Fabrication and mechanical behavior of bulk nanoporous Cu via chemical de-alloying of Cu–Al alloys

Author’s Accepted Manuscript Fabrication and Mechanical Behavior of Bulk Nanoporous Cu via Chemical De-Alloying of CuAl Alloys Fei Chen, Xi Chen, Lijie Zou, Yao Yao, Yaojun Lin, Qiang Shen, Enrique J. Lavernia, Lianmeng Zhang www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(16)30171-X http://dx.doi.org/10.1016/j.msea.2016.02.055 MSA33360

To appear in: Materials Science & Engineering A Received date: 28 December 2015 Revised date: 18 February 2016 Accepted date: 18 February 2016 Cite this article as: Fei Chen, Xi Chen, Lijie Zou, Yao Yao, Yaojun Lin, Qiang Shen, Enrique J. Lavernia and Lianmeng Zhang, Fabrication and Mechanical Behavior of Bulk Nanoporous Cu via Chemical De-Alloying of Cu-Al Alloys, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.02.055 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Fabrication and Mechanical Behavior of Bulk Nanoporous Cu via Chemical De-Alloying of Cu-Al Alloys Fei Chen1*, Xi Chen1, Lijie Zou1, Yao Yao1, Yaojun Lin1, Qiang Shen1, Enrique J. Lavernia2, Lianmeng Zhang1* 1 State Key Lab of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, China 2 Department of Chemical Engineering and Materials Science, University of California at Irvine, Irvine, CA 92697, USA *Corresponding Author: [email protected]

Abstract We report on a study of the influence of microstructure on the mechanical behavior of bulk nanoporous Cu fabricated by chemical de-alloying of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 (at.%) alloys. The precursor Cu-Al alloys were fabricated using arc melting and bulk nanoporous Cu was obtained by subsequent de-alloying of Cu-Al alloys in 20 wt.% NaOH aqueous solution at a temperature of 65 oC. We studied the microstructure of the precursor Cu-Al alloys, as well as that of the as de-alloyed bulk nanoporous Cu, using X-ray diffraction, scanning electron microscopy and energy dispersive spectrometry. Moreover, the compressive strength of bulk nanoporous Cu was measured and the relationship between microstructure and mechanical properties was studied. Our results show that the microstructure of bulk nanoporous Cu is characterized by bi-continuous interpenetrating ligament-channels with a ligament size of 110-150 nm (for Cu50Al50), 150-180 nm (for Cu40Al60) and 150-200 nm (for Cu33Al67). Interestingly the microstructure of de-alloyed Cu30Al70 is bimodal with nanopores (100’s nm) and interdispersed featureless regions a few microns in size. The compressive strength increased with decreasing volume fraction of porosity; as porosity increased 56.3 % to 73.9 %, the compressive strength decreased from 17.18 MPa to 2.71 MPa. Keywords: Nanoporous Cu; Cu-Al alloys; chemical de-alloying; compressive strength

1. Introduction Nanoporous metals are of interest, partly because their unusual microstructure reportedly leads to some notable characteristics, such as: enhanced mechanical strength electrocatalytic properties

[3]

[1]

, piezoelectricity

, surface-enhanced Raman spectroscopic properties

potential applications in energy storage

[5]

, catalyzing carrier

[6]

, detectors

[7]

[4]

[2]

,

, as well as

, etc. De-alloying,

which refers to the selective dissolution of one or more active elements out of an alloy [8], represents the most widely used method to fabricate nonoporous metals and has rapidly become an important area of research for the international scientific community, since the pioneering work of R. Li and K. Sieredzki [9]. Publications describing the fabrication of nanoporous Au using de-alloying first appeared in the 1990’s [10-13] and typically involved Au-Ag systems [14-17]. In addition to alloys based on Au-Ag, recent research has focused on the preparation of nanoporous metals using various precursor alloy systems, such as Cu-Zr [18], Cu-Mg [19], Cu-Al [20], Au-Zn [21], Cu-Mn [22], etc., by electrochemical or chemical de-alloying methods. As examples, M. Hakamada et al.

[23]

fabricated nanoporous Ni and

nanoporous Cu by electrochemical de-alloying rolled Ni0.3Mn0.7 and Cu0.3Mn0.7 alloys. Statistical analysis of 100 ligaments in the TEM images revealed that the ligament size of nanoporous Ni (8 nm ± standard deviation of 2 nm) is much smaller than that of nanoporous Cu (24 nm ± 5 nm). R. Morrish et al. [24] fabricated thin film Cu-Au alloys by alternately depositing Au and Cu followed by a heat treatment. Precursor Cu-Au alloys were de-alloyed in concentrated HNO3 solution from the starting Cu content of 75 at.% to as low as 20 at.%. They studied the effect of different Cu contents in Cu-Au alloys on the morphology of as-de-alloyed nanoporous Au. T. Song et al.

[25]

fabricated

nanoporous Cu with uniform nanopores by electrochemical de-alloying annealed Cu25Al75 foils

with thickness of 0.6 mm. They investigated the dependence of formation of nanoporous Cu on the electrochemical de-alloying conditions of the precursor Cu-Al alloy. The microstructure that forms during de-alloying at -0.1 V is coarser (average ligament width of 204 ± 85 nm) than that formed during de-alloying at -0.2 V (average ligament width of 196 ± 48 nm). De-alloying of Al and Al2Cu phases occurred at -0.5 V, which resulted in a nanoporous structure with a much smaller average ligament width of 52 ± 10 nm. In fact review of the literature shows that most of these investigations have focused on the effects of alloy composition and de-alloying conditions on the microstructure of nanoporous metals. Moreover, typical precursor alloys used in these studies are usually fabricated by melting, followed by rolling, sputtering or deposition processes, and therefore the morphology of final nanoporous metal is generally a thin film, ribbon, or foil, with thickness values that are smaller than 1 mm. There are very few studies that describe the fabrication of bulk (i.e., in contrast to film) nanoporous metals. In related work, Y. Sun et al. [26] used a two-step method to fabricate crack-free bulk nanoporous Au with millimeter dimensions, no volume change and no significant cracking. In this work, bulk Au-Ag alloy pellets (30 at.% Au) were sectioned into 4 mm × 4 mm × 0.45 mm and then annealed at 850 oC for 10 h. De-alloying was performed in multiple stages and with varying HNO3 concentrations. Isolated cracks were observed in the de-alloyed nanoporous Au, but these were few in number and did not cause specimen fracture during handling. In other work, X. M. Zhang et al.

[27]

obtained rectangular (8 mm × 8 mm × 10 mm) and cylindrical (Φ 8 mm × 20 mm)

bulk, three-dimensional bimodal porous Cu by de-alloying Gasar Cu-34.6 wt.% Mn alloy in 10 vol.% HCl for 27 h and characterized the microstructure. The Gasar Cu-34.6 wt.% Mn microstructure was characterized by lotus root like oriented column pores, homogeneously arranged in the alloy matrix. Moreover, the de-alloyed bulk nanoporous Cu contained uniformly distributed, sub-millimeter pores, with large oriented Gasar pores and small random de-alloyed pores.

Inspection of the published literature shows that the porous metals that are typically used for mechanical behavior studies normally contain micron size pores, and are fabricated using a variety of approaches, including: thermal treatment and a pore-forming agent [31,32]

[28,29,30]

, metal infiltration

, powder sintering [33,34] and unidirectional casting of molten metal [35,36], to name a few. In fact,

only very few studies of the mechanical behavior of nanoporous metals have been published. In one such study by T. Balk et al.

[37]

nanoporous Au was obtained by de-alloying Ag-30 at.% Au alloy

derived from cold-rolled and annealed plates, with a gauge length of 430 μm and a thickness of 30-400 μm. The compressive strength and tensile strength of the bulk nanoporous Au were reported to be 15 MPa and 11 MPa, respectively, with an average porosity of 30 %. In other work, Q.Q. Kong et al. [38] prepared bulk nanoporous Cu by chemically de-alloying spark plasma sintered Al-22 at.% Cu with 80 % porosity and a compressive strength of 5.8 MPa. In view of the above discussion it is evident that there is very limited information on the synthesis, microstructure and mechanical behavior of bulk nanoporous metals. Accordingly, in our work we describe a simple chemical de-alloying method to produce bulk nanoporous Cu with uniform and well dispersed pores. We selected Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 (at. %) as target compositions for several reasons. First, Cu alloys are characterized by good electrical and thermal conductivity, high strength, abrasion resistance and oxidation resistance

[39]

. Second, it is

possible to de-alloy Al out of the Cu-Al alloys because of the large difference in the standard potential between Al and Cu (-1.662 V vs. standard hydrogen electrode (SHE) for Al/Al3+ and 0.342 V vs. SHE for Cu/Cu2+) [40]. Third, de-alloying of Cu-Al alloys is relatively easy because Al cannot only be dissolved in an acidic solution, but also can be dissolved in an alkaline solution since Al is amphoteric. Accordingly, the alloys were prepared by arc melting and the de-alloying process was conducted in a 20 wt.% NaOH aqueous solution at the temperature of 65 oC. The objectives of the present study are three-fold: first, to establish the feasibility of using a

simple chemical dissolution method to prepare bulk nanoporous Cu; second, to study the influence of the nanoporous microstructure on the mechanical behavior of Cu; and third to provide fundamental insight into the factors that govern the size, distribution and geometry of the nanopores.

2. Experimental procedures 2.1 Preparation of master alloys and bulk nonoporous Cu The starting materials used in our study are commercially available pure Cu and Al powders (Cu: 99.9 %, mean particle size: 10 μm, Alfa Aesar, USA; Al: 99.9 %, mean particle size: 10 μm, Alfa Aesar, USA). The Cu and Al powders were mixed using the following proportions: 50/50, 40/60, 33/67, 30/70 at.% (Cu/Al) ratio by ball milling for 24 h and then compacted using a dry pressing facility. The precursor Cu-Al alloys were then melted three times in the arc-melting furnace (ACM-S01, ULVAC JAPAN LTD, Japan), in order to achieve a good homogeneity. Following this step, bulk precursor Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 (at.%) alloys, with dimensions of 6 mm × 6 mm × 3 mm were obtained by wire cutting. The de-alloying process of the precursor Cu-Al alloys was conducted in 20 wt.% NaOH solution to dissolve the Al but not the Cu atoms. The de-alloying temperature was maintained at a constant value of 65 oC using a water bath for 24 h. The as-de-alloyed bulk nanoporous Cu were rinsed with distilled water and dehydrated alcohol, and then fully dried and kept under vacuum drying conditions in order to minimize oxidation. 2.2 Materials characterizations X-ray diffraction (XRD) patterns of the precursor Cu-Al alloys were recorded using a Rigaku diffractometer with Cu-Kα radiation. Zeiss Ultra Plus SEM operated in back scattered electron mode at 10 kv was used to observe the phase distribution of the Cu-Al alloys.

Archimedes’ Principle was used to determine their relative density and energy dispersive spectrometer (EDS) was used to analysis their elemental contents and homogeneity. Microstructural characterization of the as-de-alloyed bulk nanoporous Cu was accomplished using an X-ray diffraction (XRD) diffractometer with Cu-Kα radiation and a Zeiss Ultra Plus scanning electron microscope (SEM) operated in secondary electron mode at 5 kV. The porosity of bulk nanoporous Cu was measured using Archimedes principle and the residual Al contents in the bulk nanoporous Cu were measured using ICP (Optima 4300 DV, American PerkinElmer, America) spectrometry test. The mechanical behavior of the bulk nanoporous Cu was determined in static compression tests conducted with a universal testing machine, at room temperature and at a strain rate of 0.5 mm/min. Each strength value was averaged over three measurements.

3. Results 3.1 Phase compositions and microstructure of Cu-Al alloys Fig. 1(a) shows the precise phase compositions of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 alloys, calculated on the basis of the line compounds in the Cu-Al phase diagram. The Cu50Al50 and Cu33Al67 alloys are supposed to be single intermetallic phase, composed of AlCu and Al2Cu, respectively. Cu40Al60 consists of AlCu and Al2Cu phases whereas Cu30Al70 contains α-Al (Cu) and Al2Cu phases. Fig. 1(b) shows the XRD patterns of the precursor Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 alloys, confirming the phase diagram predictions. The phase contents of the precursor Cu-Al alloys, calculated from the phase diagram, as well as those calculated on the basis of the Maud program [41-44] are reported in Table 1, confirming that the software predictions are consistent with the results calculated from phase diagram. The quantification of phase compositions of the Cu-Al alloys with different Cu/Al ratios facilitates comparison of the four Cu-Al alloys by establishing the correlation between phase and microstructure. The measured

density values are also shown in Table 1. It is apparent that all Cu-Al alloys have a high relative density, exceeding 98 %, indicating that the powders are almost were efficiently compacted during the melting process. The SEM images (back scattered electron images) of the Cu-Al alloys with different Cu/Al ratios prepared by arc melting are shown in Fig. 2. It is observed clearly from Fig. 2(a) and (c) that both Cu50Al50 and Cu33Al67 alloys show rather weak contrast, consistent with the fact that the two alloys are primarily single-phase alloys, with AlCu phase and Al2Cu phase, respectively. Given that in the back-scattered electron image, the higher contrast area corresponds to higher atomic number elements, the contrast of the AlCu phase will be higher than that corresponding to the Al2Cu phase. As a result, it is shown in Fig. 2(b) that Cu40Al60 is composed of the Al2Cu phase and the AlCu phase, as noted in the figure. Cu30Al70 alloy shows a different phase distribution, which solidifies as Al2Cu phase and a coexisting structure containing α-Al (Cu) and Al2Cu phases, as shown in Fig. 2(d). All these back scattered electron images of Cu-Al alloys are in agreement with their XRD results. Fig. 3 shows the EDS (energy dispersive spectrometer) images of Cu-Al alloys, while their mapping elemental weight and atomic contents are listed in Table 2. It is suggested that the elemental contents are in accordance to the original Cu/Al ratio and all precursor Cu-Al alloys are melted homogeneously. Fig. 4 shows SEM images of the microstructure of the precursor Cu30Al70 alloy and its corresponding EDS images. It is observed that the Al2Cu phase is distributed as islands in the coexisting α-Al (Cu) and Al2Cu phases, which is shown in Fig. 4(a). Coexisting α-Al (Cu) and Al2Cu phases presents a lamellar eutectic structure which can be seen from the SEM image at a higher magnification in inset of Fig. 4(b). Fig. 4(c) and (d) shows the energy dispersive spectrometer (EDS) images of Cu30Al70 alloy and the higher magnification lamellar eutectic α-Al (Cu)-Al2Cu, respectively. It is evident that a homogeneous composition of the lamellar eutectic

α-Al (Cu)-Al2Cu structure exists in the Cu30Al70 alloy. Such coexisting lamellar eutectic α-Al (Cu)-Al2Cu is proposed to be related to the eventual formation of a novel porous structure with nanopores and nanoporesand interdispersed featureless regions by subsequent de-alloying of Cu30Al70 alloy; this is addressed in the discussion section. 3.2 Phase compositions and microstructure of nanoporous Cu Fig. 5 shows the XRD patterns of bulk nanoporous Cu obtained from chemical de-alloying of the Cu-Al alloys in the 20 wt.% NaOH solution. It is seen from Fig. 5 that after de-alloying the precursor Cu-Al alloys, only Cu phase is detected in the as-de-alloyed bulk nanoporous Cu, indicating that all AlCu, Al2Cu and α-Al (Cu) in these Cu-Al alloys are almost fully de-alloyed to form Cu. Fig. 6 shows both the experimental and theoretical values of porosity of bulk nanoporous Cu. The porosity increased from 56.3 % to 73.9 % with increasing Al contents from 50 at. % to 70 at.% in the precursor Cu-Al alloys. The porosity of bulk nanoporous Cu de-alloyed from Cu30Al70 alloy is 73.9 %, which is the maximum value obtained in the present study. The ICP test results shown in Table 3 indicate that the residual Al contents in bulk nanoporous Cu is relatively small, no more than 2 wt.% after de-alloying. This is believed to be the reason why the amount of porosity present in the bulk nanoporous Cu samples is slightly lower than the predicted theoretical value. Fig. 7 shows the section-view microstructures of bulk nanoporous Cu through de-alloying of the Cu-Al alloys in the 20 wt.% NaOH solution at temperature of 65 oC. For the Cu50Al50 and Cu33Al67 single-phase alloys, a uniform porous structure can be observed in the as-de-alloyed bulk nanoporous Cu, as shown in Fig. 7(a) and Fig. 7(c), respectively. Bulk nanoporous Cu exhibits a typical bicontinuous interpenetrating ligament-channel structure with ligament size of 110-150 nm (for Cu50Al50 ) and 150-200 nm (for Cu33Al67), which can be seen from Fig. 7(c) and Fig. 7(g),

respectively. Fig. 7(b) shows the microstructure of the sample de-alloyed from Cu40Al60 two-phase alloy, a uniform porous structure can also be observed throughout the entire as de-alloyed bulk sample. The associated bicontinuous, interpenetrating ligament-channel structure, with a ligament size of 150-180 nm, is observed clearly in Fig. 7(f). While the as-de-alloyed Cu30Al70 two-phase alloy shows a different porous structure, which can be seen in Fig. 7(d) and Fig. 7(h). The microstructure in this figure is unusual in that it reveals interdispersed featureless regions a few microns in size and nanopores 100’s nm. 3.3 Mechanical behavior of nanoporous Cu Fig. 9(a) shows the compressive stress-strain curves of bulk nanoporous Cu with different amounts of porosity. The stress-strain curves typically consist of three regions. First, region I is a linear elastic region with a steep slope at a low strain. Stress increases dramatically with the increase of strain in this region. Then, region II is a stress plateau region that occurs after yielding where the stress is nearly constant with the increasing strain. Last, region III is a densification region where stress increases sharply with the increase of strain. These three regions are shown in the insert of Fig. 9(a). All bulk nanoporous Cu in this study shows similar stress-strain curves with a long elastic region, the end of which is assumed to be the compressive strength. It is observed that the samples de-alloyed from Cu50Al50 and Cu40Al60 alloys (56.3 % porosity, 110-150 nm ligaments for Cu50Al50 sample; 65.3% porosity, 150-200 nm ligaments for Cu40Al60 sample) show a ductile-like stress-strain curve during compression. The initial elastic deformation, with an initial peak stress at 17.18 MPa and 9.45 MPa respectively, is followed by a gradual transition to plastic deformation. Serrations are evident in the plateau region of both samples. The two samples exhibit a plateau stress of 4.20 MPa and 2.77 MPa, respectively and then begin to densify at ~37.1 % and ~41.2 % strain. The sample de-alloyed from Cu33Al67 (72.6 % porous,

150-180 nm ligaments) is, with an initial peak stress at 3.26 MPa, significant lower than that of the Cu50Al50 and Cu40Al60 samples. The stress-strain curve also has a characteristic of ductile porous metals in compression, including a long plateau region with a plateau stress of 1.22 MPa followed by densification beginning at ~46.3 % strain. The sample de-alloyed from Cu30Al70 (73.9 % porous) with interdispersed featureless regions a few microns in size and nanopores 100’s nm shows the lowest initial peak stress of 2.71 MPa and plateau stress of 0.50 MPa. The long plateau region ends at a strain of ~50.0 %, which represents the maximum strain among all materials studied. The Young’s modulus is calculated as the slope of region I where stress increases dramatically with increasing strain. It can be seen from Fig.9 (a) that the Young’s modulus of bulk nanoporous Cu decreases with increasing porosity while the Young’s modulus of bulk nanoporous Cu de-alloyed from Cu33Al67 and Cu30Al70 alloys shows the similar value. It should also be noted that the smoothness of plateau region increases with increasing porosity.

4. Discussion 4.1 The de-alloying mechanism of Cu-Al alloy systems The results obtained in our study provide support to the original hypothesis that AlCu, Al2Cu and α-Al (Cu) can be fully de-alloyed in the 20 wt.% NaOH solution, resulting in the formation of bulk nanoporous Cu with uniform porous structures, as seen in Fig. 7. The corresponding reaction involving Al atoms in the NaOH solution is shown in Eq. (1): 2Al+2OH-+2H2O

2AlO2-+3H2

(1)

We propose that the mechanism governing the evolution of nanopores during chemical de-alloying of Cu-Al alloys is qualitatively similar to that active during the chemical de-alloying process of Au-Ag alloys

[45]

, as seen in Fig. 8. This mechanism is described as follows. First,

consider a surface layer of Cu-Al atoms as the starting point. When this surface is immersed in an

electrolyte (20 wt.% NaOH solution), an Al atom is dissolved and a vacancy is formed, leaving the surrounding Cu atoms with an absence of atomic coordination. When compared to Al atoms located in the bulk region (e.g., away from the surface), the Al atoms that were originally laterally coordinated to the dissolved Al atom are more susceptible to dissolution since they have fewer bonds. Second, the etch front initially spreads laterally. Third, prior to corrosion of the next layer, the Cu atoms that remain on the surface region agglomerate into clusters through diffusion rather than dissolution. Fourth, the dissolution of Al atoms form corrosion channels in the bulk while Cu atoms in the interior continue to agglomerate into clusters via diffusion. Fifth, when Al atoms in the second layer of alloy are attacked, Cu atoms from that layer are released onto the surface. These Cu atoms will be transported to the base of the clusters formed during the first layer of dissolution before the next layer is attacked. Hence, there is always an island within a distance λ away. In any case, a characteristic spacing, λ, between Cu clusters is established. These clusters will coarsen and passivate given sufficient time, and thus Al atoms dissolve gradually while simultaneously Cu clusters appear as a function of time. When the Cu clusters attain a sufficient size, they will detach form the surface and form ligaments. This sequence of events results in the formation of porosity in the bulk of the alloy. It is important to note, however, that the precise de-alloying mechanism that is active will depend on phase composition. The de-alloying mechanism that governs the formation of nanopores in a single-phase alloy is relative simple. The atomic scale phase separation occurs at grain surfaces in the precursor alloy, continuously releasing Cu atoms. The agglomeration of Cu clusters contributes to the formation of porosity in the bulk of the alloy. This is consistent with the de-alloying mechanism that has been proposed for Au-Ag [45]. Hence, compared to Cu50Al50, and Cu33Al67 single-phase alloys, the de-alloying process that governs the behavior of the Cu40Al60 two-phase alloy is more complex. The number of Al atoms in Al2Cu is higher than that in AlCu.

Accordingly, when de-alloying begins, the Al atoms in Al2Cu dissolve preferentially, forming corrosion channels, which promotes further dissolution of AlCu. As the corrosion grows into the bulk region of the Cu-Al alloys, Cu atoms are left without any bonds. Prior to dissolution of Al atoms in a near-surface layer, Cu atoms on the surface spread and agglomerate into clusters. As Al atoms continue to dissolve, more Cu atoms are released and spread to the Cu clusters in the upper layer. In the case of the Cu30Al70 alloy, it is reasonable to assume that the de-alloying process that is active in this alloy is more complex than that in Cu50Al50, Cu40Al60 and Cu33Al67, because of the formation of homogeneous interdispersed featureless regions in the as-de-alloyed bulk nanoporous Cu shown in Fig. 7(d) involve diffusion of Cu atoms, consistent with the co-existence of lamellar eutectic α-Al (Cu)-Al2Cu, as observed in Fig. 2(d). The interdispersed featureless regions are likely attributable to the oriented rearrangement of Cu atoms in the lamellar eutectic α-Al (Cu)-Al2Cu as they diffuse to the position corresponding to the Al2Cu phase. Accordingly, and as illustrated in Fig. 4, the microstructure of the two-phase alloys consists of. nanopores (100’s nm) with interdispersed featureless regions a few microns in size. 4.2 The effect of porosity and pore morphology on the mechanical behavior of bulk nanoporous Cu As reported in the literature, a typical stress-strain curve formed during the compression of a homogeneous porous metals, that is containing uniform pores with approximately the same pore size, consists of three regions: first, short elastic deformation (region I); then, a stress drop that evolves into a long plateau stress (region II); and finally, a densification region at high strains (region III) [46, 47, 48]. The compressive stress-strain curve of bulk nanoporous Cu shown in Fig. 9(a) corresponds to the typical stress-strain curve of porous metals. The elastic region (region I) is closely related to the elastic deformation of the ligaments of bulk nanoporous Cu, as observed in

Fig. 7. The drop in stress in region I indicates the onset of damage to ligaments in bulk nanoporous Cu. When stress reaches the yield strength of these ligaments, bulk nanoporous Cu exhibits a yield-like behavior. There is always a long yield plateau due to the plastic collapse or brittle fracture of ligaments, marked as the stress plateau region (region II), where major energy absorption occurs. Evident serrations in this region also indicates that such bulk nanoporous Cu undergoes brittle fracture after the yield point

[49]

. As the compressive stress further increases,

bending, yielding, buckling of ligaments occur, leading to the rapid decrease of the inner pores in bulk nanoporous Cu. Consequently, the bulk nanoporous Cu deformed to densification and the stress-strain curves shows a sharp stress increase as shown in the region III. The Young’s modulus usually characterizes the deformation resistance of bulk nanoporous Cu. As previously discussed, elastic deformation occurs readily in bulk nanoporous Cu with higher porosity, as a result, Young’s modulus decreases with the increase of porosity of bulk nanoporous Cu. The Young’s modulus of bulk nanoporous Cu de-alloyed from Cu33Al67 and Cu30Al70 alloys shows a similar value due to little difference in their amount of porosity, between 72.6 % and 73.9 %. To analyze influence of porosity on the compression behavior of bulk nanoporous Cu de-alloyed from the four Cu-Al alloys, we plot Fig. 9(b) using the data from Fig. 9(a). The initial peak stress, plateau stress, and densification strain are typical variants in the three regions mentioned above. Based on the theory proposed by N. Gupta et al. [48], the points where the initial peak stress, plateau stress and densification strain values are selected or calculated are clearly marked in the insert figure of Fig. 9(a), as point 1, 2 and 3. The initial peak stress is the top stress in elastic region (region I), shown as the stress value of point 1, which is referred to as the compressive stress of bulk nanoporous Cu. The plateau stress value and densification strain value are respectively shown as the stress value of point 2 and the strain value of point 3. The plateau stress is equivalent to the lowest stress in plateau region (region II) while the densification strain

is the strain when bulk nanoporous Cu begin to be densified. In elastic region (region I) and plateau region (region II), elastic deformation occurs readily in bulk nanoporous Cu with higher porosity, as a result, the initial peak stress and plateau stress of bulk nanoporous Cu decrease with increasing porosity. In terms of the precise data, the initial peak stress decreases from 17.18 MPa to 2.71 MPa and the plateau stress decreases from 5.89 MPa to 0.57 MPa, with increasing porosity from 56.3 % to 73.9 % of as de-alloyed bulk nanoporous Cu. Some published studies [50, 51]

have shown the influence of porosity on the densification strain in the densification region

(region III) of traditional porous metals and suggest that densification strain shifts toward the higher strains with increasing porosity due to the decreasing of the mass of supporter in porous metals. Accordingly, the densification strain of bulk nanoporous shifts towards the higher strains (from 35.5 % for porosity of as de-alloyed Cu50Al50 sample to 46.3 % for porosity of as-de-alloyed Cu30Al70 sample) with increasing porosity due to the fact that the mass of Cu in bulk nanoporous Cu with a higher porosity is lower than that in bulk nanoporous Cu with lower porosity. Therefore, in region III, densification becomes difficult for bulk nanoporous Cu with a high porosity. The ligaments in bulk nanoporous Cu deform elastically since they are formed as a result of dissolution of Al atoms and rearrangement of Cu atoms, as illustrated in the de-alloying mechanism proposed in Fig. 8, thus leading to a low compressive strength (< 20 MPa). The bulk nanoporous Cu de-alloyed from Cu50Al50, Cu40Al60 and Cu33Al67 alloys, exhibits the same structure, that is bi-continuous interpenetrating ligament-channels with little difference in ligament size, thus the variation on the initial peak stress, plateau stress and densification strain of these bulk nanoporous Cu can be ascribed to variations in the volume fraction of porosity rather than to pore morphology. However, the microstructure of bulk nanoporous Cu de-alloyed from Cu30Al70, which is bimodal with nanopores (100’s nm) and interdispersed featureless regions a

few microns in size (see Fig. 7(d) and 7(h)), critically influences compressive behavior, particularly the strain required for densification. The presence of the micron-size interdispersed featureless regions in this material facilitates elastic deformation in both region I (elastic region) and region II (plateau region); this suggestion is consistent with the measured decrease in initial peak stress (2.71 MPa) and plateau stress (0.57 MPa), seen from Fig. 9(b). Moreover, the presence of these micron-size featureless regions result in a decrease in the volume fraction of nanoporous Cu, which will render densification more difficult (region III). It can be seen from Fig. 9(b) that the densification strain increases almost linearly in the case of bulk nanoporous Cu de-alloyed from Cu50Al50, Cu40Al60 and Cu33Al67, whereas it increases rapidly from Cu33Al67 to Cu30Al70; the maximum value of the densification strain, 46.3 %, is obtained for bulk nanoporous Cu de-alloyed from Cu30Al70.

5. Conclusions (1) Bulk nanoporous copper (6mm×6mm×3mm) with uniform pore morphology and high volume fraction of porosity (up to 73.9 %) is fabricated by chemical de-alloying of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 (at.%) alloys in 20 wt.% NaOH solution at temperature of 65 oC. (2) The compositions of the precursor Cu-Al alloy systems play an important role in the de-alloying process and have an important impact on the microstructure of the bulk nanoporous Cu. Bulk nanoporous Cu exhibit typical bicontinuous interpenetrating ligament-channel structure with ligament sizes of 110-150 nm (for Cu50Al50), 150-180 nm (for Cu40Al60) and 150-200 nm (for Cu33Al67). While the as-de-alloyed Cu30Al70 sample presents a novel bimodal microstructure with nanopores (100’s nm) and interdispersed featureless regions a few microns in size. (3) Bulk nanoporous Cu exhibits ductile-like behavior in compression and the compressive stress-strain curve corresponds to the typical stress-strain curve of porous metals. The

compressive strength increases with decreasing volume fraction of porosity; as porosity increases from 56.3 % to 73.9 %, the compressive strength decreases from 17.18 MPa to 2.71 MPa. Acknowledgements The work is financially supported by the National Natural Science Foundation of China (No. 51202171, No. 51472188, No. 51521001), the Specialized Research Fund for the Doctoral Program of Higher Education of China (No. 20120143120004) and the “111” project (No. B13035).

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Table 1

The phase contents and relative density of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70

alloys Table 2 The elemental contents of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 alloys obtained from EDS mapping observation in Fig. 3. Table 3 The Al and Cu contents of de-alloyed Cu-Al samples obtained from ICP tests. Fig. 1 (a) Cu-Al phase diagram showing the phase compositions of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 alloys; (b) XRD patterns of precursor Cu-Al alloys. Fig. 2

Back scattered electron (BSE) images showing the microstructure of the precursor Cu-Al

alloys: (a) Cu50Al50; (b) Cu40Al60; (c) Cu33Al67; (d) Cu30Al70. Insert in (b): the SEM image at a lower magnification. Insert in (d): the SEM image of α-Al (Cu)-Al2Cu at a higher magnification. Fig. 3

(a) EDS images of Cu50Al50; (b) EDS images of Cu40Al60; (c) EDS images of Cu33Al67; (d)

EDS images of Cu30Al70. Fig. 4

(a) SEM image showing the microstructure of the precursor Cu30Al70 alloy consisting of

Al2Cu and α-Al (Cu)-Al2Cu; (b) The lamellar eutectic α-Al (Cu)-Al2Cu; (c) EDS images of Cu30Al70; (d) EDS images of the higher magnification lamellar eutectic α-Al (Cu)-Al2Cu. Fig. 5 XRD patterns of bulk nanoporous Cu through chemical de-alloying of the Cu-Al alloys in the 20wt.% NaOH solution. Fig. 6 Fig. 7

Experimental and theoretical porosity of bulk nanoporous Cu. Section-view SEM images showing microstructure of bulk nanoporous Cu through

de-alloying of the precursor Cu-Al (a and e: Cu50Al50; b and f: Cu40Al60; c and g: Cu33Al67; d and h: Cu30Al70) alloys in the 20 wt.% NaOH solution. Fig. 8 The schematic illustration of the evolution of porosity during de-alloying: (a) alloy immersed in electrolyte; (b) dissolution of Al atoms in the first layer; (c) agglomeration of Cu atoms; (d) growth of clusters; (e) coarsening and passivation of clusters. Fig. 9 (a) Compressive stress-strain curves of bulk nanoporous Cu; (b) the relationship of stress and strain as a function of porosity of bulk nanoporous Cu; Insert in: (a) the division of three regions of compressive stress-strain curves of bulk nanoporous Cu and the initial peak stress, plateau stress and densification strain values of bulk nanoporous Cu.

Table 1 The phase contents and relative density of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 alloys Phase contents (wt.%) Relative density Calculated from phase diagram

Refined by Maud software (%)

AlCu

Al2Cu

α-Al (Cu)

AlCu

Al2Cu

α-Al (Cu)

Cu50Al50

100

/

/

100

/

/

98.8

Cu40Al60

44.69

55.31

/

45.83

54.17

/

98.5

Cu33Al67

/

100

/

/

100

/

98.9

Cu30Al70

/

93.24

6.76

/

92.58

7.42

98.5

Table 2 The elemental contents of Cu50Al50, Cu40Al60, Cu33Al67 and Cu30Al70 alloys obtained from EDS mapping observation in Fig. 3. Alloys Cu50Al50

Cu40Al60

Cu33Al67

Cu30Al70

Elements Al Cu Total amounts Al Cu Total amounts Al Cu Total amounts Al Cu Total amounts

Line style K-series

K-series

K-series

K-series

wt.% 30.24 69.76 100.00 38.78 61.22 100.00 47.28 52.72 100.00 51.15 48.85 100.00

at.% 50.52 49.48 100.00 59.86 40.14 100.00 67.87 32.13 100.00 71.15 28.85 100.00

Table 3 The Al and Cu contents of de-alloyed Cu-Al samples obtained from ICP tests. De-alloyed samples

Cu contents (wt.%)

Al contents (wt.%)

Cu50Al50

98.38

1.62

Cu40Al60

98.43

1.57

Cu33Al67

98.96

1.04

Cu30Al70

99.27

0.73