Kirkendall Effect in Creating Three-Dimensional Metal Catalysts for

May 23, 2017 - In this work, we report an innovative mechanism, the Kirkendall effect, in creating three-dimensional (3D) microporous catalysts with t...
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Kirkendall Effect in Creating Three-Dimensional Metal Catalysts for Hierarchically Porous Ultrathin Graphite with Unique Properties Jianhe Guo,† Andrew Chan,‡ Weigu Li,‡ and D. L. Fan*,†,‡ †

Materials Science and Engineering Program and ‡Department of Mechanical Engineering, The University of Texas at Austin, Austin, Texas 78712, United States S Supporting Information *

ABSTRACT: In this work, we report an innovative mechanism, the Kirkendall effect, in creating three-dimensional (3D) microporous catalysts with tunable pore sizes for the growth of hierarchic ultrathin graphite foams (HP-UGFs) with unique properties. Employing the Kirkendall effect is one of the first demonstrated for fabricating 3D porous catalysts, where tunable pores of 1.9−8.3 μm are created on 3D interconnected struts (∼100 μm). With the catalysts, we readily synthesized freestanding HP-UGFs that offer higher crystallinity and electric conductivity, larger surface area, as well as enhanced electric invariance to strains compared to those of conventional ultrathin graphite foams. A gauge factor as low as ∼10 at a strain as high as 80% is achieved owing to the unique porous corrugations created on the microstruts of the HP-UGFs. This work may inspire a new paradigm in designing and synthesizing a new type of 3D porous architecture made of 2D materials with controlled local corrugations, which could greatly benefit flexible electronics.

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INTRODUCTION Graphene, a crystalline allotrope of carbon with a twodimensional (2-D) hexagonal honeycomb lattice structure, has drawn intensive attention since its discovery owing to its unique properties, including high electron mobility, large thermal conductivity, optical transparency, and excellent mechanical strength.1−6 Although various promising applications have been demonstrated with 2D graphene-based materials,5−9 their performances have yet to be fully exploited. The key challenge is that the strong π−π stacking interactions and van der Waals forces result in self-aggregation and stacking, which further causes a compromised specific surface area. Recently, various innovative approaches for fabricating macroscopic three-dimensional (3D) graphene or ultrathin graphite superstructures have received intensive interest.10−15 When 2D carbonaceous materials are grown into 3D interconnected networks, the specific surface areas and mechanical properties can be substantially improved in addition to the desirable high electrical conductivities and thermal stability. These 3D interconnected structures facilitate many applications, including energy storage and conversion devices, flexible electronics, and heat dissipation in microelectronics.10−22 Substantial efforts have been devoted to the fabrication of 3D graphene or ultrathin graphite superstructures. Two major approaches emerge: (1) fabrication based on graphene oxide (GO) reduction and assembly and (2) template-assisted chemical vapor deposition (CVD).10−18,23−25 Although facile and manufacturable, the 3D graphene superstructures obtained from reduced GO sheets exhibit inferior electrical conductivity because of high-density defects and large junction resistance among different sheets.16,23,26,27 In contrast, 3D graphene and © 2017 American Chemical Society

ultrathin graphite superstructures obtained by template-assisted CVD growth offer excellent electrical and mechanical durability due to high crystalline quality and 3D interconnected architectures.10−12,17,18 However, the feature size of the CVDgrown graphene/graphite is often restricted to approximately 100 μm. This restriction arises from the intrinsic features of the commonly used 3D Ni foam catalysts. Although various other types of metallic catalysts are explored to overcome this challenge,24,25,28 the fabrication efficiency, rationality, or scalability are compromised. The discovery of the Kirkendall effect was first reported as early as 1947.29 However, the formation of pores in alloys due to the Kirkendall effect has long been considered negatively due to the impairment of mechanical properties of alloys. Therefore, for a long time, research on the Kirkendall effect has been focused on reducing the Kirkendall porosity. Recently, the synthesis of hollow nanospheres and nanotubes were achieved with the assistance of this effect, which placed the Kirkendall effect in a new light.30,31 In this work, for the first time, we investigated the Kirkendall effect in the creation of 3D multilevel porous metal catalysts with tunable micropores ranging from 1.9 to 8.3 μm and studied the effect systematically. With the obtained catalysts, we synthesized porous 3D ultrathin graphite foams with hierarchical porosity (HP-UGFs). Compared to conventional graphite foams (UGFs), they exhibit unique and desired properties, including lower density, larger specific surface areas, higher electric conductivity, and Received: April 12, 2017 Revised: May 23, 2017 Published: May 23, 2017 4991

DOI: 10.1021/acs.chemmater.7b01518 Chem. Mater. 2017, 29, 4991−4998

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Chemistry of Materials

cm−2 min−1 and uniform coating of Cu on the 3D Ni foams. The thickness of the Cu layer can be readily tuned by the amount of charges passing through the circuit ranging from 4 to 11 μm (Figure 1(b) and Table S1), which is essential for tuning the composition of the Ni−Cu alloy for creating desired pore sizes to be discussed later. Next, the Ni−Cu composite foams with Cu of various thicknesses are annealed at a temperature of 1000−1100 °C in a gas flow of hydrogen (H2, 5 sccm) and nitrogen (N2, 45 sccm) of 500 mTorr. The annealing process can readily control the degree of alloying of Ni−Cu. This is shown by the monotonical variation of Cu composition on the surface of struts at different annealing times as determined by energy-dispersive X-ray (EDX) spectroscopy in Figure S1. After electrochemical etching at 0.6 V (vs Ag/AgCl, 3 M NaCl; electrolyte: 2 M copper sulfate and 1 M boracic acid), we found dense arrays of micropores distributing uniformly on the interconnected microstruts of the foams across both large and small scales as shown in Figures 1(d) and 3 and Figure S2. In previous studies, porous Ni−Cu alloys can be fabricated by selective etching of the Cu-rich phase from electrodeposited Ni−Cu alloys that are phase-segregated into Cu-rich islands in a Ni-rich matrix.32 To clarify the mechanisms of pore formation in our 3D porous Ni−Cu foams, we carried out a series of characterizations. It is found that, after annealing, Cu and Ni elements are uniformly distributed among each other as shown by the EDX mapping of the struts of the foams in Figure 2(a− c). X-ray diffraction (XRD), showing one peak between those of pure Cu and Ni further suggests that the Ni and Cu are uniformly alloyed after treatment without distinguishing the two phases [Figure 2(d)]. Cross-sectional characterization

much greater invariance of electric resistance to mechanical strains with a gauge factor as low as ∼10 at a strain as high as 80%. The new 3D HP-UGFs could find an array of applications in flexible electronics. The demonstrated design and synthetic mechanisms could inspire a general approach in fabricating 2D materials with 3D multilevel porous superstructures with controlled local corrugations.

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RESULTS AND DISCUSSION Creation of 3D Hierarchically Porous Metal Catalyst by the Kirkendall Effect. Three steps were taken in the synthesis of the 3D multilevel porous Ni−Cu alloy catalysts. In brief, as shown in Figure 1(a−d), on the surface of

Figure 1. Synthesis of HP-UGFs: (a, b) Cu is uniformly plated on the commercial Ni foam; (c) Kirkendall pores are created during the annealing of Ni−Cu; (d) the micropores are exposed on the surface after removing the top layer of struts of Ni−Cu foam by electrochemical etching; (e) HP-UGFs grow on the Ni−Cu alloy foam catalysts; (f) HP-UGFs are freestanding with inherited multilevel porosity from the Ni−Cu catalysts after selective etching of the Ni−Cu catalytic foams.

interconnected struts of Ni foams (MTI Corporation), we deposited a controlled amount of Cu thin films at −1.8 V (vs Ag/AgCl, 3 M NaCl) in a three-electrode electrochemical cell setup. Here, the Ni foam is pretreated with 1 M dilute sulfuric acid (H2SO4) solution for 15 min to remove the native nickel oxide layer before the electrodeposition. The electrolyte is made of copper sulfate (CuSO4, 2M) and boracic acid (H3BO3, 1M). The high electric potential of −1.8 V (vs Ag/AgCl, 3 M NaCl) is chosen to ensure a high deposition rate of ∼1.3 mg

Figure 2. (a−c) EDX mappings of Ni−Cu alloy foams (20 min Cu deposition and annealed at 1100 °C for 5 min). (d) X-ray diffraction (XRD) characterization of pure Cu foil, electrodeposited Cu on Ni foam (20 min Cu deposition), annealed Ni−Cu alloy foams (20 min Cu deposition and annealed at 1100 °C for 5 min), and pure Ni foam, respectively. (e−g) Dependence of average pore sizes on the amount of deposited Cu, annealing temperature, and annealing time, respectively (in (e) and (g), all samples are annealed at 1100 °C; in (f), all samples are coated with 4 μm Cu with 20 min deposition). 4992

DOI: 10.1021/acs.chemmater.7b01518 Chem. Mater. 2017, 29, 4991−4998

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Chemistry of Materials reveals high-density micropores embedded under the surface of the Ni−Cu alloy struts in the Cu-enriched regions as shown in Figure 1(c). After etching the surface layer by electrochemistry, high-density embedded micropores can be readily exposed on the surface (Figure 1(d)). The formation of micropores at the interface of Cu−Ni is a hallmark of the so-called Kirkendall effect, which is attributed to the significantly higher diffusion rates of Cu atoms compared to that of Ni at an elevated temperature.33,34 Our further quantitative analysis supports this mechanism well. The diffusion coefficients (D) of Ni and Cu atoms can be determined with the Arrhenius expression given by D = C·e−E/RT, where the diffusion constants (C) are 0.78, 2.7, 0.57, and 1.77 cm2/s for Cu-in-Cu, Ni-in-Cu, Cu-in-Ni, and Ni-in-Ni, respectively. The activation energies (E) are 50.5, 56.5, 61.7, and 69.1 kcal/mol for Cu-in-Cu, Ni-in-Cu, Cu-in-Ni, and Ni-inNi diffusion, respectively.35 R is the ideal gas constant of 8.314 J K−1 mol−1, and T is the temperature in kelvin. For instance, calculations show that, at 1000 °C, diffusion coefficients of Cu and Ni are 1.67 × 10−9 and 5.38 × 10−10 cm2/s in bulk Cu, respectively, and 1.45 × 10−11 and 2.42 × 10−12 cm2/s in bulk Ni, respectively. These calculations indicate that, at an annealing temperature of 1000 °C, Cu atoms diffuse at a speed at least three times that of Ni in both Cu and Ni matrixes. We further calculated the diffusion coefficients of Ni and Cu in Ni−Cu alloys of different compositions and plotted the results in Figure S3. The results show that the diffusion coefficients of Cu atoms are always higher than those of Ni in any given composition of alloys. It is well-known that, in miscible Ni−Cu systems, diffusions of atoms occur via the vacancy-based diffusion mechanism, where atoms move through the bulk alloy by jumping along vacant lattice sites. Mesoscopically, the distinct diffusion rates of Cu and Ni result in a net mass flow from the Cu-enriched region to the Nienriched region in the struts of Ni−Cu alloy foams, which is equivalent to a net flow of vacancies in the opposite direction. With sufficient time, the vacancies in the Cu-enriched area accumulate and eventually coalesce into larger micropores that reduce the total surface energy driven by thermodynamics.33,34 In previous research, substantial efforts have been devoted to reducing porosity resulting from such a Kirkendall effect because it significantly impairs the mechanical integrity of materials as well as their thermal and electrical conductivities. However, in this work, we made the first attempt to strategically leverage the Kirkendall effect to create a new level of porosity in the entire 3D Ni foams. Furthermore, by controlling the thickness of deposited Cu, annealing temperature, and time, we can readily tune the size of Kirkendall micropores from 1.9 to 8.3 μm as shown in Figures 2(e−g) and 3. The size of micropores monotonically increases with the total mass of electrodeposited Cu from 26 to 70 mg/ cm2and is restricted by the thickness of the Cu-enriched layers. These experimental results agree with the Kirkendall effect: a higher amount of Cu in the thin film alloys leads to higher flux of vacancy diffusion and larger pores. Here, we note that the accumulated vacancies finally condense into large micropores, which lower the total specific surface energy. A similar surface driven effect has been observed in Volmer−Weber-thin-film growth, where individual crystalline grains coalesce into large ones for lowering the total surface energy.36 By the same token, according to the Kirkendall mechanism, higher annealing temperature and longer annealing time, which increases atomic diffusion rates and total diffusion flux, respectively, should also

Figure 3. Characterization of the distributions of pore sizes and scanning electron microscopy (SEM) images of Ni−Cu alloy foams with microporosity of (a, d) 1.9 ± 0.4 μm, (b, e) 4.6 ± 0.6 μm, and (c, f) 8.3 ± 1.4 μm, respectively. The pore size characterization is carried out using ImageJ.

lead to the creation of larger micropores.37 Indeed, the expected dependence has been observed experimentally as shown in Figure 2(f, g), where the pore sizes monotonically increase with the annealing temperature (1000−1100 °C) and time (5−30 min). Note that the dependence on annealing time has a limitation. We found that, when the annealing time is rather long, e.g., 2 h (at 1100 °C), the micropores will grow and finally break out on the surface of Ni−Cu alloys as shown in Figure S4. When the temperature is as low as 900 °C, no Kirkendall micropores are observed. These results further support that the formation and tunability of these micropores are governed by the Kirkendall effect. As discussed previously, we exposed the micropores embedded in the struts of the Ni−Cu foams by removing the top layers via electrochemical etching at 0.6 V (vs Ag/AgCl) as shown in Figure 1(d). With increased etching time, micropores gradually emerge at around 50−100 s and reach a maximum density at ∼500 s (Figure 4(a−f)). The morphology of the micropores is also altered from circles (Figure 4(c, d)) to crystalline polyhedrons (Figure 4(e, f)) during the etching progress. The change in the shape could be attributed to anisotropic etching and atomic reconstruction, where different crystalline planes of metals have distinct atomic etching and diffusion rates. Sharp facets and edges are created in minimizing the total surface energy. Next, the specific surface area as a function of pore size was determined by 5-point Brunauer−Emmett−Teller (BET) analysis (performed at Pacific Surface Science Inc., Oxnard, CA). It shows that the surface areas of the multilevel porous Cu−Ni alloy foams increase monotonically with the decrease in size of the micropores as in Figure 4(g, h). At a pore size of 1.9 μm, approximately 3-times increment of the total surface area is obtained compared to that of the original Ni foam. With further reduction of the pore sizes, i.e., by decreasing the deposition thickness of Cu, tuning the annealing temperature, time, and the cooling process, we expect to further expand the total surface areas. The rational approach to forming and tuning micropores on existing 3D porous structures suggests a new route in creating hierarchically porous Ni−Cu foams by cycling the process of electrodeposition, annealing, and etching at controlled conditions for specific pore sizes in each cycle, which is under investigation. The fabrication processes are advantageous because of their low cost and compatibility with conventional fabrication techniques and equipment, such as electroplating, 4993

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areas and two levels of porosity of ∼100 μm and a few of micrometers (tunable by the catalysts), respectively, largely replicating the features of the Ni−Cu alloy foams as shown in Figure 5(a−c) and Figure S6.

Figure 4. (a−f) SEM images of the surface of struts of Ni−Cu alloy foams after electrochemical etching for 50, 100, 250, 500, 750, and 1000 s, respectively (Ni−Cu alloy foams: 40 min Cu deposition and annealing at 1100 °C for 15 min). (c, inset) Effect of anisotropic chemical etching on the created micropores (scale bar: 2 μm). (g, h) Brunauer−Emmett−Teller (BET) characterization: nitrogen sorption isotherms (g) and specific surface areas (h) of Ni−Cu alloy foams with average pore sizes ranging from 1.9 to 8.3 μm compared to the original Ni foam.

Figure 5. (a−c) SEM images of HP-UGFs with tunable microporosities, templated from Ni−Cu catalysts with pore sizes of 1.9 (a), 4.6 (b), and 8.3 μm (c), respectively. (c, inset) Zoom-in SEM image (scale bar: 20 μm). (d) HRTEM image of HP-UGFs (1.9 μm catalytic template, density: 40 mg/cm3). (e) Raman spectra of HP-UGFs with the same areal density and different pore sizes. (inset) Ratio of intensity of D to G peaks versus pore size. (f) As-measured conductance of HP-UGFs (same dimensions: 1.5 cm × 0.5 cm × 200 μm) versus volumetric density and (g) the conductivity of HPUGFs normalized by total volume fraction according to the Lemlich model.

electro-etching, and annealing, for mass production. Given the unique morphologies and growth mechanisms, the Ni−Cu alloy foams could find applications ranging from electrodes in batteries to catalyst support in fuel cells owing to the 3D porous structures, enhanced surface area, and high electric conductivity.38−40 They could also be applied as catalysts with high activities for the growth of 2D materials into 3D multilevel porous structures,41,42 which is vital for next generation flexible optoelectronics.43,44 Synthesis, Characterization, and Electronic Properties of HP-UGFs with Multilevel Porous Ni−Cu Foams as Catalysts. Here, we employed porous Ni−Cu alloy foams to catalyze the growth of 3D freestanding HP-UGFs, which have potential applications in flexible electronics. The HP-UGFs inherit the tunable and hierarchical porosities of the Ni−Cu foams with enhanced surface areas. They also exhibit desirable electronic properties due to the catalytic properties of the Ni− Cu alloy foams. The fabrication of HP-UGFs is achieved by chemical vapor deposition (CVD) in a gas mixture of ethylene (C2H4) and hydrogen (H2) at 700 °C from minutes to hours followed by selective etching of the Cu−Ni foams in a solution of iron chloride (FeCl3, 1 M) and hydrochloride (HCl, 2 M) at room temperature. In particular, using ethylene as the carbon source reduces the required temperature compared to that of methane (CH4),45 which is essential for preserving the hierarchically porous morphology of the Ni−Cu foam catalysts during the high-temperature growth. The XRD characterization shown in Figure S5 indicates that the Ni−Cu alloy foams have been removed well. The details are provided in the Supporting Information. With this process, we successfully obtained freestanding ultrathin graphite foams with enhanced surface

The HP-UGFs become freestanding after the Ni−Cu alloy foams are etched. As shown in Figure S7, the HP-UGFs largely retain their original morphology of the catalytic alloy foams and can be facilely handled by tweezers. Experimentally, we found that the HP-UGFs with smaller pore sizes were more resistive to fractures during sample processing compared to those with larger micropores or those without micropores (nonporous UGFs) with the same areal mass density. Note that, at the same areal mass density, the walls of the graphite struts in HP-UGFs are even thinner than those in UGFs (without the second-level micropores) because the surface area of HP-UGFs can be a few times larger than that of the UGFs. This can be qualitatively understood from the availability of the corrugations created on the entire 3D struts of the graphite foams, which provide substantial allowance for mechanical deformation when receiving strains found in corrugated 2D materials.46,47 In previous studies, to reinforce 3D graphene made from simple Ni foams, the so-called UGFs, polymer residues such as polydimethylsiloxane (PDMS), are coated on the surface of 3D graphene,10 which adversely increase the total electric resistance and alter various electronic properties of 3D graphene. Moreover, in comparison to UGFs, HP-UGFs grown from the porous Ni−Cu foams show higher uniformity in thickness (Figure S8). Such a property can be attributed to the distinct 4994

DOI: 10.1021/acs.chemmater.7b01518 Chem. Mater. 2017, 29, 4991−4998

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(Figure 5(g)). There are two factors that could explain the observation: one is the higher uniformity of graphite obtained from Ni−Cu alloy foams as shown in Figure S8. This also applies to HP-UGFs made from Ni−Cu alloy foam catalysts (1.9−8.3 μm), where the composition of Cu on the foam surface decreases with the increase in pore size for the tested samples as shown in Figure S10. The higher fraction of Cu in the alloy leads to higher crystallinity. Therefore, the crystallinity of smaller-pore HP-UGFs can be greater than that of largerpore HP-UGFs. The analysis is supported by Raman spectroscopy, where the average ratio of D peak to G peak (ID/IG) monotonically increases with pore size in Figure 5(e). The other factor could be attributed to the larger surface areas of smaller-pore HP-UGFs that result in lower effective thickness of graphite struts at a given mass density. As discussed previously, thinner graphite has a higher normalized electric conductivity. Therefore, both factors lead toward the observation that porous HP-UGFs with reduced pore sizes have higher normalized electric conductivity (Figure 5(g)). Finally, leveraging the unique structural, mechanical, and electric properties, we demonstrated the HP-UGFs as flexible electric conductors for wearable electronics. Wearable electronic devices that can be bended and folded show remarkable advantages over rigid counterparts in revolutionizing modern lives due to their lower weight, higher flexibility, better durability, and improved comfort.54,55 Innovative advances have been made in the development of various wearable electronic and energy storage devices, such as in flexible strain sensors, biochemical sensors, lithium−oxygen batteries,56−58 zinc−air batteries59 and sodium-ion batteries.60 Flexible conductors that offer high electric conductivity, low variation in electric conductivity under strain, and large mechanical compliance are essential components for wearable electronic devices. The HP-UGFs with a second level of porous structures over the struts not only offer enhanced electric conductivity but also provide better mechanical allowance due to the local corrugations when subjected to deformation. To explore this application, we made polymer-graphite strain sensors by integrating the HP-UGF with PDMS elastomer via infiltration and curing as shown in Figure 6(a) (fabrication details are provided in the Supporting Information). For comparison purposes, we tested both HP-UGFs and nonporous

catalytic effects of Ni and Cu: the growth of graphene on the surfaces of Ni is via a carbon segregation process, whereas that on Cu is via a surface adsorption process.48 The different mechanisms result from a much lower saturation solubility of carbon in Cu than that in Ni at high temperatures. Experimentally, it is found that Cu is much more favorable for catalyzing uniform growth of single layer graphene compared to that of Ni,48 whereas with Ni as the catalyst, multilayer graphene with nonuniform thickness is often obtained due to nonequilibrium precipitation of carbon.41 Catalysts also affect the growth rate of ultrathin graphite. The lower saturation solubility of carbon in Cu gives rise to the observation that the thickness of HP-UGFs reduces with an increase in total Cu composition as well as pore size given the same growth time. Consequently, the Cu−Ni alloying of the porous foams can effectively tune the solubility of carbon for the growth of highly crystalline ultrathin graphite49 as shown in the characterizations with HRTEM (Figure 5(d)), Raman spectroscopy (Figure 5(e) and Figure S9), as well as enhanced uniformity (Figure S8). The unique HP-UGFs with hierarchical porous superstructures is a new material with many unknown physical properties. Here, we characterized the electric conductivity of HP-UGFs as a function of pore size and mass density. According to the Lemlich model, the normalized electric conductivity of porous materials is given by50 σnormalized =

3σas ‐ measured φ

where σas‑measured is the as-measured electric conductivity of the HP-UGFs (or UGFs) (Figure 5(f)) and φ is the volume fraction. The Lemlich model has been widely applied in calculating electronic conductivities of foam structures with low volume fractions.18,50,51 With this model, both the normalized conductivities of HP-UGFs and nonporous UGFs are determined as a function of mass density as shown in Figure 5(g). It is found that the normalized electric conductivities of HP-UGFs and UGFs both decrease with the density of graphite foams. Note that, for a given porous structure, the material density (ρ) is proportional to the thickness of struts (t) with longer growth time. Therefore, we can readily determine that the normalized electric conductivities of HP-UGFs and UGFs monotonically increase with reduction of thickness of struts. This finding is supported by other experiments and simulations.52,53 It can be understood from two distinct modes of electron transport in graphene or graphite layers. For single layer graphene, electrons only transport via surface channels, where the electrons move directionally at a high mobility approaching Fermi velocity.53 In contrast, two types of directional transmission channels of electrons are present for ultrathin graphite: the aforementioned in-layer surface channel and interlayer channel among different graphene sheets. When electrons are transported via interlayer channels, their motions are affected by both phonon and electron scattering and thus have a lower mobility than those via surface channels.53 As a result, for ultrathin graphite made of multilayers of graphene, the overall contribution of the high-speed surface channel decreases with the thickness of graphite until reaching the value of bulk graphite, which agrees with our observation in experiments in Figure 5(g). We also note that the HP-UGFs with smaller pores exhibit higher electric conductivities compared to those with larger pores or those without micropores at a given mass density

Figure 6. (a) Schematic illustration of the fabrication process of UGF/ PDMS composite strain sensors. Normalized change of electrical resistance (ΔR/R0) versus (b) bending radius and (c) tensile strain of HP-UGF and UGF. The UGF foams are fabricated from simple Ni foams. HP-UGF and UGF have the same density of ∼125 mg/cm3, thickness of 200 μm, and lateral dimension of 1 × 2 cm2. 4995

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UGFs integrated in PDMS, which have the same density of ∼125 mg/cm3, thickness of 200 μm, and lateral dimension of 1 × 2 cm2. When subjected to bending and tensile tests, the change of electrical resistance of HP-UGF and UGF are recorded as shown in Figure 6(b and c). The electrical resistances of both HP-UGF and UGF exhibit a similarly increasing trend with mechanical deformation, which can be attributed to the intrinsic piezoresistivity due to elastic deformation as well as extrinsic piezoresistivity due to crack generation.61 However, HP-UGFs show much lower changes of electric resistivity compared to those of UGFs, particularly when subjected to high mechanical deformations. For instance, for a bending radius of 1.7 mm, the resistivity of HP-UGFs only changed by 34.6%, compared to 66.6% obtained from UGFs. In tensile tests with strain as high as 80%, the Gauge factors of HP-UGFs and UGFs are 10 and 100, respectively. Here, the gauge factor tells the magnitude of variation of signals with strain defined by the ratio of relative resistance change (ΔR/R0) and the mechanical strain (ε). Therefore, the invariance of electric signals to mechanical deformations of HP-UGFs is only ∼50 and 10% of those of UGFs at a bending radius of 1.7 mm and in-plane strain of 80%, respectively. Furthermore, it can be found that the electric signals of UGFs increase exponentially when the bending curvature is above 40%, where cracks could be generated.61,62 On the contrary, for HP-UGFs, the electric signals do not have such an exponential increase (up to 80% strain). Moreover, the HP-UGFs offer high reproducibility in signal detection among different testing cycles, suggesting its high fidelity for applications (Figure 6(c)). Overall, these characterizations indicate that the unique second-level porosity created on the microstruts of HP-UGFs is highly desirable for lowering piezoelectric resistance. The concept of creating multilevel porous structures point toward an important route in designing and fabricating flexible electric conductors.

Article

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b01518. Electrodeposition conditions and analysis of Cu thin films; additional SEM and optical images, composition and pore size characterizations of Ni−Cu alloy foams; additional SEM and optical images, XRD, XPS, Raman, and density characterizations of UGFs and HP-UGFs; and fabrication details of HP-UGFs and HP-UGFs/ PDMS composites (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

D. L. Fan: 0000-0002-4724-2483 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We are grateful for the support of the National Science Foundation (Grant CMMI 1563382) and Welch Foundation (Grant F-1734).

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REFERENCES

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CONCLUSIONS In summary, we investigated the fundamental mechanism of Kirkendall effects in the synthesis of multilevel porous Ni−Cu foams. The Kirkendall pores are often considered as defects generated during the annealing processes that impair the mechanical integrity of materials, whereas here it is employed strategically to create hierarchically porous microsuperstructures of catalysts, which may inspire a new approach in synthesizing an array of 3D porous micro/nanostructures. By tuning reaction conditions, we can readily tune the Kirkendall pores from 1.9 to 8.3 μm created on the interconnected 100 μm struts of the 3D foams. The as-obtained multilevel porous Ni−Cu foams can readily catalyze the growth of a unique type of HP-UGFs, which inherit the superstructures of the catalysts and offer much larger specific surface areas, higher electric conductivity, and significantly enhanced invariance in electric resistance under mechanical strains compared to those of commonly made UGFs. These unique properties are highly desirable for an array of applications in electronics and energy storage and conversion devices. In particular, the low variation of electric conductivity under strain due to the local corrugations created on the struts of foams, makes them highly potent for flexible energy and electronic devices. This work could also inspire a new paradigm in creating 2D materials with 3D multilevel porous architectures with controlled local corrugations. 4996

DOI: 10.1021/acs.chemmater.7b01518 Chem. Mater. 2017, 29, 4991−4998

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DOI: 10.1021/acs.chemmater.7b01518 Chem. Mater. 2017, 29, 4991−4998