Au-group (Ag, Au) and Pt-group (Ni, Pd, Pt) metals have lower surface diffusion coefficients than Cu and are defined as LSD. The chemical composition has been designed based on the differences in diffusion coefficients, and the micro-alloying of 1 at % LSD metals with the Ti60Cu40 amorphous precursor alloy results in the formation of bi-continuous nanoporous copper (NPC) with finer nanoporous structure. LSD-stabilized NPCs have the smallest characteristic pore sizes of 7 nm and 6 nm after dealloying amorphous Ti60Cu39Pd1 and Ti60Cu39Pt1 precursor alloys, while NPC had a pore size of 39 nm after dealloying the amorphous Ti60Cu40 alloy. The refining factor increases approximately from 3.7 for Ti60Cu39Ag1 to 1780 for Ti60Cu39Pt1 precursors due to the dramatic decrease in the surface diffusivity during both preferential dissolution and rearrangement of Cu adatoms. The elaboration efficiencies of Ti60Cu40 alloy with addition of 1 at.% Pt-group elements are higher than those of Au-group elements. The homogeneous distribution of LSD elements in both the precursors and final stabilized NPCs played a key role in restriction of the long-distance diffusion of Cu adatoms. LSD-stabilized NPCs are able to have an ultrafine nanoporosity with a pore size almost one order smaller than that from LSD-free alloys.
- amorphous precursor alloys
- nanoporous copper
- surface diffusion
- elaboration of nanoporous copper
Nanoporous metals (NPMs), a representative type of nanostructured materials, possess intriguing properties to generate promising potentials for various important applications, including catalysis, sensors, actuators, fuel cells, micro-fluidic flow controllers, and so forth [1, 2, 3, 4]. NPMs with a variety of superior physical-chemical properties arisen from their unique pore structure, large specific surface area, and high electrical conductivity have attracted great interests to explore their electrocatalytic properties and greatly extend their potential applications in catalysts, electrochemical sensing, and energy systems . Dealloying primarily originated from the phenomenon of selective corrosion has attracted more attention recently because it has been regarded to be an effective method to fabricate NPMs with a three-dimensional bi-continuous interpenetrating ligament-channel structure at the nanometer scale [5, 6, 7, 8]. Nanoporous copper (NPC) is cost-effective and readily fabricated via dealloying process due to the high electrochemical stability. NPCs with different morphologies have been obtained from numerous binary alloy systems including Zn-Cu [7, 9], Mg-Cu , Al-Cu [7, 11, 12, 13, 14, 15, 16], Ni-Cu , Mn-Cu [17, 18], Zr-Cu , and Ti-Cu .
The characteristic pore sizes of NPCs obtained from the systems mentioned above are relatively larger than those of nanoporous gold (NPG) or nanoporous platinum (NP Pt) and change from few tenth nanometers to few hundredth nanometers, particularly for Zr-Cu system with a pore size of 500 nm . The pore sizes have a significant effect on the mechanical properties of NPMs. The smallest size of nanopores is of order of 3.5 nm for the Cu-Pt system . The yield strength of nanometer-sized NPG ligaments has been improved from ~880 MPa to 4.6 GPa as the NPGs’ pore size was refined from 50 nm to 10 nm . As has been reported, the rough NPCs have a relative lower yield strength (i.e., 128 ± 37 MPa with a ligament size of 135 ± 31 nm , 86 ± 10 MPa as the ligaments with a size of 300–500 nm of NPCs fabricated by one-step dealloying of the melt-spun Al-50 at.% Cu alloy ). It is thus of importance to fabricate NPMs with finer nanostructure with smaller pore size and ligament scales. On the other hand, the elaboration of NPCs is helpful for the enhancements of the catalytic performance and sensitivity for various gaseous phases or metallic ions. Effective ways to reduce the characteristic nanopore sizes have been reported to be dealloying at low temperature ; chemical composition design of the precursor alloys, for example, Ag-Au-Pt , Al-Pt-Au , Ti-Cu-Au [26, 27], Ti-Cu-Ag , Ti-Cu-Ni , and Ti-Cu-Pd/Pt ; and modification of the solution chemistry by using organic acids  and by introducing the macromolecules of polyvinylpyrrolidone [32, 33]. The chemical composition design is considered to be an effective way to change the surface diffusion and rearrangements of the adatoms of the noble elements in the precursor alloys since these noble elements take effect from inside to outside. However, the uniform nanoporosity of NPCs is of importance for enhancing the mechanical properties and catalytic performances. Final nanoporous structure is affected by many factors, such as the chemical compositions and initial microstructure of the precursor alloys, the solution chemistries of dealloying solutions, and the experimental conditions (i.e., temperatures, etc.) [12, 34, 35]. The crystalline precursor alloys are extensively dealloyed to prepare NPMs (i.e., coarsen crystalline Al-Cu alloys , nanocrystalline Ag-Au, Ag-Au-Pt alloys [22, 24]). While the intermetallic phases or secondary phases exist in the matrix, the final nanoporous structures inherit the characteristics of their initial microstructure of precursor alloys. The characteristics of the casting structures, the intermetallics, and the phase segregations typical of the crystalline alloys, such as Al-Cu [11, 12, 13, 14, 15, 16], Zn-Cu , Mg-Cu , Mn-Cu , Zr-Cu , and Ti-Cu [12, 13] alloys are inherited by the final porous structures. The kind and the chemical composition of the initial intermetallic phases and microstructure usually cause the formation of multi-modal nanoporous structures [9, 10, 11, 12, 13, 14, 15, 16, 17, 19]. On the contrary, the amorphous alloys exhibit many advantages, especially in their uniform distribution of the alloy constituents without the segregation in the chemical composition and heterogeneity in the microstructure . It is believed that the absence of grain boundary, the large-scaled phase segregations, and intermetallics of amorphous precursors contributes to high uniformity in NPMs. Amorphous alloys with disordered atomic-scale structure, an absence of the weak sites (i.e., grain/phase boundaries typical of crystalline materials) , are good to be precursor alloys for the fabrication of the uniform NPMs in some extent. The formation of highly uniform and ultrafine nanoporous structures has been realized for several amorphous alloy systems, including Mg-Cu-Y , Ti-Cu-Au [26, 27], Ti-Cu-Ag , and Ti-Cu-Ni  ternary alloys.
2. Materials and methods
On the basis of binary Ti60Cu40 alloy, ternary alloys with nominal compositions of Ti60Cu39M1 (M: Ni, Pd, Pt, Ag, Au) were designed and prepared by arc melting of high purity (purity > 99.99 mass%) of raw metals. The LSD elements were divided into two groups: G-I Au-group (Ag, Au) and G-II Pt-group (Ni, Pd, Pt). The surface diffusion coefficients have been reported to be 1.1 × 10−24 m s−2 for Pd and 3.6 × 10−26 m s−2 for Pt and 2.2 × 10−23 m s−2 for Au . The surface diffusion coefficient of Ag in a vacuum is two orders lower than that of Cu as reported [12, 39, 40]. That of Cu is 1.1 × 10−18 m s−2, more than two orders higher than LSD elements. The surface diffusivity of Ni adatoms in the electrolyte was one tenth of that of Cu adatoms . The ribbon samples were fabricated by melt spinning with a dimension of 20 μm in thickness and 2 mm in width. The starting Ti60Cu39M1 amorphous precursor alloys were treated under the free immersion condition for 43.2 ks in 0.03 M HF solution at 298 K. The detailed information of the experimental procedures has been supplied in the previous publications [27, 28, 29, 30]. An X-ray diffractometer (XRD, Rigaku 4200) was employed to identify the change in the lattice constants, crystalline states of dealloyed alloys and microstructure of precursor alloys and dealloyed alloys. The porous morphologies of the nanoporous NPC and LSD-stabilized samples were observed by a scanning electron microscope (JEOL, JIB-4610F). The chemical composition of the as-dealloyed samples was analyzed by an energy dispersive X-ray spectroscope (JEOL, JIB-4610F). Transmission electron microscopes (JEOL, HC2100, and ARM200) were used to observe the internal porous structure, the nanoporosity of dealloyed alloys, and crystalline characteristics of the LSD-stabilized Cu ligaments in high-resolution TEM modes. The nanoporosity was mainly analyzed on the basis of the average size of the nanopores and ligaments. The characteristic pore size of as-dealloyed alloys was confirmed by the single chord method for over 125 sites on SEM and TEM images.
3. Elaboration of nanoporous copper via the chemical composition design
3.1. Characteristics of LSD-added Ti-Cu precursor alloys and nanoporous counterparts
The XRD patterns of the as-spun Ti60Cu40, Ti60Cu39Ni1, Ti60Cu39Pd1, and Ti60Cu39Pt1 ribbons present one strong and broad diffraction peak at 41° and a weak diffraction peak at 70–75°, indicating an amorphous structure, as shown in Figure 1. Other LSD-substituted Ti60Cu40 precursor alloys were still amorphous states, indicating that the micro-alloying of 1 at.% LSD elements had not generated the heterogeneous microstructure or phase segregation. The high-resolution TEM images (HRTEM) and corresponding selective area diffraction patterns (SADP) of as-spun Ti60Cu40 and Ti60Cu39M1 show that as-spun LSD-substituted Ti60Cu39M1 alloys had an amorphous structure without the crystalline clusters and phase segregations . As is well known, because G-I LSD elements, Ag and Au, have a similar crystal structure to Cu, the addition of minor amounts of G-I elements has thus no significant influence on the structure of as-spun alloys [26, 27]. In G-II elements, Ni, Pd, and Pt, three strong diffraction peaks were observed in the XRD patterns of the dealloyed Ti60Cu39Pd1 and Ti60Cu39Pt1 ribbons after dealloying for 43.2 ks in the 0.03 M HF solution. The diffraction peaks were identified to be fcc Cu and had slightly shifted to low diffraction angles due to the expansion of the LSD-substituted unit cells. The small peak around 36° was considered to be from the Cu2O phase. The absence of strong diffraction peaks from crystalline Pd and Pt phases indicates that CuPd or CuPt solid solution formed after dealloying. These XRD patterns were similar to those of Ag-, Au-, and Ni-stabilized NPCs [26, 27, 28, 29]. Figure 2 shows the change in the lattice constant, α, and the grain sizes, L, with 1 at% addition of Au-group (G-I) and Pt-group (G-II). The standard lattice constants are reported to be 0.3524 nm for Ni, 0.3608 nm for Cu, 0.3891 nm for Pd, 0.3924 nm for Pt, 0.4078 nm for Au, and 0.4085 nm for Ag, respectively. The lattice constants of NPCs were estimated to be 0.3615–0.3627 nm on the basis of the XRD data of Cu (111) peaks. The lattice constants became larger when the added elements had larger lattice constants except in the case of Ag. As indicated in Figure 2, the Cu lattice expanded more when Pd and Pt were micro-alloyed into the Ti60Cu40 alloy. The added Ni, Pd, or Pt atoms were thus considered to substitute the Cu atoms in the Cu lattice to a higher extent than in the Ag- or Au-added cases, resulting in the expansion of the Cu lattice constants from 0.3615 nm to 0.3627 nm in Figure 2. In previous works it was found that the Au or Ag phase was formed after dealloying the Ti60Cu39Au1 and Ti60Cu39Au1 precursor alloys [26, 27, 28, 29]. The absence of diffraction peaks from the Pd or Pt phase in the XRD patterns of dealloyed Ti60Cu39Pd1 and Ti60Cu39Pt1 ribbons and the large expansion of the lattice constants of the NPCs indicated that more Pt-group atoms invaded the Cu lattice than Au-group atoms (i.e., Ag addition and Au addition) for NPCs from precursor alloys micro-alloyed with LSD elements [26, 27, 28, 29]. On the other hand, the grain sizes were reduced to 15 nm and 13 nm when the Pd and Pt were added into Ti60Cu40 ribbons. The decrease in the grain sizes was considered to be due to the retardation of the self-diffusion of Cu adatoms by LSD elements [27, 29, 30]. The diffusion distance of Cu adatoms under free diffusion patterns is prevailed in a long-distance diffusion mode [2, 5, 10, 19, 20, 21]. However, the long-distance self-diffusion of Cu adatoms was interrupted and restricted by the LSD adatoms during the rearrangement of adatoms and resulted in an accumulation of Cu and LSD adatoms in a smaller scale. Consequently, smaller grains were formed when the Pt-group elements (Pd, Pt) and Au element were used to stabilize NPCs.
The nanoporous surface morphologies of dealloyed Ti60Cu39M1 (M: Cu, Ag, Au, Ni, Pd, Pt) alloys are shown in Figure 3. The mean pore size of the nanoporous structure of as-dealloyed Ti60Cu40 ribbons was about 71 nm after dealloying in 0.03 M HF solution. The characteristic length scales of ligaments of NPC obtained from Ti60Cu40 alloy were 74 nm in 0.03 M HF solution. The characteristic sizes for the nanopores and ligaments were confirmed to be for 41 and 48 nm for dealloyed Ti60Cu39Ag1 precursor, 16 nm and 27 nm for Ti60Cu39Au1 precursor, 12 nm and 26 nm for Ti60Cu39Ni1 precursor, 9 nm and 24 nm for Ti60Cu39Pd1 precursor, and 8.5 nm and 31 nm for Ti60Cu39Pt1 precursor. Figure 4 shows a typical bright field TEM image (BFI), their corresponding selected area diffraction pattern (SADP) of Ti60Cu39M1 (M: Cu, Ag, Au, Ni, Pd, Pt) amorphous alloys after dealloying for 43.2 ks. A bi-continuous porous microstructure was formed with a characteristic pore size of 7 nm for the dealloyed Ti60Cu39Pd1 ribbon and ca. 6 nm for the dealloyed Ti60Cu39Pt1 ribbon, respectively (Figure 4e and 3f). The diffraction rings in the SAD patterns were assigned to Cu (111), (200), (220), and (311) (JCPDS card No. 02-1225), and the inner ring was done for Cu2O (111) (JCPDS card No. 74-1230). The diffraction rings from the Cu2O phase were absent in the inset SADP in Figure 4e. The ligament sizes were confirmed to be ca. 23 nm for the dealloyed Ti60Cu39Pd1 ribbon and ca. 30 nm for the dealloyed Ti60Cu39Pt1 ribbon. The BFI images mainly reflected the internal nanoporous structures with a finer porosity in comparison to those in the surface regions. It is considered that the finer nanoporous structure in the internal parts mainly resulted from the lower concentration gradients inside the channels. As shown in Figure 4e and f, the pore size of Pd-stabilized and Pt-stabilized NPCs had a mean pore size of 7 and 6 nm, respectively. As shown in Figure 4b–d, the characteristic pore size of NPCs stabilized by the micro-alloying of Ni, Ag, and Au [27, 28, 29] has been reported to be 11 nm, 28 nm, and 12 nm on the basis of TEM analysis, respectively. The mean pore size of dealloyed Ti60Cu40 ribbons confirmed by TEM observation was 39 nm in Figure 4a . The pore size decreased more than one order due to the addition of either Pd or Pt as shown in Figure 4e and f. The characteristic scale length of the nanopores and ligaments of dealloyed Ti60Cu39Pd1 and Ti60Cu39Pt1 ribbons decreased dramatically. On the basis of XRD, TEM, and SEM-EDX analysis, the residue is considered to be fcc CuPd and fcc CuPt solid solution. The similarity existed in all LSD-substituted Ti60Cu39M1 (M: Ag, Au, Ni, Pd, Pt). Commonly the Cu(LSD) solid solutions can be regarded as the residual phases after dealloying. However, the selective dissolution is slightly different which caused the formation of Ag and Au phases [27, 28]. Although NPCs from Ti60Cu39Ag1 precursor had a large final nanopore, the elaborating behavior still happened at the initial dealloying stages with a high refining efficiency . As shown in Figure 5, the distribution of Au LSD element was profiled. The uniform distribution of Au elements can be confirmed here. For other precursor alloys, the residual LSD elements had similar profiles which benefited from the uniform distribution of these LSDs in the amorphous precursor alloys.
3.2. Effects of LSDs on surface diffusion
On the basis of the surface diffusion-controlled coarsening mechanism, the surface diffusivity, Ds, at various dealloying temperatures was estimated by the equation :
where k is Boltzmann constant (1.3806 × 10−23 J K−1), γ is surface energy, t is the dealloying time (43,200 s), d(t) is the pore size at t, T is the temperature, and α is the lattice constant. The pore size of NPCs confirmed by TEM micrographs and lattice constants calculated from XRD data were adopted for calculation of Ds. The surface energy of Cu has been reported to be 1.79 J m−2 [40, 43]. The surface energy of micro-alloyed elements has been reported to be 1.24 J m−2 for Ag, 1.50 J m−2 for Au, 2.0 J m−2 for Ni, 2.0 J m−2 for Pd, and 2.49 J m−2 for Pt [40, 43]. The concentration of Au-group elements (Ag, Au) and Pt-group elements (Ni, Pd, Pt) in the precursor alloys was 1 at%. The concentration of added elements (Ag, Au, Ni, Pd, Pt) in NPCs after dealloying should theoretically be 2.5 at% if the dissolution of Cu in HF solution is not considered. Therefore, the surface energy of LSD-stabilized NPCs is considered to be very close to that of Cu. The surface energy of Cu, 1.79 J m−2, was adopted for the calculation of Ds. The characteristic nanopore sizes were summarized in Figure 6a, which has been illustrated above. On the other hand, the surface diffusivities of these LSD-stabilized alloys in 0.03 M HF solutions can be modified. As shown in Figure 5b, the value of Ds was estimated to be 2.5 × 10−18 m2 s−1 for Ti60Cu40 ribbon, 6.0 × 10−20 m2 s−1 for Ti60Cu40 ribbon, 9.7 × 10−21 m2 s−1 for Ti60Cu39Au1 ribbon and 2.0 × 10−21 m2 s−1 for Ti60Cu39Ni1 ribbon, 2.3 × 10−22 m2 s−1 for Ti60Cu39Pd1 ribbon, and 1.3 × 10−22 m2 s−1 for Ti60Cu39Pt1 ribbon, respectively. The surface diffusivity decreased more than four orders of magnitude due to the addition of Pd and Pt. Compared with Ds of the Au-group-stabilized NPCs, the decrease in Ds in the Pt-group-stabilized NPCs was remarkable.
The refining factor, R, is defined as the ratio between the surface diffusivity, Ds, of NPC obtained from amorphous Ti60Cu40 precursor (numerator) and NPCs stabilized by LSDs (denominator):
As shown in Figure 6c, the estimated refining factor was 968 for Ti60Cu39Pd1 ribbon and 1780 for Ti60Cu39Pt1 ribbon, respectively. In other cases, the value of R changed from 4 for the dealloyed Ti60Cu39Ag1 ribbon, 111 for the dealloyed Ti60Cu39Ni1 ribbon, and 231 for the dealloyed Ti60Cu39Au1 ribbon, respectively. The refining efficiency of the micro-alloying of the Pt-group elements into the precursor Ti60Cu40 alloy was one order higher than that of the Au-group cases. The lower value of R for Ti60Cu40 alloy stabilized by Au-group (G-I) elements could take into account from several aspects: (1) the larger self-diffusion coefficients of Ag and Au than that of Cu [44, 45, 46, 47, 48] and (2) the large misfit in the atomic radii . As has been reported before [44, 45, 46, 47, 48], the self-diffusion coefficients of Au- and Pt-group elements are ranked as following: Ag > Cu > Au > Ni > Pd > Pt. As such, there are more possibilities for Ag adatoms to meet with other Ag adatoms during the diffusion and rearrangements processes of dealloying to form clusters and to develop the Ag phase because it has a larger self-diffusion coefficient than Cu atoms [44, 48]. The self-diffusion coefficient of Au is slightly smaller than that of Cu . Because the misfit in the atomic radius between Cu and Ag and Cu and Au is about 13%, the incorporation of Ag-Au atoms into the Cu lattice becomes more difficult than that of Cu/Cu atoms. On the other hand, the self-diffusion coefficients of Pt-group (G-II) elements are more than two orders lower than that of Cu element [47, 48], and the formation of the crystalline phase of G-II elements (Ni, Pd, Pt) was interrupted by the fast diffusion of Cu adatoms. Furthermore, the misfits in the atomic radius between Cu and Pd and Pt elements are ca. 8%, and the incorporation of Pd/Pt into the Cu lattices appears more easily. The changes in the lattice constants of the NPCs stabilized by G-I and G-II indicated by XRD patterns in Figure 1 and Refs. [27, 28] support the present hypothesis. As has presented above, the formation of the Ag and Au crystalline phases after dealloying is considered to be due to the difference in self-diffusion behaviors and misfits of the atomic radius. The minor addition of Group-I LSDs causes the formation of the crystalline Ag/Au phases, and the refining factors of Group I were small. On the other hand, the micro-alloying by Group-II LSDs mainly resulted in the invasion of the Cu lattices, with some Cu atoms in the lattice substituted by Ni, Pd, and Pt atoms, forming solid solutions, and the refining factors for Group II were high. In some extent, these LSD atoms in the lattice are considered to be the main contributors for the elaboration of NPCs.
The diffusion in the interfacial regions between Cu, LSD adatoms, and Ti adatoms played a key role in the formation of ultrafine NPCs. For example, Evangelakis investigated the diffusion of Au adatoms on Cu and self-diffusion of Cu adatoms by using the molecular dynamic method . They have found that diffusion of Cu adatoms takes place exclusively by hopping from one adatom position to an adjacent one and that multiple jumps are frequent at low temperatures. On the other hand, Au adatoms on the NPC ligaments hopped less frequently than in the case of Cu self-diffusion. The migration energy required for hopping of Au on Cu was almost twice that of the corresponding energy for Cu adatoms. While the diffusion of Au on Cu was more difficult than the diffusion of Cu on Cu, this phenomenon is compatible with the binding energy of Au adatoms, which is found to be 2.77 eV for Au adatoms and 2.26 eV for Cu adatoms. The Au adatoms relaxed at a distance −15% smaller than the bulk interlayer distance, but in the case of the Cu adatoms, it was −10% for the same quantity . The Cu adatoms diffused quickly in and out by themselves; however, the Au adatoms at the activated sites diffused out slowly and hardly ever diffused back to the NPC ligaments. Consequently, Au adatoms gradually accumulated outside the ligaments during the dealloying [27, 30]. As a result, Au adatoms uniformly distributed outside the ligaments after dealloying for 43.2 ks via the hopping mechanism, as shown in Figure 5. When more Au was added to TiCu amorphous alloys, the surface coverage of Au adatoms increased during dealloying, suggesting that the accumulation rate of Au adatoms in the more concentrated solutions was higher due to the faster migration of Au and Cu adatoms . Because the accumulated Au adatoms built up a continuous outmost diffusion barrier, the behaviors of diffusion and rearrangements of Cu and Au adatoms can be fulfilled to form ultrafine NPCs from both the LSD-substituted Ti60Cu40 amorphous precursor alloys in HF solutions. The diffusion behavior for other LSD-stabilized nanoporous structures was considered to be similar. However, the accumulation of G-II Pt-group elements (Ni, Pd, and Pt) with much slower diffusion rates and Ni, Pd, and Pt cluster or grains with much smaller size less than 7 nm tended to form, which are not sensitive to X-ray . As has been reported , a bimodal nanoporous structure with a pore size of 10 nm and 20 nm has been fabricated from Al75Pd17.5Au7.5 precursor alloys by successive dealloying. The initial heterogeneous microstructure consisting of Al2Au- and Al3Pd-type intermetallics causes the formation of a bimodal nanoporous structure. The similar evolution of bimodal or multimodal nanoporosity on precursor alloys heterogeneous in microstructure has been reported [11, 12, 14, 20, 41]. Amorphous precursor alloys with homogeneously distributed Cu, Ti, and LSD atoms also have an important refining effect. The uniform distribution of added Pd and Pt atoms in final NPCs resulted in smaller nanopores, as shown in Figures 3 and 4. In 2008, fine nanoporous AuPt alloys with a pore size of about 5 nm were fabricated from an Ag65Au29Pt6 precursor . The alloying of 6 at% Pt into the Ag65Au35 alloy effectively reduced the pore size from 10 to 20 nm to about 4 nm, which also supports the present results. However, the high cost of Au and Pt weakens their potential application. By minor addition of the 1 at% Pt-group elements, it is possible to elaborate NPCs down to a pore size of approximately 6 nm, comparable to high-cost AuPt nanoporous in nanoporosity. So far the LSD-stabilized NPCs had a relative small nanopores and narrow ligaments comparable to these of NP Au, NP Pd, and NP Pt. If catalytic Au, Pd, or Pt monolayer is electrodeposited on ultrafine NPCs, the catalytic performance of this kind of cost-effective porous materials is able to be close to those Au, Pd, or Pt catalysts. Meanwhile, if the oxides are assembled on ultrafine NPC templates, the new-developed nanocomposites can be used as promising pseudocapacitors to contribute in the energy conversion fields.
The 1 at% minor addition of low surface diffusive (LSD) elements in two groups (G-I, Au-group (Ag, Au), and G-II, Pt-group (Ni, Pd, Pt)) is able to elaborate the nanoporous Cu structure dealloyed from micro-alloyed Ti60Cu40 alloy efficiently. The chemical compositions of Ti60Cu40 alloys stabilized by the addition of G-II metals were shown to be more effective in refining nanoporous structure than the addition of G-I metals. Nanoporous Cu with a pore size of less than 7 nm was obtained from Ti60Cu39Pd1 and Ti60Cu39Pt1 ribbons after dealloying. The residue of dealloyed Ti60Cu39Pd1 and Ti60Cu39Pt1 ribbons was confirmed to be bimetallic solid solutions, such as fcc Cu(Pd) or Cu(Pt) solid solutions, and fcc Cu(Au), Cu(Ag), and Cu(Ni) solid solutions formed after dealloying. The refining factor increases approximately from 3.7 for the Ti60Cu39Ag1 precursor alloy to 1780 for the Ti60Cu39Pt1 precursor alloy. The elaboration was attributed to the dramatic decrease in the surface diffusivity during both preferential dissolution and rearrangement of Cu adatoms. The refinement efficiency of the micro-alloying of the G-II LSDs in Pt-group elements was almost one order higher than that of the G-I LSDs in Au-group elements. The homogeneous distribution of LSD elements in both of the amorphous precursor alloys and the final stabilized NPCs played a key role in refining the NPCs. The strategy outlined in this work has the potential to be applied to other alloy systems to obtain other ultrafine nanoporous metals with comparable nanoporosity to those high-cost catalysts.
The authors gratefully acknowledge the financial support from the Ministry of Education, Culture, Sports, Science, and Technology (MEXT) Japan through Grant-In-Aid for Science Research in a Priority Area on “Research and Development Project on Advanced Materials Development and Integration of Novel Structured Metallic and Inorganic Materials” and a Grant-in-Aid for Young Scientists (B) under Grant No. 24760567. This work is also financially supported by the Natural Science Foundation of China under Grant no. 51671106 and Natural Science Foundation of Jiangsu Province under Grants nos. BK20171424 and BK20151536. The authors would also like to acknowledge Jiangsu Collaborative Innovation Center for Advanced Inorganic Function Composites, the International S&T Cooperation Program of China (2015DFA51430).
Conflict of interest
The authors declare no conflict of interest.