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Synthesis of Nanostructured/Macroscopic Low-Density Copper Foams Based on Metal-Coated Polymer Core−Shell Particles Sung Ho Kim,*,† Nick Bazin,‡ Jessica I. Shaw,† Jae-Hyuck Yoo,† Marcus A. Worsley,† Joe H. Satcher Jr.,† John D. Sain,† Joshua D. Kuntz,† Sergei O. Kucheyev,† Theodore F. Baumann,† and Alex V. Hamza† †

Nanoscale Synthesis and Characterization Laboratory, Lawrence Livermore National Laboratory, 7000 East Avenue, Livermore, California 94550, United States ‡ Atomic Weapons Establishment, Aldermaston, Reading RG7 4PR, United Kingdom S Supporting Information *

ABSTRACT: A robust, millimeter-sized low-density Cu foam with ∼90% (v/v) porosity, ∼30 nm thick walls, and ∼1 μm diameter spherical pores is prepared by the slip-casting of metal-coated polymer core−shell particles followed by a thermal removal of the polymer. In this paper, we report our key findings that enable the development of the low-density Cu foams. First, we need to synthesize polystyrene (PS) particles coated with a very thin Cu layer (in the range of tens of nanometers). A simple reduction in the amount of Cu deposited onto the PS was not sufficient to form such a lowdensity Cu foams due to issues related to foam collapse and densification upon the subsequent polymer removal step. Precise control over the morphology of the Cu coating on the particles is essential for the synthesis of a lower density of foams. Second, improving the dispersion of PS−Cu particles in a suspension used for the casting as well as careful optimization of a baking condition minimize the formation of irregular large voids, leading to Cu foams with a more uniform packing and a better connectivity of neighboring Cu hollow shells. Finally, we analyzed mechanical properties of the Cu foams with a depth-sensing indentation test. The uniform Cu foams show a significant improvement in mechanical properties (∼1.5× modulus and ∼3× hardness) compared to those of uncontrolled foam samples with a similar foam density but irregular large voids. Higher surface areas and a good electric conductivity of the Cu foams present a great potential to future applications. KEYWORDS: copper, metal foam, porous materials, core−shell particles, electroless deposition

1. INTRODUCTION Nanostructured porous metals are a fascinating class of materials because they could combine useful properties of metals with extreme features of porous materials such as low density, high surface area, and a high strength-to-weight ratio.1−5 This unique combination provides an ideal environment for new applications in areas of catalysts, sensors, hydrogen storage, batteries and supercapacitors, separation membranes, biocompatible scaffolds, and lightweight composite materials.1−3,6−8 However, it still remains a challenge to produce large-area continuous films or macroscopic monoliths (∼millimeters and above) with precisely controlled porosity and density, pore geometry, and uniformity.9−17 The development of a simple, versatile synthetic route is highly required for metal foams to find widespread technological applications. Here, we present a metal-coated polymer core−shell particle approach to nanostructured/macroscopic copper foams that allows control over the pore size and foam density, leading to a good uniformity and superior mechanical properties of resultant foams. © XXXX American Chemical Society

Sacrificial template approaches such as colloidal crystal templating and dealloying are the most widely applied, but they depend on a diffusion process through very narrow channels and hence are mostly limited to thin films with an intermediate density range.9−17 Previously, we developed a synthetic route to porous copper monoliths that start as a liquid suspension of Cu nanoparticles and polystyrene (PS) spheres.11 Despite great flexibility of this bottom-up synthesis,11 porosity and pore size of the resulting copper foams were limited. To overcome these limitations in conventional approaches, we have explored the application of metal-coated polymer core− shell particles to the preparation of metal foams.18−23 In this approach, the coating of colloidal particles (core) with metals (shell) is performed prior to the casting of the core−shell particles, as we previously demonstrated in the synthesis of ultralow density gold foams.20 This does not require the tedious Received: September 28, 2016 Accepted: November 23, 2016

A

DOI: 10.1021/acsami.6b12320 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces Scheme 1. Synthesis of Copper Foams

and further dried at ambient conditions. Finally, to remove the PS template and prepare porous Cu foams, the PS−Cu was baked in a tube furnace at 400 °C for 4 h under 4% H2/Ar atmosphere. Additional details on the synthesis of copper foams are described in the Supporting Information. 2.2. Characterization. Monolithic PS−Cu sample cast in the plaster of Paris mold had a cylindrical shape, and the Cu foam held the shape even after the baking process. Bulk densities of the samples were thus determined from the measurement of the weight and sizes:

step to infiltrate the small voids of prearranged template crystals with the metal. Recently, Kränzlin et al. prepared macroscopic parts of copper-based foams using copper-coated zinc oxide (ZnO−Cu) core−shell particles.21 However, the copper layer in such foams had a thickness of ∼200 nm, and irregular larger voids between loosely packed particles still existed in the copper foams.21 The objective of this study is to expand a metal-coated core−shell particle approach to be more suitable for the preparation of low-density copper foams. First, we present a synthetic route to significantly lower the thickness of the copper layer (down to the tens of nanometers thick range) and maintain Cu shells from collapsing, which is a key that enables us to synthesize very low density copper foams (∼10% relative to full density Cu). Second, we demonstrate that a welldispersed suspension of metal-coated core−shell particles minimizes the formation of irregular large voids, resulting in metal foams with extremely uniform spherical pores. Finally, we investigate the effect of these fine-tuning steps in synthetic conditions on mechanical properties of copper foams, which leads to a better understanding of the formation mechanism.

W

ρ= π×

2

( d2 )

×h

(1)

where ρ is the bulk density, W is the weight, d is the diameter, and h is the height of a sample. Morphology of fractured (and indented) samples was investigated with a Jeol JSM-7401F scanning electron microscope (SEM) at an acceleration voltage of 2−3 kV in a lower secondary electron image (LEI) mode without any additional conductive coating. BET surface areas were measured by nitrogen adsorption porosimetry with a Micromeritics Instrument ASAP 2000 after degassing at 70 °C for 12 h. Raman and photoluminescence (PL) spectra of samples were obtained by a lab-built Raman spectroscopy system. A Raman excitation laser (532 nm, 2.5 mW) and a 20× objective lens (NA: 0.42, Mitutoyo) were employed, and a HgNe calibration lamp (Newport) was used for the wavenumber calibration. The spectrometer was cooled with a liquid nitrogen, and slits with different gratings of 1200 and 150 lines/mm were used for the measurement of Raman shift and PL spectra, respectively. Mechanical properties of copper foams were characterized with a MTS XP Nanoindenter equipped with a 200 μm diameter spherical indenter. Before indentation, copper foam monoliths were attached to silicon wafers with epoxy and diamond-turned to have a flat surface. A series of load−unload indents were carried out in laboratory air at room temperature. The loading rate was continuously adjusted to keep a constant representative strain rate of 10−2 s−1. For every cycle, the unloading rate was kept constant and equal to the maximum loading rate of the cycle.27 The uniformity of copper foams was further characterized with an Xradia Micro XCT-200 X-ray microscope. X-ray projection images were collected (0.25° angular spacing, 745 projections) and reconstructed to provide 3D X-ray computed tomography (CT) data sliced vertically and horizontally.

2. EXPERIMENTAL SECTION 2.1. Synthesis of Copper Foams. Scheme 1 summarizes the synthetic procedure of copper foams used in this study: (I) incorporation of Pd seeds into PS particles (seeding); (II) electroless deposition of Cu onto the Pd-seeded PS particles (Cu coating); (III) assembling copper-coated PS particles into a monolith (casting); and (IV) removal of the PS template to form Cu foams (baking). First, Pdseeded PS particles (PS−Pd) were prepared by treating carboxylateterminated polystyrene (PS−COOH, 1 μm diameter, 2.6% solidslatex, Polysciences Inc.) with (i) polyethylenimine (PEI) solution,18 (ii) K2PdCl4 solution, and (iii) the reduction with NaBH4. Second, coating of copper onto PS−Pd particles was performed by electroless Cu plating using formaldehyde (HCHO) as a reducing agent at a basic condition.11,24−26 In this preparation, a solution of CuSO4·5H2O (1 g), potassium sodium tartrate (2.6 g), and NaOH (660 mg) in water (100 mL) was used as a Cu plating solution. Different densities of copper foams (∼0.7 to ∼1.1 g/cm3) were obtained by changing the ratios of Cu plating solution (ranging from 20 to 60 mL) to PS−Pd (25 mg) (see Figures S1 and S2). It should be mentioned that that the solution-based, multistep nature of this porous copper preparation limited accuracy and precision with an error of about 10−15% in the Cu foam densities although significant optimization steps were made in this study. Typically, for the synthesis of an ∼0.9 (±0.1) g/cm3 copper foam, 30 mL of the Cu plating solution was added to the mixture of PS−Pd (25 mg), NaOH (1M, 150 mL), and water (300 mL) under stirring. The addition of formaldehyde (37%, 28 mL for 1 a.u. of HCHO addition) gradually turned the color of the suspension from light blue to reddish brown and then to completely colorless within a few minutes, and particles started to precipitate. After 30 min of the reaction, reddish brown PS−Cu precipitates were first washed with excess water and then redispersed by additional stirring in water and a short sonication using a tip sonicator (Cole-Parmer Instruments CP502). Third, the concentrated suspension of PS−Cu particles in water (∼2 mL) was poured into a Teflon tube (a diameter of 4.8 mm) in a plaster of Paris mold, and the mold was transferred into an ultrasonic bath (Branson 2510). Once a PS−Cu monolith was cast in the mold, the monolith was carefully separated from the Teflon tube

3. RESULTS AND DISCUSSION Scheme 1 shows a schematic of our metal-coated core−shell particle approach to prepare low-density copper foams. First, we need to synthesize copper-coated core−shell particles using 1 μm diameter polystyrene (PS) spheres. Although the fundamental aspects of the electroless copper plating process were already covered,24 there appeared to be little known about the growth and structure of electroless deposits of very thin films on nonconducting substrates, to the best of our knowledge. A simple reduction in the amount of metals deposited onto a template was not sufficient to form lower density metal foams due to issues related to foam collapse and densification upon the subsequent template removal step. Figure 1a shows SEM images of PS−Cu monoliths prepared B


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ACS Applied Materials & Interfaces

Figure 1. Control over the morphology of Cu deposits on PS particles. SEM images of (a) PS−Cu monoliths and (b) copper foams, prepared with different amounts of formaldehyde (HCHO). (c) Different densities of Cu foams were obtained depending on the amount of HCHO and the morphology of Cu coatings. (d) Increase in the HCHO addition varied morphologies of Cu deposits from discrete large copper particles to small particles to continuous film. Scale bars in panels a and b are 1 μm.

Figure 2. Process to redisperse large precipitates into individual particles. (a, b) Electroless Cu plating reaction led to not only the coating of copper onto PS particles but also the agglomeration of copper-coated particles. Large precipitates were redispersed by washing with excess water and subsequent stirring in water for ∼3 h. (c, d) SEM images of (c) PS−Cu monoliths prepared from a direct suspension of large precipitates or (d) a well-dispersed suspension after a redispersion process. (e) Black monoliths (PS−CuO) with a needle-like morphology were often obtained from samples prepared using insufficient washing or drying at high temperature.

with different amounts of formaldehyde (HCHO) at both low and high magnifications (see Table S1 for the composition and

Figure S3 for additional SEM images). Interestingly, the grain size of Cu deposits varies largely, with the amount of HCHO C


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Figure 3. Optimization of baking conditions to remove the PS template. (a, b) Changes in weights, volumes, and densities were measured for 4 different samples before and after baking at either 400 °C (samples S1 and S3) or 450 °C (samples S2 and S4). (c) Selected SEM image of a fractured Cu foam sample baked at 400 °C. Scale bar is 1 μm.

(with tens of nanometers thickness) on PS spheres is a key that enables us to synthesize a low density of copper foams. The electroless copper plating reaction in water using excess formaldehyde leads to not only the coating of copper onto PS but also the agglomeration of copper-coated particles, as shown in a schematic and photos (Figure 2a,b). We often observed that samples prepared from a suspension of large precipitates were easily broken at handling or mechanical machining. SEM image of the sample in Figure 2c revealed the formation of large voids surrounded with the PS−Cu agglomerates. To remove these irregular large voids and improve the uniformity of the pore structure in metal foams, we wanted to redisperse large PS−Cu agglomerates into separate individual particles and prepared a well-dispersed suspension prior to the casting process. As demonstrated in photos in Figure 2b, they were redispersed in water through a washing of the precipitates with excess water and a subsequent stirring in excess DI water at a neutral pH for ∼3 h. The well-dispersed PS−Cu suspension results in a quite uniform monolith composed of an ideal closepacking of spheres with no large voids (Figure 2d). It is worth mentioning that letting a Cu deposition (plating) reaction sit for several hours without these washing/stirring steps did not yield a complete redispersion. Here, the pH of water was an important parameter. The use of water at acidic pH caused the complete etching of Cu coatings, and basic pH caused the oxidation of them to CuO. Although not completely understood, changes in experimental conditions such as incomplete washing or casting and drying in a high temperature oven often resulted in the formation of a black monolith with a needle-like morphology (Figure 2e). Elemental analyses on the area of needles with the SEM/EDS technique indicate that the morphological transformation was caused by the oxidation of Cu foams to CuO (data not shown). This can be explained by Pourbaix diagrams to graphically show the regions of thermodynamic stability for the metal/electrolyte pair. Typically, copper has several predominant regions for elementary Cu (immunity), Cu2O or CuO (passivity), and dissolved species of Cun+ or Cu(OH)n2−n (corrosion). According to the calculation of Pourbaix diagrams for copper,33 the stability area of elementary Cu was mainly located near at a neutral pH, and the area became smaller as we moved to lower or higher pH. At low pH, copper was dissolved, and copper

addition acting as a reducing agent. Morphological transitions from separated large Cu grains to smaller grains to almost continuous thin film are observed with an increase in HCHO from 0.02 a.u. (1×, 0.56 mL of 37% HCHO addition to 450 mL of NaOH/H2O solution) to 1 a.u. (50×, 28.0 mL), all with the same pretreatment (seeding) and the same amounts of PS−Pd beads and the Cu plating solution. It is worth mentioning that at a minimum concentration (0.02 a.u.) of HCHO the supernatant became totally clear and colorless indicating that no copper ions were left after the Cu plating reaction. Figure 1b shows SEM images of those copper foams baked at 400 °C for 4 h under a 4% H2/Ar gas. Upon the removal of a PS template, the samples of PS−Cu monoliths coated with discrete large Cu grains (prepared with less HCHO) collapse and the copper foams lose their initial porosity. In a sharp contrast, the diameter of a Cu shell in the copper foam prepared with 1 a.u. of HCHO is nearly identical with that in the PS−Cu monolith. It explains how different densities of copper foams are obtained depending on the morphologies of Cu deposits and coverage of the Cu coating in particles (Figure 1c). Although some differences were reported as to details on the reaction order and kinetics for each species during electroless copper plating, they reported that the rate of Cu reduction increases with formaldehyde concentration.28−32 For example, El-Raghy and Abo-Salama studied oxidation and reduction processes involved in electroless copper deposition using a two-chamber galvanic cell and determined an empirical equation of the rate constant: r = k[Cu 2 +]0.37 [OH−]0.254 [HCHO]0.082 [tartrate]0.194 (g cm−2 hr −1)

(2)

where k is the rate constant (i.e., 82.9 at room temperature), and [X] is the concentration of the components including formaldehyde, copper, hydroxide, and tartrate ions.31 This indicates that ∼50× HCHO addition (1 a.u. vs 0.02 a.u.) may result in ∼1.4× increase in the rate of Cu deposition. Our study demonstrates that the concentration of HCHO addition plays a crucial role in controlling the growth and shape of metallic Cu deposits on PS templates, as schematically represented in Figure 1d. Further studies on the effect of different components and temperature will be an interesting work, which is above the scope of this study. Formation of very thin conformal Cu films D


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Figure 4. Characterization of Cu foams by Raman spectroscopy. (a) Peaks in the range of ∼1600, ∼1400, and ∼1100 cm−1 in the Raman spectra indicate that a trace of amorphous carbons left in the Cu foams baked at 400 °C. (b) Strong photoluminescence (PL) at ∼590 nm represents that the native Cu2O with a bandgap of ∼2.1 eV was formed.

copper monoliths (data not shown), the synthetic route to low density copper foams through the formation of copper oxides and the subsequent reduction was not considered in this study. Raman spectroscopy was also used to characterize further structural details of the Cu foams. The Raman shift spectra of Cu foams (baked at 400 °C) in Figure 4a showed several peaks in the shift range of 800−2000 cm−1. Typically, the peaks at ∼1600 cm−1 (G mode) and ∼1450 and ∼1150 cm−1 (transPA) were assigned to be due to the bond stretching of all sp2 carbon atoms and the sum and difference combinations of C C chain stretching and CH wagging modes, respectively.34,35 This indicates that a trace of amorphous carbons were formed and still left in the final Cu foams, although almost complete removal of the PS was observed from weight loss and SEM images above. Here, we also observed that the background of Raman spectra increased with the wavelength of Raman shift. To understand the origin of the increasing slope, we did another Raman/photoluminescence (PL) measurement using a different grating (150 line/mm). Interestingly, our Cu foam sample has a very strong PL, as shown in Figure 4b. Cuprous oxide (Cu2O) has a bandgap of ∼2.1 eV (∼590 nm), and cupric oxide (CuO) has a bandgap of ∼1.2 eV (∼1033 nm). The strong PL at ∼590 nm in Figure 4b represents that native Cu2O was formed during the storage under ambient condition. It is worth remembering that Cu foam samples were prepared under a reducing atmosphere at high temperature. No detectable oxidation was observed by SEM/EDS from fresh samples under a normal storage procedure (similarly to Figure 6f). We think that this might be attributed to sensitivity of Raman and PL spectroscopy. This study proves that Raman spectroscopy is a crucial tool to reveal the details of Cu foams. Otherwise we could not detect them. Further study on the kinetics of Cu oxidation using Raman spectroscopy is outside the scope of this study and will be treated in future publications. Copper foam samples prepared by annealing at 400 °C for 4 h under a flow of 4% H2/Ar gas are strong enough to be handled and machined into a customized shape. Here, a depthsensing indentation experiment, also often called nanoindentation, is carried out to characterize mechanical properties of the copper foam sample (sample A, ∼0.8 g/cm3) prepared by our standard procedure. For a comparison, another sample (sample B, ∼0.8 g/cm3) with a similar density but some irregular large voids and the evidence of pore collapse was purposely prepared by reducing the amount of HCHO addition and the duration for redispersion, while maintaining other

ions were formed. Oxidizing environments at slightly alkaline pH oxidized elementary copper to Cu2O, which at a higher potential was further oxidized to CuO. Increasing temperature also reduced the predominant area of Cu. Further details on the effect of these experimental parameters are outside the scope of current work. Baking PS−Cu samples under a reducing atmosphere at high temperature is used to remove a PS template, promote interparticle bonding and mechanical properties of metal foams, and reduce surface oxides in Cu shells at the same time.11 Typical thermogravimetric analysis (TGA) curves of PS and PS−Cu samples showed an almost complete decomposition of PS in the temperature range from ∼250 to ∼420 °C, although the decomposition patterns were characteristic for each sample (see Figure S4). To optimize baking conditions, four different samples with initial densities of ∼1.1 and 1.2 g/ cm3 were baked at either 400 or 450 °C for 4 h in a flow of 4% H2/Ar gas. Changes in weights, volumes, and densities were measured before and after baking (Figure 3a,b). No additional weight loss is observed with increasing baking temperature from 400 to 450 °C. Samples baked at 450 °C also show a serious increase in densities of the Cu foams, which might be attributed to an aggregation of Cu shells involving interparticle bonding, neck formation, and partial collapse. Figure 3c is the selected example of SEM images taken from a fractured Cu foam sample baked at 400 °C for 4 h in a flow of 4% H2/Ar gas. The clear demonstration of hollow Cu shells indicates that the PS was removed totally after the calcination process (see more SEM images of samples before and after the calcination process in Figures 1, S2, and S3). As mentioned earlier, we observed that sizes of Cu grains decreased with an increase in HCHO addition, and they eventually transformed into almost continuous thin films. The escape of gaseous species generated during a thermal annealing is of interest, especially for the 1 a.u. samples. Holes observed at many Cu shells could be one route. Considering that not every shell has such holes and the PS−Cu sample at 1 a.u. has a slight surface roughness, we expect that the Cu thin film might be composed of smaller grains packed with some voids rather than a complete dense layer. Voids between smaller grains are expected to be another crucial supplementary route for the gas escape. It is worth mentioning that purposely/accidentally oxidized samples also left no detectable CuO (by SEM/EDS) after the annealing at this condition, and a complete reduction to Cu occurred. Because of a larger density increase during the reduction of oxidized E


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Figure 5. Mechanical properties of copper foams. (a) Uniform copper foam (sample A) was compared with a less uniform sample (sample B) with a similar density but containing irregular large voids and some pore collapse. (b) Load−displacement curves obtained with a spherical indenter. (c) Stress−strain curves for copper foams. (d) SEM image of sample A after the indentation test.

Figure 6. Characterization of low-density Cu foams. (a) Macroscopic copper foam machined to a 3 mm diameter cylinder for X-ray computed tomography (CT). (b) Typical X-ray projection image from the side (0° rotation) and (c) a selected slice (123 of 375) of the reconstructed 3D CT data. (d−f) SEM/EDS data for the copper foam.

synthetic conditions, including the same Cu/PS ratio and annealing conditions (Figure 5a). Figure 5b shows typical load−displacement (P−h) curves obtained with a 200 μm diameter spherical indenter. Most of the indenter displacement in the A and B copper foams is accommodated plastically, and only a small portion is elastically recovered on unloading. We first observe that the response of these two samples to the initial loading is significantly different, despite their quite similar density. Spherical indentation data can be converted to indentation stress−strain curves, as follows:

σ=

P A

(3)

ϵ=

4a 3πR

(4)

σ = Er ϵ

(5)

where σ is the average contact pressure (indentation stress), A is the contact area projected on the contact plane, ϵ is the indentation strain, a is the contact radius, R is the indenter radius, and Er is the reduced Young’s modulus. More details on F


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the fact that our copper foams have a thinner wall of ∼30 nm thickness, compared to the porous shell with ∼200 nm thickness in the previous work.21 Typically, surface areas of metal foams are given in m2 m−3 rather than m2 g−1 due to the higher density of metals. The surface area of our Cu foam corresponds to 3.4 × 106 m2 m−3, which exceeds commercially available metal sponges by a factor of 100−1000.37 The electric conductivity of the Cu foam samples with a density of ∼0.9 g/ cm3 was also measured with the four-point probe method. Bulk copper foams have a sheet resistance of ∼1.6 Ω/square. Such a low sheet resistance confirms a good electrical conductivity through the whole monolith. This work was motivated by the need of low-density copper foam targets for laser-driven high-energy-density (HED) physics experiments. In addition to a target fabrication, high surface area, highly conductive copper foams of this study can be a promising candidate in other potential applications. For example, nanostructured coppers and copper-in-charcoal (Cu/ C) have been shown to be effective catalysts for a variety of important synthetic transformations.38 Cupric oxide (CuO) is a p-type semiconductor with narrow band gap (Eg = 1.2 eV) and is known for its applications in microconductors, Li-ion battery anode materials, and chemical conversion catalysts.39−41 Bulk oxide formation or selective patterning of porous copper oxide from our conductive Cu foams may provide interesting new routes to these applications.

indentation experiments and the analysis could be found elsewhere.27,36 Figure 5c shows indentation stress−strain curves for both copper foams. The initial slope of the stress−strain curves reflects the Young’s modulus. In the first linear section with a strain up to 15−20%, elastic deformation mainly controls the total deformation of the copper foam samples. For larger strain, the slopes of the two curves gradually decrease, which could be attributed to plastic deformation of copper shells (yielding). Sample A shows a significant improvementin the indentation behavior (∼1.5× higher modulus and ∼3× larger hardness compared to those of sample B), although both samples have a quite similar shell thickness and foam densities. One reason is that a more uniform packing and a better connectivity of neighboring copper shells give rise to a higher resistance against the elastic deformation and the collapse of samples. Figure 5d shows SEM image of sample A after the indentation test, revealing no evidence of radial cracks or pileup around residual indentation impressions. Finally, we present general characteristics of a typical Cu foam sample with a density of ∼0.9 g/cm3 (∼10% relative to full density Cu) analyzed with various characterization methods. This monolithic copper foam was first machined to a 3 mm diameter cylinder in a diamond-turning process on a precision lathe (Figure 6a). Good machinability is a very important requirement to our specific application. This machined part was further used for X-ray computed tomography (CT). Figure 6b,c includes selected examples of an X-ray projection measured from the side (0° rotation) and of a slice of the reconstructed CT data near the center, respectively. The resolution of X-ray CT analysis depends on sample size (e.g., ∼3 mm) divided by the number of pixels. The effective pixel resolution is ∼6 μm in this study, which is much larger than the Cu wall thickness (∼30 nm) or pore sizes (∼1 μm) of Cu foams. X-ray CT analysis is used to monitor the uniformity of foams as well as the formation of larger defects or cracks. The pixel color (or value) represents the relative attenuation coefficient of the material at that spatial location. Specifically, black represented air, and the level of gray/white is correlated with the variation in density of the Cu foam. The CT analyses show that the material density and structure are quite uniform, especially near the middle of the sample, although some small cracks and voids are often observed (see Figures 6c and S5). SEM/EDS images in Figure 6d−f represent further details on surface morphology and composition. Unlike our previous example of gold/silver deposition on PS particles,20 the copper has a morphology of conformal coatings. No detectable oxidation was observed by SEM/EDS from fresh samples under a normal storage procedure. The diameter of Cu shells roughly corresponds to the size of the PS bead template, indicating negligible shrinkage upon the template removal. In addition to microscopic characterizations, a method to determine the Cu shell thickness is also developed. Considering that the weight loss (26 mg) during the baking comes from a complete decomposition of a PS template (i.e., Wbefore − Wafter = WPS and Wafter = WCu), weights of PS (26 mg) and Cu foam (41 mg) are used to calculate a relative volume ratio of a PS core and a Cu shell by using bulk densities of PS (1.05 g/cm3) and Cu (8.96 g/cm3). The thickness of the conformal Cu coating in the sample is determined to be about 29 nm (see eqs S1−S3). The surface area is also an important property of metallic foams. Our Cu foam (∼0.9 g/cm3) has a high BET surface area of 3.8 ± 0.2 m2/g, which is about 1.8−3.3 times higher than those of the well-known reference.21 This is due to

4. CONCLUSIONS Low-density copper foams (∼10% relative to full density Cu) with a uniform pore structure are synthesized using coppercoated polystyrene core−shell particles prepared by electroless Cu deposition. Precise control over morphologies of the core− shell particles and a thickness of copper shells in the range of tens of nanometers opens a new door to low-density copper foams that are beyond the capabilities of traditional templating approaches. Fine tuning of experimental conditions enables us to minimize the formation of irregular large voids and get metal foams with extremely uniform spherical pores. Depth-sensing indentation demonstrates that a more uniform packing and a better connectivity of neighboring copper shells play an important role in improving mechanical properties of the copper foam. Despite multiple steps required, this study provides a unique, scalable, easily accessible synthetic route to macrosized monoliths (∼millimeters and above) of metal foams with precisely controlled porosity and density and pore geometries of tens of nanometer thickness shells. These high surface area, highly conductive copper foam materials will open many new exciting opportunities to benefit both academia and industries in areas of materials science, catalysis, and energy applications.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.6b12320. Details on the experimental procedure on copper foam synthesis, determination of wall thickness, SEM images of different densities of Cu foams, TGA thermograms, and X-ray CT image (PDF) G


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(17) Rickerby, J.; Steinke, J. H. G. Current Trends in Patterning with Copper. Chem. Rev. 2002, 102, 1525−1549. (18) Ji, T.; Lirtsman, V. G.; Avny, Y.; Davidov, D. Preparation, Characterization, and Application of Au-Shell/Polystyrene Beads and Au-Shell/Magnetic Beads. Adv. Mater. 2001, 13, 1253−1256. (19) Shchukin, D. G.; Caruso, R. A. Template Synthesis of Porous Gold Microspheres. Chem. Commun. 2003, 1478−1479. (20) Nyce, G. W.; Hayes, J. R.; Hamza, A. V.; Satcher, J. H. Synthesis and Characterization of Hierarchical Porous Gold Materials. Chem. Mater. 2007, 19, 344−346. (21) Kränzlin, N.; Niederberger, M. Wet-Chemical Preparation of Copper Foam Monoliths with Tunable Densities and Complex Macroscopic Shapes. Adv. Mater. 2013, 25, 5599−5604. (22) Zhang, K.; Tan, X.; Wu, W.; Tang, Y. Template Synthesis of Low-Density Gold Foams: Density, Microstructure and Compressive Strength. Mater. Res. Bull. 2013, 48, 3499−3504. (23) Zhang, K.; Tan, X.; Zhang, J.; Wu, W.; Tang, Y. TemplateDealloying Synthesis of Ultralow Density Au Foams with Bimodal Porous Structure. RSC Adv. 2014, 4, 7196−7201. (24) Bindra, P.; White, J. R. Fundamental Aspects of Electroless Copper Plating. In Electroless Plating-Fundamentals and Application; Mallory, G. O., Hajdu, J. B., Eds.; William Andrew: New York, 1990; pp 289−329. (25) Hamada, S.; Kudo, Y.; Tojo, T. Preparation and Reduction Kinetics of Uniform Copper Particles from Copper(I) Oxides with Hydrogen. Colloids Surf. 1992, 67, 45−51. (26) Sun, K.; Wei, T. S.; Ahn, B. Y.; Seo, J. Y.; Dillon, S. J.; Lewis, J. A. 3D Printing of Interdigitated Li-Ion Microbattery Architectures. Adv. Mater. 2013, 25, 4539−4543. (27) Kucheyev, S. O.; Hamza, A. V.; Satcher, J. H., Jr.; Worsley, M. A. Depth-Sensing Indentation of Low-Density Brittle Nanoporous Solids. Acta Mater. 2009, 57, 3472−3480. (28) Feldman, B. J.; Melroy, O. R. The Mechanism of Electroless Cu Deposition: Extraction of The Oxidative and Reductive Electrochemical Half-Cell Currents from A Complete Bath. J. Electrochem. Soc. 1989, 136, 640−643. (29) Schumacher, R.; Pesek, J. J.; Melroy, O. R. Kinetic Analysis of Electroless Deposition of Copper. J. Phys. Chem. 1985, 89, 4338− 4342. (30) Sard, R. The Nucleation, Growth, and Structure of Electroless Copper Deposits. J. Electrochem. Soc. 1970, 117, 864−870. (31) El-Raghy, S. M.; Abo-Salama, A. A. The Electrochemistry of Electroless Deposition of Copper. J. Electrochem. Soc. 1979, 126, 171− 176. (32) Oita, M.; Matsuoka, M.; Iwakura, C. Deposition Rate and Morphology of Electroless Copper Film from Solutions Containing 2,2′-Dipyridyl. Electrochim. Acta 1997, 42, 1435−1440. (33) Beverskog, B.; Puigdomenech, I. Revised Pourbaix Diagrams for Copper at 25 to 300°C. J. Electrochem. Soc. 1997, 144, 3476−3483. (34) Ferrari, A. C.; Robertson, J. Interpretation of Raman Spectra of Disordered and Amorphous Carbon. Phys. Rev. B: Condens. Matter Mater. Phys. 2000, 61, 14095−14107. (35) Ferrari, A. C.; Robertson, J. Origin of the 1150-cm−1 Raman Mode in Nanocrystalline Diamond. Phys. Rev. B: Condens. Matter Mater. Phys. 2001, 63, 121405. (36) Oliver, W. C.; Pharr, G. M. An Improved Technique for Determining Hardness and Elastic Modulus Using Load and Displacement Sensing Indentation Experiments. J. Mater. Res. 1992, 7, 1564−1583. (37) Dietrich, B. Pressure Drop Correlation for Ceramic and Metal Sponges. Chem. Eng. Sci. 2012, 74, 192−199. (38) Buckley, B. R.; Butterworth, R.; Dann, S. E.; Heaney, H.; Stubbs, E. C. Copper-in-Charcoal” Revisited: Delineating The Nature of The Copper Species and Its Role in Catalysis. ACS Catal. 2015, 5, 793− 796. (39) Liu, Y.; Liao, L.; Li, J.; Pan, C. From Copper Nanocrystalline to CuO Nanoneedle Array: Synthesis, Growth Mechanism, and Properties. J. Phys. Chem. C 2007, 111, 5050−5056.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: (+1) 925-423-2072. Fax: (+1) 925-422-3570. ORCID

Sung Ho Kim: 0000-0002-0994-8098 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was performed under the auspices of the U.S. Department of Energy by Lawrence Livermore National Laboratory under Contract DE-AC52-07NA27344.

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REFERENCES

(1) Dumee, L. F.; He, L.; Lin, B.; Ailloux, F.-M.; Lemoine, J.-B.; Velleman, L.; She, F.; Duke, M. C.; Orbell, J. D.; Erskine, G.; Hodgson, P. D.; Gray, S.; Kong, L. The Fabrication and Surface Functionalization of Porous Metal Frameworks − A Review. J. Mater. Chem. A 2013, 1, 15185−15206. (2) Vukovic, I.; Punzhin, S.; Vukovic, Z.; Onck, P.; De Hosson, J. T. M.; ten Brinke, G.; Loos, K. Supramolecular Route to Well-Ordered Metal Nanofoams. ACS Nano 2011, 5, 6339−6348. (3) Tappan, B. C.; Steiner, S. A.; Luther, E. P. Nanoporous Metal Foams. Angew. Chem., Int. Ed. 2010, 49, 4544−4565. (4) Velev, O. D.; Kaler, E. W. Structured Porous Materials via Colloidal Crystal Templating: From Inorganic Oxides to Metals. Adv. Mater. 2000, 12, 531−534. (5) Kränzlin, N.; Niederberger, M. Controlled Fabrication of Porous Metals from The Nanometer to The Macroscopic Scale. Mater. Horiz. 2015, 2, 359−377. (6) Cheng, I. C.; Hodge, A. M. Morphology, Oxidation, and Mechanical Behavior of Nanoporous Cu Foams. Adv. Eng. Mater. 2012, 14, 219−226. (7) Biener, J.; Wittstock, A.; Zepeda-Ruiz, L. A.; Biener, M. M.; Zielasek, V.; Kramer, D.; Viswanath, R. N.; Weissmuller, J.; Baumer, M.; Hamza, A. V. Surface-Chemistry-Driven Actuation in Nanoporous Gold. Nat. Mater. 2009, 8, 47−51. (8) Biener, M. M.; Ye, J.; Baumann, T. F.; Wang, Y. M.; Shin, S. J.; Biener, J.; Hamza, A. V. Ultra-Strong and Low-Density Nanotubular Bulk Materials with Tunable Feature Sizes. Adv. Mater. 2014, 26, 4808−4813. (9) Jiang, P.; Cizeron, J.; Bertone, J. F.; Colvin, V. L. Preparation of Macroporous Metal Films from Colloidal Crystals. J. Am. Chem. Soc. 1999, 121, 7957−7958. (10) Vukovic, I.; Brinke, G. T.; Loos, K. Block Copolymer TemplateDirected Synthesis of Well-Ordered Metallic Nanostructures. Polymer 2013, 54, 2591−2605. (11) Hayes, J. R.; Nyce, G. W.; Kuntz, J. D.; Satcher, J. H.; Hamza, A. V. Synthesis of Bi-Modal Nanoporous Cu, CuO and Cu2O Monoliths with Tailored Porosity. Nanotechnology 2007, 18, 275602. (12) Hayes, J. R.; Hodge, A. M.; Biener, J.; Hamza, A. V.; Sieradzki, K. Monolithic Nanoporous Copper by Dealloying Mn−Cu. J. Mater. Res. 2006, 21, 2611−2616. (13) Ressnig, D.; Antonietti, M. Ultrafast Syntheses of Silver Foams from Ag2NCN: Combustion Synthesis versus Chemical Reduction. Chem. Mater. 2014, 26, 4064−4067. (14) Bi, Y.; Ren, H.; Chen, B.; Chen, G.; Mei, Y.; Zhang, L. Characterization of The Microstructures of Copper-Based Aerogels on The Sol−Gel Process. J. Sol-Gel Sci. Technol. 2012, 63, 140−145. (15) Sisk, C. N.; Hope-Weeks, L. J. Copper(II) Aerogels via 1, 2Epoxide Gelation. J. Mater. Chem. 2008, 18, 2607−2610. (16) Tang, Y.; Yeo, K. L.; Chen, Y.; Yap, L. W.; Xiong, W.; Cheng, W. Ultralow-Density Copper Nanowire Aerogel Monoliths with Tunable Mechanical and Electrical Properties. J. Mater. Chem. A 2013, 1, 6723−6726. H


Research Article

ACS Applied Materials & Interfaces (40) Jiang, X.; Herricks, T.; Xia, Y. CuO Nanowires Can Be Synthesized by Heating Copper Substrates in Air. Nano Lett. 2002, 2, 1333−1338. (41) Grugeon, S.; Laruelle, S.; Herrera-Urbina, R.; Dupont, L.; Poizot, P.; Tarascon, J.-M. Particle Size Effects on the Electrochemical Performance of Copper Oxides toward Lithium. J. Electrochem. Soc. 2001, 148, A285−A292.

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