu–11.7Al– 2.49Mn alloy were determined as 183, 216, 156 and 142 °C, respectively. The CuAlMn powders were blended with micro amount of fluoride (the mixture of AlF3 and CaF2, less than 0.08 wt.%) to remove the surface oxidation layer by the formation of gaseous AlOF from the reaction between Al2O3 and AlF3. The reaction can be accelerated further by the presence of liquid phase which appeared in Al–Ca fluoride. The mixture was then mechanically milled for 1.5 h under an atmosphere of argon to get the final well-mixed powders as shown in Fig. 1c. It can be seen that milled powders have irregular shapes and coarse surfaces. These features can improve the coupling between CuAlMn powders and facilitate the pressing process of green compact. Ball milled CuAlMn powders and NaCl particulates with a chosen size were weighed using an electronic balance with an accuracy of 0.01 g. They were blended together in a stirrer at a rotation speed of
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Q. Wang et al. / Materials Letters 65 (2011) 2735–2738
Fig. 1. Raw materials used in the SDP. (a) Starting CuAlMn powders; (b) DSC result of starting CuAlMn powders; (c) ball milled mixture of CuAlMn and fluoride powders; (d) NaCl particulates.
80 rpm for 10 min. The volume percentages of NaCl were varied from 50% to 70% with a gap of 5%. Well mixed powders with a small amount of ethanol were uniaxially pressed into a steel die with an inner diameter of 20 mm under pressures ranged from 400 to 450 MPa.
Relatively high pressures were adopted in the present study to obtain high density green compacts. However, excessive high pressure should be avoided due to the existence of brittle NaCl particulates that are liable to be crushed under too high pressure.
Fig. 2. Typical structure of the porous CuAlMn shape memory alloys. (a), (c) and (d) cross-section of the specimens perpendicular to the pressing direction; (b) longitudinal section of the specimen shown in (a).
Q. Wang et al. / Materials Letters 65 (2011) 2735–2738
2737
Fig. 3. Microstructures of porous CuAlMn shape memory alloys. (a) and (b) microstructures of the matrix before and after quench treatment; (c) and (d) flexural fractographs of the specimens before and after quench treatment.
The specimen density was determined from its weight and physical dimensions, from which the volume fraction of pores was calculated by the following equation:
The sintering was performed in atmospheres of high-purity argon and hydrogen with a ceramic crucible to hold the compact. Before the melting of NaCl, a holding at 785 °C for 3 h was carried out to densify the CuAlMn matrix. The final sintering was conducted at temperatures ranging from 980 to 985 °C for 2–3 h. Under such a high temperature that is well above the 801 °C melting point of NaCl, the NaCl particulates began to melt and flow from the compact, leaving pores in the material and forming porous specimens. After cooling naturally to room temperature in the furnace, the resultant specimens were washed with water to remove the residual NaCl (the specimen mass was tracked regularly. The dissolution process was terminated when the specimen mass was stable and/or decreased to the predicted value equivalent to the mass of starting CuAlMn powders).
with Vf as the volume fraction of pores; ρp and ρb as the densities of the porous and bulk CuAlMn specimens, respectively. The measured porosity ranged from 52.6% to 73.1% corresponding to aforementioned volume percentages of NaCl particulates. The mean diameter of pores is an average of data from more than 10 pores. It ranged from 0.4 to 2.0 mm.
Fig. 4. XRD patterns of the furnace cooled and quenched porous CuAlMn shape memory alloys. 1–7: M18R martensitic phase; 8, 9: γ2 precipitated phase.
Fig. 5. Quasi-static compressive stress–strain curves of the porous CuAlMn shape memory alloys.
Vf = 1−
ρp ρb
ð1Þ
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Q. Wang et al. / Materials Letters 65 (2011) 2735–2738
Quasi-static uniaxial compressive tests were performed on a universal material testing machine at a strain rate of 10 − 2 s − 1 at room temperature. Cylindrical specimens used in compressive tests with 10 mm diameter and 15 mm height were cut by electro-discharge machining. The cross-sectional area used in the calculation of the stresses is the total area of pores and solid on the cross-section of specimen. Optical microscopy and scanning electron microscopy (SEM) were used to characterize the sintered porous CuAlMn specimens. The pores of the specimens were filled with a cold-hardening epoxy resin before grinding and polishing for optical examination. 3. Results and analysis Fig. 2 shows the typical structures of the porous CuAlMn shape memory alloys with an average pore size of 0.4 mm. The porosities of the specimens shown in Fig. 2a, c and d are 67.5%, 52.6% and 73.1%, respectively. Fig. 2b shows the longitudinal section of the specimen shown in Fig. 2a. It can be seen that the granular macroscopic pores are open and uniformly distributed over the matrix of all specimens. The smooth surfaces and stable metal frames of the porous CuAlMn specimens suggest that effective bonding between CuAlMn powders had been achieved before NaCl melted at 801 °C. Therefore, leakage of liquid NaCl would not affect the morphology of the porous CuAlMn specimen, and the final sintering temperature would not affect the shape of the pores either but merely elevate the densification level of the CuAlMn matrix. As previously stated, after final sintering the resultant specimens were cooled naturally to room temperature in the furnace. Fig. 3a illustrates the microstructure of a furnace cooled porous CuAlMn specimen. From the figure, a large number of white strips can be found on the gray metal matrix, while martensite phase cannot be seen clearly. This can be ascribed to the decomposition of β phase during slow cooling [12]. XRD analysis results shown in Fig. 4 indicated that there exists hard brittle γ2 precipitate in furnace cooled specimens. Considering that accumulation of γ2 phase may worsen the performance of SMAs, furnace cooled specimens were therefore quenched after the solution treatment. Fig. 3b shows the microstructure of the specimen subjected to quench treatment after holding at 850 °C for 1 h. It is clear that the CuAlMn matrix is mainly composed of lath-shaped matensite. XRD analysis results showed that the main phase in quenched porous CuAlMn alloys is M18R martensite. This is also concordant with the results reported in [13]. Fig. 3c and d illustrate the comparison of flexural fractographs between the furnace cooled and quenched specimens. Numerous dimples over the fracture surface can be observed from Fig. 3d, while the furnace cooled specimen shows significantly decreased number and size of dimples as shown in Fig. 3c, i.e. the main fractographic appearance is brittle. In order to clarify the influence of microstructures on mechanical properties, a quasi-static compressive test was carried out. Fig. 5 shows the nominal stress–strain curves of the porous CuAlMn specimens with different porosities and heat treatment histories. To understand the detailed stress–strain curves at the lower scale, an inner figure is also presented. From the figure the elastic modulus values of the furnace cooled specimen with the porosity of 60% and the quenched specimens with the porosities of 60.3% and 67.5% were determined as 1.89, 1.12 and 0.50 GPa, respectively. And the yield strength values of these specimens are 18.47, 3.06 and 2.54 MPa, respectively. It is apparent that all compressive stress–strain curves consist of typical three regions, i.e.
elastic, plateau and densification regions. Their long and relatively flat plateau regions give them a good compressive energy-absorption property. The influence of porosity on the compressive behaviors of the present porous CuAlMn SMA is well illustrated: increasing the porosity results in a decreasing stress–strain curve, a short elasticity region and a relatively long plateau region. It is very interesting that the furnace cooled specimen has much higher elastic modulus and plateau region than that of quenched specimen although the two specimens have nearly the same porosity. It can also be found that the quenched specimens have rather smooth stress–strain curves, showing typical deformation characteristics of ductile porous metals. While the stress–strain curve of furnace cooled specimen starts smoothly, but exhibits serrated deformation characteristic at relative high strains. This may be related to the decrease of ductility as a result of the accumulation of γ2 precipitate at boundaries or interfaces in the CuAlMn matrix. At relative high strains during compressive deformation, brittle collapse of pores occurs. Consequently, stress decreases as a result of the decrease of effective cross section in the specimen. With continuous increase of strains, collapsed pore walls start to come into contact, leading to the increase of stress as a result of the increase of effective cross sections. So repeatedly, serrated stress– strain curve appears. Therefore it is obvious that the different compressive deformation behaviors of the quenched and furnace cooled specimens are closely related to the change of microstructures before and after heat treatments. 4. Conclusions A porous CuAlMn SMA has been successfully fabricated via the sintering–dissolution process by using NaCl particulates as spaceholders for the first time. Resultant material exhibits uniformly distributed and interconnected macroscopic pores. Non-martensite phase is present in the matrix of furnace cooled specimens due to the decomposition of β phase during slow cooling. After the solution and quench treatments, the M18R martensite becomes the main phase in the matrix. Quasi-static compressive test reveals that the main deformation manners of the furnace cooled and quenched specimens are brittle and ductile, respectively. It is obvious that the change of main phase in the CuAlMn matrix should be responsible for this. Acknowledgments The study was supported by the Key Program of Natural Science Foundation of Tianjin (no. 09JCZDJC22200) and the Natural Science Foundation of Hebei Province (no. E2009000060). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
Qidwai MA, Entchev PB, Lagoudas DC, et al. Int J Solids Struct 2001;38:8653. Brailovski V, Trochu F. Biomed Mater Eng 1996;6:291. Hu GX, Zhang LX. Comput Mater Sci 2008;42:558. Li DS, Zhang YP, Ma X, et al. J Alloys Compd 2009;47:L1. Hosseini SA, Sadrnezhaad SK, Ekrami A. Mater Sci Eng C 2009;29:2203. Wang QZ, Han FS, Gao ZY, et al. J Alloys Compd 2006;425:200. Wang B, Zhang E. Int J Mech Sci 2008;50:550. Zhao Y, Sun D. Scr Mater 2001;44:105. Zhao Y, Fung T, Zhang L, et al. Scr Mater 2005;52:295. Wang QZ, Cui CX, Liu SJ, et al. Mater Sci Eng A 2010;527:1275. Bansiddhi A, Dunand DC. Acta Biomater 2008;4:1996. Wang QZ, Han FS, Wu J, et al. Phys Status Solidi A 2006;203:825. Chen J, Li Z, Zhao YY. J Alloys Compd 2009;480:481.
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
Fabrication of a porous CuAlMn shape memory alloy by the sintering–dissolution process Qingzhou Wang ⁎, Chunxiang Cui, Qian Wang, Najun Yan School of Materials Science and Engineering, Hebei University of Technology, Road No. 1, Dingzigu, Hongqiao District, Tianjin 300130, China
a r t i c l e
i n f o
Article history: Received 14 March 2011 Accepted 21 May 2011 Available online 30 May 2011 Keywords: Porous materials Shape memory materials Powder technology
a b s t r a c t A porous CuAlMn shape memory alloy with adjustable pore characteristics and mechanical properties has been successfully fabricated via the sintering–dissolution process for the first time. Micro amount of fluoride was added to accelerate the sintering of CuAlMn powders covered with the oxide film. The resultant material exhibits uniformly distributed and interconnected macroscopic pores. In order to eliminate possible adverse effects of the microstructures resulted from slow cooling on the properties of the resultant materials, appropriate solution and quench treatments were carried out. The microstructures and quasi-static compressive deformation behaviors of the resultant materials were also examined before and after treatments. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Initiating from early studies in Russian and further developments in China, porous shape memory alloys (SMAs) have been extensively investigated [1]. They have lower density and the possibility of further enhancing energy absorption and damping capabilities. Up to now investigations of porous SMAs are mainly focused on the Ni–Ti alloys [2–5]. But the high cost is its main limitation for engineering applications. As a result, Cu-based SMAs are attracting more and more attention in recent years. In our previous work a porous Cu-based SMA was fabricated via an air pressure infiltration process [6]. However, in the fabrication process there are still strict restrictions to the superheating temperature of Cu-based SMA melt, the choice of dissoluble space-holders as well as the velocity of infiltration process due to the relatively high melting point of Cu-based SMAs (over 1000 °C). Fast solidification of melting metals enhances the level of difficulty in techniques. It is therefore necessary to develop new techniques to offer a wider range of structures and properties while continuing to improve original infiltration technique. Powder metallurgy technique possesses unique advantages in fabrication of porous metals [7,8]. Among the powder metallurgical techniques sintering–dissolution process (SDP) is a promising route for manufacturing open-celled porous metals [8–10]. The most outstanding technical advantage of SDP is much wider range of structure parameters, thus providing more opportunities for resultant porous metals to meet the desired behaviors. Thus, in the present study the SDP is employed to fabricate a porous CuAlMn SMA using NaCl particulates as space-holders. It is interesting to note that ⁎ Corresponding author. Fax: + 86 22 60204125. E-mail address: [email protected] (Q. Wang). 0167-577X/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2011.05.092
using NaCl as space-holders possesses numerous advantages, such as low cost, fast dissolution in water, reduced corrosive attack of metal during dissolution, free of toxicity, in addition to significant widening of structure parameters such as porosity, pore shape, pore size and homogeneity of pore distribution, etc. [11]. 2. Experimental An SDP process was employed to fabricate the porous SMA specimens from the Cu–11.7Al–2.49Mn (wt.%) SMA powders with a purity of more than 99% (shown in Fig. 1a) prepared by an atomization method and domestic NaCl particulates (shown in Fig. 1d). From the figure it can be seen that the CuAlMn powders with an average diameter of about 15 μm are round or elliptic, while NaCl particles are around 0.4 mm in size and cuboidal in shape. Fig. 1b shows the DSC analysis result of the starting CuAlMn powders. The characteristic temperatures As, Af, Ms and Mf of the Cu–11.7Al– 2.49Mn alloy were determined as 183, 216, 156 and 142 °C, respectively. The CuAlMn powders were blended with micro amount of fluoride (the mixture of AlF3 and CaF2, less than 0.08 wt.%) to remove the surface oxidation layer by the formation of gaseous AlOF from the reaction between Al2O3 and AlF3. The reaction can be accelerated further by the presence of liquid phase which appeared in Al–Ca fluoride. The mixture was then mechanically milled for 1.5 h under an atmosphere of argon to get the final well-mixed powders as shown in Fig. 1c. It can be seen that milled powders have irregular shapes and coarse surfaces. These features can improve the coupling between CuAlMn powders and facilitate the pressing process of green compact. Ball milled CuAlMn powders and NaCl particulates with a chosen size were weighed using an electronic balance with an accuracy of 0.01 g. They were blended together in a stirrer at a rotation speed of
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Fig. 1. Raw materials used in the SDP. (a) Starting CuAlMn powders; (b) DSC result of starting CuAlMn powders; (c) ball milled mixture of CuAlMn and fluoride powders; (d) NaCl particulates.
80 rpm for 10 min. The volume percentages of NaCl were varied from 50% to 70% with a gap of 5%. Well mixed powders with a small amount of ethanol were uniaxially pressed into a steel die with an inner diameter of 20 mm under pressures ranged from 400 to 450 MPa.
Relatively high pressures were adopted in the present study to obtain high density green compacts. However, excessive high pressure should be avoided due to the existence of brittle NaCl particulates that are liable to be crushed under too high pressure.
Fig. 2. Typical structure of the porous CuAlMn shape memory alloys. (a), (c) and (d) cross-section of the specimens perpendicular to the pressing direction; (b) longitudinal section of the specimen shown in (a).
Q. Wang et al. / Materials Letters 65 (2011) 2735–2738
2737
Fig. 3. Microstructures of porous CuAlMn shape memory alloys. (a) and (b) microstructures of the matrix before and after quench treatment; (c) and (d) flexural fractographs of the specimens before and after quench treatment.
The specimen density was determined from its weight and physical dimensions, from which the volume fraction of pores was calculated by the following equation:
The sintering was performed in atmospheres of high-purity argon and hydrogen with a ceramic crucible to hold the compact. Before the melting of NaCl, a holding at 785 °C for 3 h was carried out to densify the CuAlMn matrix. The final sintering was conducted at temperatures ranging from 980 to 985 °C for 2–3 h. Under such a high temperature that is well above the 801 °C melting point of NaCl, the NaCl particulates began to melt and flow from the compact, leaving pores in the material and forming porous specimens. After cooling naturally to room temperature in the furnace, the resultant specimens were washed with water to remove the residual NaCl (the specimen mass was tracked regularly. The dissolution process was terminated when the specimen mass was stable and/or decreased to the predicted value equivalent to the mass of starting CuAlMn powders).
with Vf as the volume fraction of pores; ρp and ρb as the densities of the porous and bulk CuAlMn specimens, respectively. The measured porosity ranged from 52.6% to 73.1% corresponding to aforementioned volume percentages of NaCl particulates. The mean diameter of pores is an average of data from more than 10 pores. It ranged from 0.4 to 2.0 mm.
Fig. 4. XRD patterns of the furnace cooled and quenched porous CuAlMn shape memory alloys. 1–7: M18R martensitic phase; 8, 9: γ2 precipitated phase.
Fig. 5. Quasi-static compressive stress–strain curves of the porous CuAlMn shape memory alloys.
Vf = 1−
ρp ρb
ð1Þ
2738
Q. Wang et al. / Materials Letters 65 (2011) 2735–2738
Quasi-static uniaxial compressive tests were performed on a universal material testing machine at a strain rate of 10 − 2 s − 1 at room temperature. Cylindrical specimens used in compressive tests with 10 mm diameter and 15 mm height were cut by electro-discharge machining. The cross-sectional area used in the calculation of the stresses is the total area of pores and solid on the cross-section of specimen. Optical microscopy and scanning electron microscopy (SEM) were used to characterize the sintered porous CuAlMn specimens. The pores of the specimens were filled with a cold-hardening epoxy resin before grinding and polishing for optical examination. 3. Results and analysis Fig. 2 shows the typical structures of the porous CuAlMn shape memory alloys with an average pore size of 0.4 mm. The porosities of the specimens shown in Fig. 2a, c and d are 67.5%, 52.6% and 73.1%, respectively. Fig. 2b shows the longitudinal section of the specimen shown in Fig. 2a. It can be seen that the granular macroscopic pores are open and uniformly distributed over the matrix of all specimens. The smooth surfaces and stable metal frames of the porous CuAlMn specimens suggest that effective bonding between CuAlMn powders had been achieved before NaCl melted at 801 °C. Therefore, leakage of liquid NaCl would not affect the morphology of the porous CuAlMn specimen, and the final sintering temperature would not affect the shape of the pores either but merely elevate the densification level of the CuAlMn matrix. As previously stated, after final sintering the resultant specimens were cooled naturally to room temperature in the furnace. Fig. 3a illustrates the microstructure of a furnace cooled porous CuAlMn specimen. From the figure, a large number of white strips can be found on the gray metal matrix, while martensite phase cannot be seen clearly. This can be ascribed to the decomposition of β phase during slow cooling [12]. XRD analysis results shown in Fig. 4 indicated that there exists hard brittle γ2 precipitate in furnace cooled specimens. Considering that accumulation of γ2 phase may worsen the performance of SMAs, furnace cooled specimens were therefore quenched after the solution treatment. Fig. 3b shows the microstructure of the specimen subjected to quench treatment after holding at 850 °C for 1 h. It is clear that the CuAlMn matrix is mainly composed of lath-shaped matensite. XRD analysis results showed that the main phase in quenched porous CuAlMn alloys is M18R martensite. This is also concordant with the results reported in [13]. Fig. 3c and d illustrate the comparison of flexural fractographs between the furnace cooled and quenched specimens. Numerous dimples over the fracture surface can be observed from Fig. 3d, while the furnace cooled specimen shows significantly decreased number and size of dimples as shown in Fig. 3c, i.e. the main fractographic appearance is brittle. In order to clarify the influence of microstructures on mechanical properties, a quasi-static compressive test was carried out. Fig. 5 shows the nominal stress–strain curves of the porous CuAlMn specimens with different porosities and heat treatment histories. To understand the detailed stress–strain curves at the lower scale, an inner figure is also presented. From the figure the elastic modulus values of the furnace cooled specimen with the porosity of 60% and the quenched specimens with the porosities of 60.3% and 67.5% were determined as 1.89, 1.12 and 0.50 GPa, respectively. And the yield strength values of these specimens are 18.47, 3.06 and 2.54 MPa, respectively. It is apparent that all compressive stress–strain curves consist of typical three regions, i.e.
elastic, plateau and densification regions. Their long and relatively flat plateau regions give them a good compressive energy-absorption property. The influence of porosity on the compressive behaviors of the present porous CuAlMn SMA is well illustrated: increasing the porosity results in a decreasing stress–strain curve, a short elasticity region and a relatively long plateau region. It is very interesting that the furnace cooled specimen has much higher elastic modulus and plateau region than that of quenched specimen although the two specimens have nearly the same porosity. It can also be found that the quenched specimens have rather smooth stress–strain curves, showing typical deformation characteristics of ductile porous metals. While the stress–strain curve of furnace cooled specimen starts smoothly, but exhibits serrated deformation characteristic at relative high strains. This may be related to the decrease of ductility as a result of the accumulation of γ2 precipitate at boundaries or interfaces in the CuAlMn matrix. At relative high strains during compressive deformation, brittle collapse of pores occurs. Consequently, stress decreases as a result of the decrease of effective cross section in the specimen. With continuous increase of strains, collapsed pore walls start to come into contact, leading to the increase of stress as a result of the increase of effective cross sections. So repeatedly, serrated stress– strain curve appears. Therefore it is obvious that the different compressive deformation behaviors of the quenched and furnace cooled specimens are closely related to the change of microstructures before and after heat treatments. 4. Conclusions A porous CuAlMn SMA has been successfully fabricated via the sintering–dissolution process by using NaCl particulates as spaceholders for the first time. Resultant material exhibits uniformly distributed and interconnected macroscopic pores. Non-martensite phase is present in the matrix of furnace cooled specimens due to the decomposition of β phase during slow cooling. After the solution and quench treatments, the M18R martensite becomes the main phase in the matrix. Quasi-static compressive test reveals that the main deformation manners of the furnace cooled and quenched specimens are brittle and ductile, respectively. It is obvious that the change of main phase in the CuAlMn matrix should be responsible for this. Acknowledgments The study was supported by the Key Program of Natural Science Foundation of Tianjin (no. 09JCZDJC22200) and the Natural Science Foundation of Hebei Province (no. E2009000060). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
Qidwai MA, Entchev PB, Lagoudas DC, et al. Int J Solids Struct 2001;38:8653. Brailovski V, Trochu F. Biomed Mater Eng 1996;6:291. Hu GX, Zhang LX. Comput Mater Sci 2008;42:558. Li DS, Zhang YP, Ma X, et al. J Alloys Compd 2009;47:L1. Hosseini SA, Sadrnezhaad SK, Ekrami A. Mater Sci Eng C 2009;29:2203. Wang QZ, Han FS, Gao ZY, et al. J Alloys Compd 2006;425:200. Wang B, Zhang E. Int J Mech Sci 2008;50:550. Zhao Y, Sun D. Scr Mater 2001;44:105. Zhao Y, Fung T, Zhang L, et al. Scr Mater 2005;52:295. Wang QZ, Cui CX, Liu SJ, et al. Mater Sci Eng A 2010;527:1275. Bansiddhi A, Dunand DC. Acta Biomater 2008;4:1996. Wang QZ, Han FS, Wu J, et al. Phys Status Solidi A 2006;203:825. Chen J, Li Z, Zhao YY. J Alloys Compd 2009;480:481.