Dealloying of Quasi-High Entropy Alloys: Fabrication of Porous Noble Metals/Metal Oxides

Abstract

High entropy alloys (HEAs) have been widely studied due to their special crystal structure, but their bulk structure and low specific surface area limit their further application in broader fields. In this work, the dealloying of precious metal Cu₃₅Pd₃₅Ni₂₅Ag₅ quasi-HEAs is performed. Porous noble metals with micro prism array structure and porous noble metal PdO/Ag₂O/NiO oxides with nano “ligament/pore” structure are obtained by constant potential dealloying and free dealloying, respectively. In this way, the porosification of quasi-HEAs and noble metal oxides is achieved. Moreover, the effects of dealloying parameters on pore morphology and phase structure of dealloyed materials are studied, and the evolution mechanisms of pore structures of different dealloying products are discussed. The work provides strategies for the preparation of porous precious metal quasi-HEAs and porous noble metal oxides by the dealloying method. These products present great potential for application as functional materials in hot fields such as catalysis and energy storage.

1. Introduction

High entropy alloys (HEAs) exhibit excellent mechanical, physical, and chemical properties, such as high strength, fracture resistance, high hardness, ductility, wear resistance, and corrosion resistance [1,2,3,4,5,6,7,8]. Currently, they are mostly used as structural materials. In addition, they also have enormous potential in thermal, electrical, magnetic, catalytic, energy storage, and other aspects [9,10,11,12,13,14]. The traditional alloy design concept [15,16] believes that the more components an alloy involves, the more likely it is to produce hard and brittle intermetallics with complex crystal structures. The appearance of these phases will, in most cases, deteriorate the mechanical properties of the material. Compared with traditional alloys, HEAs have a relatively short research history. With the expansion of the research scope of HEAs, the concept of HEAs has evolved from equimolar single-phase solid solution to heterogeneous solid solution with unequal molar ratios [17]. The mechanical properties of HEAs, such as wear resistance, fatigue resistance, high strength, and the combination of ductility and high fracture toughness, have become a research hotspot [17,18,19,20,21,22,23]. It is worth noticing that most of the reported HEAs currently contain transition metals [24,25,26], such as Ni, Fe, Co, Pd, etc. For example, FeCoNiCrTi_0.2 HEA has good ductility at room temperature and low temperature [27]. The strength-to-weight ratio of the nanocrystalline HEA Al₂₀Li₂₀Mg₁₀Sc₂₀Ti₃₀ is roughly equivalent to that of ceramics, and it still maintains high hardness after annealing [28]. HEAs have received widespread attention as structural materials [29,30,31,32]. On the one hand, due to the limitations of preparation technology, HEAs are mostly prepared into bulk materials, resulting in low specific surface area of the materials and greatly limiting their use as functional materials. On the other hand, very few precious metal HEAs have been reported. In this situation, there is an urgent need to develop new types of precious metal HEAs.

The commonly used preparation methods for HEAs include arc melting [33], mechanical alloying [34], laser cladding [35], and so on. These traditional methods mostly focus on preparing large-sized HEAs, such as CoCrFeMnNi HEA [36], which can be prepared through mechanical alloying and spark plasma sintering, making it difficult to reduce its size to micrometer or even nanoscale [37,38,39]. At present, there are many preparation strategies for HEAs, but there is a lack of post-processing technology to utilize HEAs as intermediates to further synthesize other new materials. Carbon thermal shock technology [40] and MOF-template method [41] were used to synthesize HEA nanoparticles, which show excellent catalytic effects. For example, the five-component PtPdRhRuCe HEA nanoparticles prepared by carbon thermal shock technology were used as an ammonia oxidation catalyst, which receives a conversion rate of close to 100% and selectivity to reactants of more than 99% [40]. However, the above processes are more complex and costly.

In the current work, a low-cost and simple process for dealloying a kind of quasi-HEAs was developed. Porous quasi-HEAs with micro prism array structure and nanoporous PdO/Ag₂O/NiO composites were successfully prepared under different process parameters through constant potential dealloying and free dealloying [42], respectively. At the same time, this work investigated the influence of dealloying parameters on the pore morphology and phase structure of the dealloyed materials and revealed the evolution mechanism of different pore structures. The processing technology used in this study provides a new approach for preparing more porous materials and HEA-derived materials and provides assistance for the development of related materials in catalysis and energy storage fields.

2. Materials and Methods

Cu, Pd, Ni, and Ag ingots with a purity of 99.99 wt.% (purchased from China New Metal Materials Technology Co., Ltd., Beijing, China) were selected as raw materials. The Cu₃₅Pd₃₅Ni₂₅Ag₅ alloy ingots were first fabricated by the previously reported arc melting method [43]. Then, the Cu₃₅Pd₃₅Ni₂₅Ag₅ quasi-HEA ribbons were obtained by the melt-spinning method (with a spraying pressure of 0.1 MPa and a copper roller speed of 2000 r/min). The ribbons present a width of about 1.5 mm and a thickness of about 20~30 μm.

Table 1. List of phase structure, atomic size difference (δ), enthalpy of mixing (∆H_mix), dimensionless parameter (Ω), and the corresponding references.

Alloy	Phase	δ	∆Hmix (KJ·mol−1)	Ω	VEC	Ref.
AuPdAgPt	single fcc	2.2	−2.0	7.9	5.5	[21]
AuPdAgPtCu	dual fcc	4.4	−6.6	3.0	4.6	[21]
AuPdAgPtCuNi	dual fcc	5.4	−2.2	8.8	5.5	[21]
Fe50Mn30Co10Cr10	fcc + hcp	0.6	−0.9	15.6	7.6	[20]
Cu35Pd35Ni25Ag5	dual fcc	4.7	−5.1	2.7	6.4	This work

lists the phase structures and atomic differences (δ), entropy of mixing (∆S_mix), enthalpy of mixing (∆H_mix), dimensionless parameter (Ω), and corresponding references of representative (quasi-)HEAs. The calculation method of the parameters is presented in Supporting S1 and Table S1. As with traditional HEAs, the values of δ, ∆H_mix, and Ω of Cu₃₅Pd₃₅Ni₂₅Ag₅ completely conform to the formation law of solid solution phase (−15 KJ·mol⁻¹ ≤ ∆H_mix ≤ 5 KJ·mol⁻¹, δ ≤ 6.6, and Ω ≥ 1.1). However, the mixing entropy of Cu₃₅Pd₃₅Ni₂₅Ag₅ (∆S_mix = 1.23R) is smaller than the lowest conformational entropy of 1.39 R for high-entropy alloys. To describe it more accurately, we define the current material as a quasi-HEA.

Electrochemical dealloying was performed through a three-electrode system at 298 K by an electrochemical workstation (CHI 660E). The Cu₃₅Pd₃₅Ni₂₅Ag₅ ribbons for testing are 3 cm long. The electrolyte is solution L_a (

Table 2. Two types of corrosive liquids required for this work.

Solution Number	Composition
La	1 M HNO3 and 0.5 M H2SO4
Lb	3 M HCl and 1 M HNO3

), while the potentials are 600 mV, 750 mV, 850 mV, and 950 mV, respectively. The working electrode is the Cu₃₅Pd₃₅Ni₂₅Ag₅ ribbon, the counter electrode is platinum foil, and the reference electrode is Ag/AgCl (saturated NaCl). Free dealloying was carried out by immersing Cu₃₅Pd₃₅Ni₂₅Ag₅ ribbons in 20 mL L_b (Table 2) solution at 298 K. The sample was taken out after dealloying and cleaned with deionized water three times. Then, it was placed in a vacuum-drying oven for further tests.

The phase compositions of the samples were confirmed by X-ray diffractometer (XRD, D8 Advance, Bruker, Karlsruhe, Germany) using Cu K_α radiation (λ = 1.5406 Å) with a scanning range of 30° to 90° at a scan rate of 5°/min. The surface morphology of the samples was observed using scanning electron microscopy (SEM, JSM 6700F, JEOL, Tokyo, Japan), and the chemical composition of the samples was analyzed using built-in energy dispersive spectroscopy (EDS, Ultim MaxX, Oxford, Oxfordshire, UK). The ligament/pore structure of the samples was observed using transmission electron microscopy (TEM, JEM-2010FEF, JEOL, Tokyo, Japan). The elemental composition and valence states of the sample surface were determined using X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi, Thermo Fisher Scientific, Waltham, MA, USA). The specific surface area and pore size distributions of samples were determined by using the Brunauer–Emmett–Teller (BET) and Barrett–Joyner–Halenda (BJH) methods, respectively.

3. Results

3.1. Cu₃₅Pd₃₅Ni₂₅Ag₅ Ribbons

Figure 1 shows the XRD spectra of the Cu₃₅Pd₃₅Ni₂₅Ag₅ ribbon. From Figure 1a, it can be seen that the ribbon exhibits a dual-phase fcc structure (fcc 1 phase and fcc 2 phase) [44]. Figure 1b,c shows the enlarged XRD spectra of (200) and (311) crystal planes, respectively. From Figure 1b,c, it can be seen that the fcc 2 phases, as the shoulder peak of the fcc 1 phases, are located at a slightly higher Bragg angle, indicating that the lattice parameters of the fcc 2 phase are slightly smaller than those of the fcc 1 phase.

The following is known:

2d_hklsinθ = λ

(1)

$${\mathrm{d}}_{\mathrm{hkl}}=\mathsf{\alpha }/\sqrt{h^2+k^2+l^2}$$

(2)

Combining Equations (1) and (2), the lattice constants (α) of the two phases can be quantitatively described as follows:

$$\mathsf{\alpha }=\mathsf{\lambda }\sqrt{h^2+k^2+l^2}/2\sin \mathsf{\theta }$$

(3)

According to the diffraction peak positions corresponding to the (200) crystal plane, the lattice parameters of the fcc 1 phase and fcc 2 phase are calculated to be 0.2053 nm and 0.2049 nm, respectively, indicating that the difference in lattice constants between the two phases is very small.

SEM images of the Cu₃₅Pd₃₅Ni₂₅Ag₅ ribbon shown in Figure 2a confirm the dual phase structure of the quasi-HEA. It can be inferred that the fcc 1 phase corresponds to the matrix phase with a grain size of 0.5~1.5 μm, and the fcc 2 phase relates to uniformly dispersed nanoparticles in the range of 50~100 nm (Figure 2a). In order to further reveal the distribution of various elements in the two phases, surface scanning energy spectrum analysis was performed on the area shown in Figure 2a, and the results are shown in Figure 2b–e. From the signal distribution maps of Cu-K, Pd-K, Ni-K, and Ag-K shown in the figure, it can be seen that there is no obvious element segregation or aggregation phenomenon in the selected area, indicating that the two phases are highly similar in chemical composition. This is consistent with the conclusion of the lattice constant analysis above. Combined with the analysis of XRD and SEM, it can be concluded that the precursor ribbon is a homogeneous two-phase solid solution quasi-HEA.

3.2. Micron Prism Array Structure

In the process of constant potential corrosion, the selection of constant potential plays a crucial role in the formation of dealloyed products [45,46]. Figure 3 shows the open circuit potential–time curve of ribbon constituent elements in solution L_a. From Figure 3, it can clearly be seen that the open circuit potentials of Ni, Cu, Ag, and Pd elements in the solution L_a are −95 mV, −55 mV, 220 mV, and 490 mV, respectively. When the dealloying potential is between −95 mV and 490 mV, certain elements in the quasi-HEA ribbon will be selectively removed to obtain porous precious metals. When the dealloying potential is greater than 490 mV, the quasi-HEA ribbon will be etched as a whole to obtain porous HEA. In order to maintain the solid solution phase structure of the precursor and achieve porous materials, the dealloying potential applied in this experiment is greater than 490 mV. In order to seek the optimal dealloying potential and study the effect of dealloying potential on the pore morphology of as-corroded samples, the precursor ribbons were dealloyed at potentials of 600 mV, 750 mV, 850 mV, and 950 mV for 3 h in this experiment.

Figure 4 shows the XRD detection results of the precursor ribbon after dealloying at a potential of 850 mV in solution L_a for 3 h. Figure 4a shows that the peak shape of XRD of the precursor HEA ribbon before and after dealloying are similar. This indicates that the dealloyed sample retains the two-phase solid solution phase structure. The difference is that the intensity of the diffraction peak of the dealloyed sample is significantly reduced compared to the precursor. This phenomenon may be related to the filtration of some elements and the formation of microporous structures in the as-corroded samples. In addition, compared to the XRD peak position of the precursor ribbon, the XRD peak position of the dealloyed sample shifts to a lower Bragg angle by approximately 0.38° (Figure 4b,c). This phenomenon is attributed to the etching of more Cu elements in the solid solution phase.

Figure 5 shows the SEM cross-sectional view of the quasi-HEA ribbon after dealloying at the potential of 850 mV in solution L_a for 3 h. From Figure 5a, it can be seen that the dealloyed sample exhibits a uniform porous prismatic structure with a small number of microparticles attached to the surface. Figure 5b is the high-power SEM image of the prism structure. The diameter of the prism is approximately 0.3~0.5 μm. An open pore structure is formed between independent prisms, and the particle surface also exhibits an open, bicontinuous nano ligament/pore structure (Figure 6c) [47]. Figure 5d is the cross-sectional view of the dealloyed sample. The prism runs through the entire ribbon and is arranged in an array with a length of approximately 8.3 μm (much less than the thickness of the precursor ribbon). This is because, at a high dealloying potential, some elements are corroded away [48]. In addition, the diameter at the top of the prism is significantly smaller than the diameter at the bottom. This is due to the large temperature gradient generated during the cooling process, which causes the grain size on one side close to the copper roller to be significantly smaller than the other side, thus forming a prism structure with continuous changes in diameter in the cross-sectional direction. From Figure 5e–i, it can be seen that the elements in the matrix are uniformly distributed, while the particle composition is rich in Ag. The structure of the quasi-HEA changes from the dual-phase solid solution structure (before corrosion) to the composite structure of Ag-rich particles attached micro prism array (after dealloying). This result is consistent with the change of the double peaks in the XRD of Figure 4.

In this work, both SEM and XRD data confirm the existence of the dual-phase FCC structure. For XRD patterns, all the constituent elements of Cu₃₅Pd₃₅Ni₂₅Ag₅, including Cu, Pd, Ni, and Ag, are FCC structures generally (Figure 1). In addition, next to all the main peaks in XRD patterns, corresponding satellite peaks appear (Figure 1b,c and Figure 4b,c). This phenomenon is attributed to the presence of two phases with the same phase structure but differ in composition. For SEM and EDS results, it can be clearly seen that there are a large number of precipitated phase particles on the surface of the precursor matrix (Figure 2a), which confirms that the precursor is composed of two kinds of phases. During the dealloying, the matrix phase and precipitate phase are etched, respectively (Figure 5a). The element mappings reveal the difference in the composition of the two phases, in which the content of Ag in the precipitate phase is significantly higher than that in the matrix phase (Figure 5e–i). From the above discussion, we can see that the phase structure influences the dealloying process. The dealloying products often inherit the phase structure characteristics of the precursor. So, even in the same alloy system, different composition ratios and/or phase structures will directly affect the selectivity, as well as the morphology and composition of as-obtained products during the dealloying process. In this study, the matrix phase and precipitate phase are etched to form micro prism arrays and porous particles, respectively.

To study the effect of dealloying potential on the porous morphology of dealloyed samples, the precursor quasi-HEA ribbons were dealloyed at potentials of 600 mV, 750 mV, 850 mV, and 950 mV for 3 h, respectively. Figure 6 shows the SEM plane views of the surface microstructure of the sample after dealloying. The results indicate that the dealloying potential significantly affects the microstructure of the corroded sample. When the dealloying potential is low at 600 mV (Figure 6a), only a few inconspicuous microcracks appear at the grain boundaries. This indicates that the dealloying does not obtain sufficient driving force at low potential. When the dealloying potential increases to 750 mV (Figure 6b), cracks at grain boundaries become apparent, but an impermeable pore structure still exhibits. When the dealloying potential further increases to 850 mV (Figure 6c), the corroded sample displays a micro prism array structure. In this situation, the porosity of the sample is greatly improved. This greatly increases the specific surface area of the material. When the dealloying potential increases to 950 mV (Figure 6d), the dealloyed sample presents an irregular prism morphology, and a large number of small pores appear on the micron prisms. It is speculated that this phenomenon may be due to uneven corrosion of the alloy at a high potential, leading to local phase detachment, which greatly affects the regularity of the pore morphology. Figure 6e,f shows the bar charts of the relationship between the size of ligament/pore and the dealloying potential. The measurement method of ligament and pore size is presented in Supporting S2 and Figure S1. It can be seen that as the dealloying potential increases, the ligament size reveals a decreasing trend first and then slightly increasing. The minimum ligament size value (0.41 μm) is achieved at a potential of 850 mV. Meanwhile, the pore size presents a continuously increasing trend with an increase in dealloying potential. Different dealloying potentials result in different reaction driving forces, which have a significant impact on the morphology of the sample. Based on the above experimental results, 850 mV is the optimal dealloying potential in this article for obtaining a good porous prism structure.

In order to further characterize the pore characteristics of the sample after dealloying at 850V for 3 h, nitrogen adsorption–desorption experiments were carried out (Figure S2). From Figure S2a, it can be seen that the adsorption–desorption curve of the sample exhibits a type IV isotherm and contains a type H1 hysteresis loop, indicating its mesoporous characteristics. The specific surface area of the sample is revealed to be 66.3 m²/g. In addition, Figure S2b shows the pore size distribution curve of the sample, and it can be clearly found that the pore size of the sample is mainly in the range of 5–20 nm. The current material exhibits mesoporous properties with a promising pore size distribution, which is conducive to the reaction of the reactants and the release of the products and can maximize the utilization of the exposed active sites, which is conducive to the realization of higher catalytic activity.

Table 3. Chemical compositions (at.%) for different regions by SEM-EDS measurements.

EDS Region	Cu	Pd	Ni	Ag
1 (Figure 2a)	35.60	34.83	24.43	5.14
2 (Figure 5b)	29.96	33.00	30.36	6.67

shows the comparison of EDS detection results between the original ribbon and the dealloyed sample. The results indicate that the composition of the dealloyed product is approximately Cu₃₀Pd₃₃Ni_30.4Ag_6.6, which is different from the composition of the precursor ribbon but remains similar. The composition of the dealloyed sample meets the requirements for HEAs (the atomic proportion of each component is between 5~35%), showing that the dealloyed product is still a quasi-HEA.

In order to systematically study the effect of dealloying time on the porous structure of dealloyed micro prism array, precursor quasi-HEA ribbons were immersed in solution L_a and dealloyed for 0.5 h, 1 h, 2 h, and 3 h at a potential of 850 mV, respectively. Figure 7a–d show SEM images of porous micro prism array for different dealloying times. After dealloying for 0.5 h, obvious pores appear at the grain boundaries of the sample (Figure 7a). When the dealloying time reaches 1 h, the pores at the grain boundaries expand and fully connect, and adjacent prisms separate from each other to form independent prisms (Figure 7b). From Figure 7c,d, it can be seen that as the dealloying time increases, the diameter of the prism continues to decrease, and the pores among the prisms continue to increase. Figure 7e,f show the quantitative relationship between the dealloying time and the diameter of the prism and between the dealloying time and the pore size, respectively. The interrelation between the prism diameter d₁, pore size d₂, and dealloying time t measured from Figure 7a–d is fitted using the following equation [49]:

$$\mathrm{lnd}={\displaystyle \frac{1}{n}}\mathrm{lnt}-{\displaystyle \frac{E}{nRT}}+{\displaystyle \frac{1}{n}}\ln (\mathrm{K}{\mathrm{D}}_0)$$

In the equation, n is the coarsening index, T is the temperature, E is the activation energy for ligament growth, R is the gas constant, and D₀ and K are constants. The fitting results show that ln d (d₁ and d₂) and lnt obtain good linear relationship, and the fitting coefficients R₂ of the two curves reach 0.93 and 0.96, respectively. This indicates that the fitted prism diameter and pore size follow a linear pattern with the dealloying time.

Figure 8 is the illustration reflecting the evolution mechanism of the micro prism array structure during the dealloying process. Figure 8a shows the illustration of the precursor quasi-HEA ribbon. At the beginning of the dealloying, due to the fast corrosion rate and high corrosion power of atoms at the grain boundaries, a large number of lattice vacancies are formed at the grain boundaries [50]. With the increase in dealloying time, the grain boundaries rapidly evolve into cracks, and then the cracks gradually aggregate, causing columnar grains to separate from each other, as shown in Figure 8b. In order to better describe the evolution process of columnar grains after separation as dealloying time increases, Figure 8c shows the enlarged view of a single grain. When the dealloying process reaches a certain level, atoms near the surface are removed on a large scale, and then a large number of lattice vacancies and dislocations are rapidly generated inside the adjacent grains, as shown in Figure 8d. At the same time, newly exposed atoms are transported from the surface to the internal vacancies, causing volume shrinkage of columnar grains, as shown in Figure 8e. During this repetitive process, the driving force that causes the continuous volume shrinkage of columnar grains is caused by the plastic deformation caused by a large number of lattice dislocations inside the grains [51]. Finally, the micro prism arrays shown in Figure 8f are formed.

3.3. Nanoporous PdO/Ag₂O/NiO Composites

In order to prove the feasibility of chemical dealloying of the precursor quasi-HEA ribbon in solution L_b, a three-electrode system was used to test the open-circuit potential of the Cu, Pd, Ni, and Ag pure metals in solution L_b. Figure 9 shows the open-circuit potentials of the Cu, Ni, Ag, and Pd pure metals in solution L_b at 298K, which are −276 mV, −253 mV, −57 mV, and 493 mV, respectively. The results show that among the four elements, Cu exhibits the highest chemical activity in the solution L_b, and the large potential difference between the elements contributes to the driving force for the dealloying process.

Traditional dealloying methods can etch precursor alloys into porous metals [52,53,54]. In a highly oxidizing environment, when dealloying, the porous metal product formed by the corrosion of the precursor alloy will be oxidized, resulting in the formation of a porous oxide composite. This method is used in this work towards quasi-HEA ribbons in highly oxidizing solution L_b to form a porous oxide composite.

The XRD spectra of the corrosion sample in Figure 10 were obtained by free dealloying the precursor quasi-HEA ribbon in solution L_b for 7 h. From Figure 10a, it can be seen that the dealloyed sample is mainly composed of PdO (JCPDS card No. 46-1211), NiO (JCPDS card No. 65-2901), and Ag₂O (JCPDS card No. 41-1104) phases. Based on this result, it can be preliminarily inferred that the corrosion product is a PdO/Ag₂O/NiO composite. It is worth noting that no oxide peaks related to the Cu element are detected in the sample. This indicates that the Cu element is completely removed due to its high electrochemical activity, while the remaining elements are oxidized by highly oxidizing solutions [55]. From the intensity of the diffraction peak, it can be preliminarily determined that the relative amount of PdO is much greater than the amount of NiO and Ag₂O (Figure 10b,c) because the Pd element exhibits the highest electrochemical potential and stability in the solution L_b.

In order to further analyze the valence states of surface elements on the dealloyed product, XPS testing was conducted (Figure 11). The investigation spectrum in Figure 11a shows the core photoionization signals of Cu, Pd, Ni, Ag, C, and O, as well as the Auger signals of Cu and Ag. In addition, a small amount of Cl 2p peaks can be observed, which can be attributed to the residual corrosive agent on the surface of the sample [56]. No Cu 2p XPS peak is observed in Figure 11b. This result further confirms that Cu has been completely removed after dealloying. Figure 11c shows the Pd 3d XPS spectrum. The gap between Pd 3d_5/2 and Pd 3d_3/2 is 5.3 eV, which is consistent with the spin-orbit dual-state separation values reported in the literature [57]. The Pd 3d_5/2 spectrum exhibits a main peak at 337.9 eV and a secondary peak at 340.1 eV, corresponding to PdO and PdO₂, respectively. The Ni 2p spectrum is fitted as two characteristic peaks and a satellite peak (Figure 11d). For the Ni 2p_3/2 spectrum, the two main peaks located at 855 eV and 856.6 eV are attributed to Ni²⁺ and Ni³⁺, respectively [58]. According to previous reports, the formation of NiO can be confirmed [59]. Ag 3d two-photon can be observed clearly in Figure 11e. The centers of Ag 3d_5/2 peak and Ag 3d_3/2 peak are 367.7 eV and 373.7 eV, respectively. The spin-orbit splitting energy between the two is 6 eV. This is consistent with the univalent oxidation state of Ag. This result indicates that the Ag element exists in the form of Ag⁺(Ag₂O) in the corrosion product [60]. Figure 11f shows the high-resolution spectrum of O 1s. The main peak can be decomposed into three peaks: 530.3 eV, 531.8 eV, and 533.3 eV. These peaks correspond to O 1s OM (metal oxide), OH, and OH₂, respectively. Therefore, XPS testing of Cu, Pd, Ni, Ag, and O elements reveals that the main dealloying product is PdO/Ag₂O/NiO composite material, which is consistent with the analysis result of XRD.

Figure 12a–c show the TEM images of quasi-HEA ribbons after free dealloying for 3, 5, and 7 h in solution L_b. From Figure 12a, it can be seen that after 3 h of dealloying, the sample presents a nanoparticle morphology with a diameter size of 20~45 nm. When the dealloying time reaches 5 h (Figure 12b), adjacent nanoparticles tended to connect and assemble into a nano “ligaments/pores” structure. From the enlarged view of the connected area of the nanoparticles (Figure 12d), it can be seen that there is a clear interface between adjacent nanoparticles. It is worth noting that the presence of these interfaces in polycrystalline materials can often be used to create high-energy surfaces, thereby improving catalytic efficiency [61]. In the bright field TEM image (Figure 12c), the dealloyed sample exhibits a transparent nano “ligament/pore” structure. This is similar to the formation process of porous Fe₃O₄ thin films reported in the previous literature. As the nanoparticles are connected together during the dealloying process, the previously formed porous structure becomes clearer, forming a complete layer of porous Fe₃O₄ thin film [62]. The transparent porous structure can provide a more active site for electroactive substances, which is conducive to the entry and exit of electrolyte solutions and is expected to achieve higher catalytic and energy storage properties of quasi-HEAs. The selected area electron diffraction shown in Figure 12e shows that the as-corrosion sample is polycrystalline. The calibrated crystal planes in the diffraction ring correspond to the diffraction peaks in XRD, shown in Figure 10. Based on the above analysis, the synthesis of the nanoporous PdO/Ag₂O/NiO composite can be determined. Figure 12f,g shows the HRTEM image of the dealloyed sample. The lattice stripes marked in the HRTEM image correspond to the lattice planes of PdO (200), Ag₂O (111), and NiO (111), with lattice spacings of 0.282 nm, 0.273 nm, and 0.242 nm, respectively. The above results indicate the formation of nanoporous PdO/Ag₂O/NiO composites. Figure 12h shows the schematic illustration of the evolution mechanism of the nanoparticles. After Cu atoms are preferentially removed, the remaining Pd, Ni, and Ag atoms undergo self-oxidation to form PdO/Ag₂O/NiO nanoparticles. With the increase in dealloying time, nanoparticles are interconnected into ligaments to reduce surface energy. In this situation, it is in the transitional stage of structural transformation between particles and ligaments. When the particles do not have enough time to assemble, a clear interface between the nanoparticles will be created. With the further increase in dealloying time, the nanoparticles completely transform into a nano “ligament/pore” structure and undergo coarsening.

In this work, electrochemical dealloying and free dealloying with two kinds of corrosion liquids were designed. The Cu₃₅Pd₃₅Ni₂₅Ag₅ quasi-HEA synthesized in this paper has good corrosion resistance in many common electrolytes. Usually, in conventional corrosive fluids (NaCl, NaOH, HCl, HNO₃, etc.), there are only some corrosion pits on the surface of the material, while continuous dealloying cannot occur to create a designed porous structure. Therefore, some special etching conditions are tried in this paper. In addition, the aim of this work is to obtain multi-component porous noble metals/metal oxides by dealloying quasi-HEAs. The above factors should be considered in the selection of proper electrolytes.

For the fabrication of multi-component porous noble metals, the ordinary corrosion solution cannot provide enough power to actuate the dealloying in the current quasi-HEA system, so an external potential is applied to carry out constant potential dealloying. In order to achieve a better dealloying effect, a variety of etching fluids have been tried, and the mixture of HNO₃ and H₂SO₄ (L_a) used in this paper shows the best effect. In this process, all the constituent elements undergo some degree of etching due to the relatively high corrosion potential, forming a porous micron prism array structure, which is the first report of such a structure after dealloying. For the synthesis of porous noble metal oxides, strong acids (L_b, HCl, and HNO₃ with 3:1 in mole ratio, aqua regia) are selected to occur not only dealloying but also spontaneous oxidation. In this case, only the copper element with the lowest potential of four elements is corroded away, while other elements are maintained and further oxidized. Concentrated sulfuric acid cannot be used in this process to avoid the formation of metal sulfides [63].

As for why two kinds of different dealloying products can be obtained in this paper, it can be understood from the following aspects. In the field of dealloying, porous metals [64,65,66] and metallic oxides [67,68,69] are two common and well-known dealloying products. Generally, in an alloy system, the relatively active elements are preferentially etched away in dealloying, leaving relatively inert elements to form a porous noble metal structure. When the corrosive fluid contains a certain amount of dissolved oxygen, it first etches the relatively active elements in the dealloying process, and the relatively inert elements are self-assembled and further oxidized to metal oxides. Even when the same material is subjected to different corrosion conditions (acid or alkaline etchant, corrodent concentration, duration), various dealloyed products with different compositions and structures will be obtained [70,71].

In this study, electrochemical dealloying and free dealloying were carried out to create porous noble metals and metal oxides, respectively. To obtain porous noble metals, electrochemical dealloying is performed with potentials higher than 490 mV to ensure all constituent elements can be etched away in this process. This is because the object of etching is a quasi-HEA, which possesses good corrosion resistance. If a low corrosion potential is applied, only ordinary pitting will be generated, or Cu/Ni elements with low potential will be etched, which will not be able to obtain multi-component porous precious metal materials. To fabricate porous noble metallic oxides, a strong acid (aqua regia) is selected for free dealloying. This ensures that after the copper element with the lowest potential is etched out, the remaining elements can self-assemble and self-oxidize to the corresponding noble metallic oxides. However, in general, noble metal oxides are difficult to obtain under mild dealloying conditions. As a result, porous PdO/Ag₂O/NiO oxides are obtained by dealloying in aqua regia solution, revealing that a dealloying strategy can synthesize a wealth of product types.

4. Conclusions

This article provides a new strategy for the porosification of quasi-HEAs and a preparation strategy for porous noble metal oxide composites using constant potential dealloying and free dealloying, respectively. The precursor Cu₃₅Pd₃₅Ni₂₅Ag₅ quasi-HEA ribbon is dealloyed at a constant potential of 850 mV in solution L_a. The porous Cu₃₀Pd₃₃Ni_30.4Ag_6.6 HEA with a micro prism array structure is successfully prepared. The effects of dealloying potential and time on the morphology of dealloyed micro prism array are studied. The evolution mechanism of micro prism arrays is also discussed. When the Cu₃₅Pd₃₅Ni₂₅Ag₅ HEA ribbons are freely dealloyed in solution L_b, the Cu element is selectedly removed, while Pd, Ni, and Ag atoms self-oxidate to form PdO/Ag₂O/NiO nanoparticles. In order to reduce surface energy, the nanoparticles are connected to form a ligament and finally form a “ligament/pore” structure. The work provides strategies for developing novel dealloying products of quasi-HEAs and may promote the development of dealloyed quasi-HEAs in catalysis, energy storage, and other applications.